Effects of Ionizing Radiation on the Layered Semiconductor Tungsten

Home , Silicon dioxide, Tungsten diselenide

The Pennsylvania State University

The Graduate School

College of Earth and Mineral Sciences

EFFECTS OF IONIZING RADIATION ON THE LAYERED

SEMICONDUCTOR TUNGSTEN DISELENIDE

A Thesis in

Materials Science and Engineering

Roger Craig Walker II

Submitted in Partial Fulfillment of the Requirements for the Degree of

Master of Science

August 2017

The thesis of Roger C. Walker II was reviewed and approved* by the following:

Joshua A. Robinson Associate Professor of Materials Science and Engineering Thesis Adviser

Saptarshi Das Assistant Professor of Engineering Science and Mechanics

Suzanne Mohney Professor of Materials Science and Engineering Materials Science and Engineering Program Chair

*Signatures are on file in the Graduate School.

Abstract

The miniaturization of electronic devices has been a critical for the successes of space exploration. Scaling down individual transistors allows integrated circuits to be more powerful and flexible, while reducing their weight and volume. This is highly attractive for space electronics given the cost of launching objects into space. Semiconducting two-dimensional materials (2DMs), such as tungsten diselenide (WSe2), are of recent interest for developing ultimately-scaled transistors. These materials have a layered structure with a thickness of individual atoms at the monolayer limit, and are compatible with existing device fabrication processes. However, the space environment imposes an additional constraint on the device performance of radiation hardness. Space is filled with high energy ionizing radiation – i.e. X- rays, gamma rays, protons, electrons, and heavy ions – that can damage and destroy devices. As such, the stability of WSe2 against various forms of ionizing radiation should be critically examined to determine if it possesses sufficient radiation hardness for space applications. In this work, the stability of WSe2 against X-rays, low energy plasma, protons, and heavy metal ions has been studied. WSe2 was prepared using a combination of transfer via mechanical exfoliation and growth via metal-organic chemical vapor deposition (MOCVD). Mechanical exfoliation is used to transfer micrometer-thick flakes onto substrates. To obtain nanometer thick films and individual layers, MOCVD was used. In this process, WSe2 is formed by controllably reacting gas phase pre-cursors of its constituent elements at elevated temperatures. The pre- cursors react on the surface of a substrate (e.g. silicon carbide (SiC)), nucleate as a WSe2 crystal, and grow into individual domains or coalesced films. X-ray exposure was achieved using the soft X-rays generated in an X-ray photoelectron spectroscopy (XPS) system for 12 hours in ultra- high vacuum. Plasma was generated by the ionization of helium and oxygen gas under medium vacuum conditions using an RF generator. Protons and heavy ions were generated by stripping the electrons from gases of hydrogen, iron, and silver, and were accelerated into the WSe2 samples under high vacuum conditions. Total fluences are varied from 1014 particles/cm2 to ~1018 particles/cm2, which covers the typical expected values for space electronics. Modification to the WSe2 induced by ionizing radiation is primarily studied using XPS, a highly surface sensitive characterization technique. It is found that each form of ionizing radiation has a different impact on the WSe2. Soft X- rays induce a charge in the WSe2 that is dependent on the domain size and surface coverage. Isolated domains of MOCVD-grown WSe2 are subject to a small surface potential on the order of 100 meV, which is attributed to interface states formed at the edges of these domains. Fully coalesced films of MOCVD-grown WSe2 were not significantly affected by X-rays, even at a fluence of 1.86 × 1018 photons/cm2. The low energy plasma can be used to convert the top layers of WSe2 into a mix of tungsten oxide (WOx) and selenium oxide. MOCVD-grown WSe2 is found to be more susceptible to oxidation than the exfoliated material, perhaps due to a greater defect density. Surprisingly, the extent of oxidation is not very sensitive to changes in plasma parameters such as pressure or gas flow rate, suggesting that much of the oxidation is due to air exposure after the treatment. Knowledge of the effects of protons and heavy ions with energies in the mega-electron-volt (MeV) range are the most directly related to knowledge of WSe2 stability in the space environment. It was determined that WSe2 flakes exfoliated onto SiC were not affected by 2 MeV proton bombardment up to a dose of 1015 protons/cm2. Once a dose of 1016 /cm2 had been achieved, there was significant charge transfer between the WSe2 and the SiC substrate that

iii arose from damage to the SiC. The SiC was observed to turn black due to the large generation of vacancies corresponding to a deep acceptor lying 1.1 eV below the conduction band edge. The damage to SiC and charging in WSe2 are expected to permanently disrupt device operation, despite a lack of physical damage to the WSe2. No physical changes such as oxidation occurred in the WSe2 due to this bombardment, suggesting that it is highly stable against cosmic radiation that mostly consists of protons. As such, the choice of substrate will be critical in tuning the radiation hardness of WSe2-based devices. Additionally, this damage threshold is several orders of magnitude greater than what was observed for damage to MoS2-based devices on SiO2. As such, top-gated 2DM-based devices using semi-insulating SiC as a substrate are suggested for future miniaturized space electronics. 1 MeV protons are found to have a qualitatively similar effect as the 2 MeV protons. By using lower energy protons (i.e. 0.2 MeV and 0.04 MeV), the vacancy-rich region in the SiC is found to move closer to the sample surface, and the charging in the WSe2 is reduced. In contrast, exposure to heavy metal ions such as iron and silver at a fluence of 1016 ions/cm2 leads to significant physical damage for both the exfoliated WSe2 and the SiC substrate. The WSe2 is converted into a mixture of Se-poor WSe2 and WOx due to selenium volatilization during the bombardment and oxidation once exposed to air. The material continues to lose selenium and gain oxygen over time even when stored in medium vacuum. The SiC turns black and becomes fully amorphous. Significant concentrations of silicon oxide are generated after air exposure, suggesting that carbon is preferentially sputtered out from the sample due to ion bombardment. However, the SiC is stable under storage in medium vacuum and does not significantly change from its damaged state. The valence band spectra revealed that the heterostructure undergoes a semiconducting to metallic transition due to the extensive damage. As such, a device in the space environment based on the WSe2/SiC heterostructure would be expected to be permanently destroyed after exposure to this fluence of heavy metal ions. Complimentary proton and ion studies using MOCVD-grown material are in the planning stages and highly recommended as future work. This work has elucidated some of the physical responses of WSe2 to ionizing radiation. This material appears to be robust against soft X-rays but is susceptible to low energy plasma, high energy protons, and high energy ions. Further work determining the sensitivity of WSe2-based devices to X-rays, protons and ions is recommended to measure the changes to the electronic properties of this material, as well as critical points of failure in the devices. To accomplish this, the stability of the WSe2/dielectric interface must also be studied, and the compatibility of high-k dielectric materials, such as HfO2, with 2DMs must be determined.

Table of Contents List of Figures ...... vii List of Tables ...... x Acknowledgements ...... xi Chapter 1 – Introduction and Literature Review ...... 1 1.1 A Brief Review of Two-Dimensional Materials ...... 1

1.2 Electronic Properties of Tungsten Diselenide (WSe2) ...... 6 1.3 The Near-Earth Space Radiation Environment ...... 8 1.4 Radiation Damage to Semiconducting Materials and Devices ...... 12

1.5 Radiation Damage to WSe2 ...... 15 Chapter 2 – Experimental Procedure ...... 18

2.1 Mechanical exfoliation of WSe2 ...... 18

2.2 Metal-organic chemical vapor deposition (MOCVD) of WSe2 ...... 19

2.3 Stability of MOCVD-grown WSe2 in the XPS environment ...... 21

2.4 Oxygen plasma treatment of WSe2 ...... 22

2.5 Proton and heavy ion irradiation of exfoliated WSe2 ...... 24

Chapter 3 – Stability of MOCVD WSe2 in the XPS environment ...... 27

3.1 X-ray attenuation and energy absorption in WSe2 ...... 27 3.2 Low energy electron flood gun utility in XPS ...... 29

3.3 Charging in WSe2 from soft X-rays and low energy electrons ...... 30

3.3 Comparison between WSe2 and MoS2 ...... 32

Chapter 4 – Stability of exfoliated and MOCVD WSe2 in air and rapid oxidation with low energy plasma treatments ...... 36

4.1 Stability of exfoliated and MOCVD WSe2 in air and medium vacuum ...... 36

4.2 Oxidation of exfoliated and MOCVD WSe2 when exposed to low energy plasma ...... 37 4.3 Impact of plasma parameters and exposure time on the oxidation process ...... 43

Chapter 5 – Impact of Proton Irradiation on Exfoliated WSe2 ...... 49 5.1 Effect of 2 MeV protons at various fluence levels ...... 49 5.2 Effect of 1 MeV, 200 keV, and 40 keV protons at a fluence of 1016 protons/cm2 ...... 52 5.3 Proton impact on the SiC substrate and its optical properties ...... 53

5.4 Proton impact on WSe2 and SiC valence band offset ...... 56

5.5 Comparisons between WSe2 and MoS2 ...... 60

Chapter 6 – Impact of Heavy Ion Irradiation on Exfoliated WSe2 ...... 62

6.1 Heavy metal ion damage to exfoliated WSe2 ...... 62 6.2 Heavy metal ion damage to SiC substrates ...... 64

6.3 Heavy metal ion modification of WSe2 and SiC valence band spectra ...... 67 Chapter 7 – Conclusions and Future Work ...... 69 References ...... 71

List of Figures

Figure 1.1 – The periodic table of 2DMs. Color legend: black = mono-elemental [6]; blue = tri- chalcogenides [7]; yellow = di-chalcogenides [8]; red = mono-chalcogenides [9,10]; purple = 2D nitrides [11,12]; orange = 2D oxides [13–15]. For simplicity, only the more cationic elements are shaded for compounds. 2D silicon carbide is not shown here, but has also been recently synthesized [16]...... 1 Figure 1.2 – The responsivity (R) of various layered materials to visible and UV light. R > 1 has been deemed necessary for practical applications involving 2DMs [33]...... 4 Figure 1.3 – Larger view of WSe2 down the c-axis (top) and in-plane (bottom) ...... 6 Figure 1.4 – The motion of particles trapped in the van Allen radiation belt [77] ...... 9 Figure 1.5 – Energy distribution of trapped protons [76] ...... 10 Figure 1.6 – Flux distribution of trapped protons [76] ...... 11 Figure 1.7 – Individual impact craters on the surface of WSe2 by bombardment via 5 keV argon ions and imaged via STM at an image size of 45 nm by 45 nm [53] ...... 16 Figure 1.8 – Atomically-resolved STM image (8 nm by 8 nm) of a corrugated WSe2 surface [53] ...... 16 Figure 1.9 – The change in resistance of a few-layer WSe2-based device on a SiO2/Si substrate due to helium ion bombardment [74]...... 17 Figure 2.1 – Optical image of exfoliated WSe2 flakes ...... 18 Figure 2.2 – XPS W 4f spectra of exfoliated WSe2 without and with the heat treatment. Only one doublet is observed, corresponding to WSe2. Intensity for all XPS measurements is in units of counts per second (CPS)...... 19 Figure 2.3 – Schematic MOCVD reactor...... 20 Figure 2.4 – Kratos Axis Ultra spectrometer ...... 22 Figure 2.5 – M4L plasma asher tool. Image copyright PVA TePla America...... 23 Figure 2.6 – WSe2/SiC samples mounted onto SiO2 pieces and stored in a vacuum chamber .... 24 Figure 2.7 – Thermal monitoring of proton and ion exposed samples ...... 25 Figure 2.8 – PHI VersaProbe II spectrometer ...... 26 Figure 3.1 – X-ray attenuation by the layered semiconductors MoS2 and WSe2...... 27 Figure 3.2 – X-ray attenuation and absorption by WSe2 and the three substrates...... 29 Figure 3.3 – Shifts in binding energy of the VBM and W 4f electrons in WSe2 ...... 31 Figure 3.4 – Comparison between metal (W 4f) and chalcogen (Se 3d) peak shifts ...... 31 Figure 3.5 – Scanning electron microscopy images of the following samples: a) WSe2 on graphite paper; b) WSe2 on Al2O3 c) WSe2 on SiC; d) MoS2 on Al2O3; e) MoS2 on SiC ...... 32 Figure 3.6 – Binding energy shifts of the metal peaks of all studied 2DMs ...... 33 Figure 3.7 – Binding energy shift of the VBM for MoS2 on sapphire ...... 34 Figure 4.1 – Changes in the W 4f spectra of exfoliated WSe2 due to storage in medium vacuum or air exposure for three days. The small contribution of 5p3/2 core shell electrons is not shown.37 3/2 Figure 4.2 – Changes in the W 4f and 5p spectra of MOCVD-grown WSe2 due to storage in medium vacuum or air exposure for three days...... 37 Figure 4.3 – W 4f spectra of exfoliated and as-grown WSe2 before and after the standard plasma treatment. Units on the vertical axis are counts per second (CPS) normalized by the intensity of the major W 4f peak for WSe2...... 39

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Figure 4.4 – Se 3d spectra of exfoliated and as-grown WSe2 before and after the standard plasma treatment. Units on the vertical axis are CPS normalized by the intensity of the major Se 3d peak for WSe2...... 39 Figure 4.5 – W 4f spectra of the control samples of exfoliated and as-grown WSe2. Units on the vertical axis are CPS normalized by the intensity of the major W 4f peak for WSe2. Note that the as-grown material starts out with some oxide...... 41 Figure 4.6 – Se 3d spectra of the control samples of exfoliated and as-grown WSe2. Units on the vertical axis are CPS normalized by the intensity of the major Se 3d peak for WSe2...... 41 Figure 4.7 – AFM scans of the following samples: (a) MOCVD WSe2 grown on epitaxial graphene (EG) prior to treatment; (b) MOCVD WSe2 on EG after treatment; (c) MOCVD WSe2 grown on 6H-SiC prior to treatment; (d) MOCVD WSe2 on SiC after treatment. The treatment used here is the standard treatment discussed in Section 2.4...... 42 Figure 4.8 – Changes in the concentration of tungsten oxide and selenium oxide due to low energy plasma exposure when the treatment time is varied...... 44 Figure 4.9 – Changes in XPS spectra of MOCVD-grown WSe2 on SiC due to plasma treatment and subsequent storage in air. There is no significant change in the extent of oxidation with air exposure (72.55% WOx after treatment vs. 73.55% WOx after storage in air)...... 45 Figure 4.10 – Changes in the concentration of tungsten oxide and selenium oxide due to low energy plasma exposure when the flow rate of oxygen gas is varied...... 46 Figure 4.11 – Changes in the concentration of tungsten oxide and selenium oxide due to low energy plasma exposure when the chamber pressure is varied. Control samples are at zero...... 47 Figure 4.12 – Changes in the concentration of tungsten oxide and selenium oxide due to low energy plasma exposure when the RF power is varied. Control samples are at zero...... 48 Figure 5.1 – Binding energy shifts in W 4f and Se 3d core shell levels due to 2 MeV protons at various fluence levels ...... 50 Figure 5.2 – W 4f and 5p3/2 core shell electron spectra from before, after, and two weeks after exposure to 1016 protons/cm2 with an energy of 2 MeV ...... 51 Figure 5.3 – Binding energy shifts in W 4f and Se 3d core shell levels from proton exposure at various energies at a fluence of 1016 protons/cm2 ...... 53 Figure 5.4 – Binding energy shifts in Si 2p and C 1s core shell levels due to 2 MeV protons at various fluence levels ...... 54 Figure 5.5 – Binding energy shifts for Si 2p and C 1s core shell levels from proton exposure at various energies at a fluence of 1016 protons/cm2 ...... 54 Figure 5.6 – (a) Camera picture of samples prior to irradiation; (b) camera picture of samples after irradiation. Samples were exposed to 2 MeV protons unless stated otherwise...... 55 Figure 5.7 – Absorbance spectra of the SiC samples with and without irradiation ...... 55 Figure 5.8 – Plot of initial valence band offsets between WSe2 and SiC and changes induced to the VBO by proton exposure at 1016 protons/cm2 ...... 58 Figure 5.9 – Binding energy shifts for Mo 3d and S 2p core shell levels from proton exposure at various energies at a fluence of 1016 protons/cm2 ...... 60 Figure 5.10 – Binding energy shifts for Si 2p and C 1s core shell levels from proton exposure at 16 2 various energies at a fluence of 10 protons/cm from the MoS2/SiC samples ...... 61 Figure 6.1 – W 4f spectra of as-exfoliated samples (black) and after irradiation (red) ...... 63 Figure 6.2 – Se 3d spectra of as-exfoliated samples (black) and after irradiation (red) ...... 63 Figure 6.3 – Changes in C:Si ratio, Se:W ratio, SiOx:SiC ratio, and WOx:WSe2 ratio due to heavy metal ion exposure and storage in medium vacuum for two weeks...... 64

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Figure 6.4 – Optical image (a) and height profile (b) of exfoliated WSe2 flakes on SiC ...... 65 Figure 6.5 – Si 2p spectra before (black) and after (red) heavy metal ion irradiation ...... 65 Figure 6.6 – C 1s spectra before (black) and after (red) irradiation ...... 66 Figure 6.7 – Valence band spectra of the WSe2/SiC heterostructure before (black) and after (red) heavy metal ion irradiation ...... 67

List of Tables

Table 1.1 – WSe2 device chart ...... 8 Table 3.1 – Table of mass attenuation coefficients, densities, attenuation coefficients and band gap values for WSe2 and the three substrates ...... 29 Table 5.1 – Tabulated binding energy shift values used for Figure 5.1 ...... 50 Table 5.2 – Tabulated binding energy shift values used for Figure 5.3 ...... 53 Table 5.3 – Tabulated peak shifts for bulk single crystal WSe2 ...... 56 Table 5.4 – Tabulated initial valence band offsets between WSe2 and SiC and changes induced to the VBO by proton exposure at 1016 protons/cm2 ...... 57 Table 5.5 – Core level separation data used to calculate the VBO between WSe2 and SiC ...... 58 Table 5.6 – Tabulated binding energy shifts for exfoliated MoS2 ...... 60

Acknowledgements

I would like to personally thank my adviser (Dr. Joshua A. Robinson) and collaborators (Tan Shi and Dr. Igor Jovanovic) for their support. I would also like to thank my group members and friends for their encouragement over the years.

This research was funded by the Defense Threat Reduction Agency under grant HDTRA1-14-1- 0037 and was completed with the assistance of the Penn State Materials Characterization Laboratory, the Penn State Nanofabrication Facility, and the Michigan Ion Beam Laboratory.

Chapter 1 – Introduction and Literature Review

1.1 A Brief Review of Two-Dimensional Materials Two-dimensional materials (2DMs) are materials with covalent or ionic bonding within a plane of atoms and van der Waals bonding outside of that plane, giving them a layered structure. They can be formed from various elements across the periodic table as shown in Figure 1.1. 2DMs are commonly derived from bulk solid counterparts; for example, layers of the 2DM graphene can be derived by exfoliating bulk pieces of graphite [1]. Research on the bulk forms of layered solids goes back decades, and single layers of 2DMs have been confirmed and studied as early as the 1980’s [2]. However, interest in single layers was relatively minimal until 2004, after the isolation and re-discovery of graphene after mechanical exfoliation from bulk graphite [1]. After that point, much more research has been carried out to stabilize graphene and other 2DMs in order to understand their basic electronic properties and determine their possible applications [3–

5].

shaded for compounds. 2D silicon carbide is not shown here, but has also been recently synthesized [16].

2DMs can be grouped into various classes, as shown in Figure 1.1, and include monoelemental materials (e.g. graphene [17]), 2D nitrides (e.g. hexagonal boron nitride [12]), transition metal dichalcogenides (e.g. molybdenum disulfide (MoS2) and tungsten diselenide

(WSe2) [8]), 2D oxides (e.g. molybdenum trioxide [18]), and Group III and IV monochalcogenides (e.g. gallium selenide [9]). Although these materials are all considered 2D at the single layer limit, their properties are highly dependent on their elemental composition and individual crystal structures. For example, graphene layers are composed entirely of sp2 bonded carbon atoms with hexagonal symmetry, and have metallic conductivity due to a small overlap of the conduction and valence bands [1]. More precisely, the two bands meet at a single point, referred to as the Dirac point. Hexagonal boron nitride has the same structure as graphene, but its differing elemental composition and more ionic in-plane bonding makes it an insulator with a band gap of ~6 eV [19].

These differences in properties are used to support different applications for these materials, each matching the 2DM. Graphene, being metallic, is of interest for electrical contacts to both

2D [20] and 3D [21] semiconductors. Semiconducting 2DMs such as MoS2 and WSe2 are of interest as the channel materials in nano-scale electronics due to their layer-dependent band gaps.

Layered semiconductors tend to increase in band gap energy as the number of layers is reduced due to quantum confinement effects and a reduction of interlayer interactions. The general trend can be described using the following relation that was recently deduced for phosphorene: Eg = Eo

β + C / n , where Eg is the band gap, Eo is the bulk band gap, n is the number of layers, and C and β are fitting parameters [22]. As such, the biggest changes occur in ultra-thin samples, opening up the possibility of using different numbers of layers to generate heterostructures [23]. It also

2 means the precise control over the number of layers of semiconducting materials will be needed for practical devices.

Narrow band gap materials (i.e. less than 3 eV) are of interest for channel materials in transistors [24], IR communications [25], and solar cells [26], while wider band gaps (i.e. greater than 3 eV) would be of interest as dielectric materials [27], visible-blind / solar-blind photodetectors [28], and in power electronics [29]. The band gap determines the selectivity of photodetectors based on that material. Zero band gap 2DMs such as graphene have a broadband sensitivity to all light wavelengths. On the other hand, 2DMs such as single layer GaS and GaSe are proposed to be optimal for selective UV light detection due to a wide band gap of ~ 3 eV as a single layer [9,30]. The quality of a photodetector to specific wavelengths is characterized by its responsivity. The responsivity of various 2DMs to visible and ultraviolet light is shown in Figure

1.2. Graphene by itself generally has low responsivity, but adding graphene to another material can be used to make a high responsivity heterostructure. On the other hand, certain individual

2DMs such as MoS2 and WSe2 are also able to achieve large responsivity values. A wide variety of light sensors have been developed based on 2DMs, and the knowledge gained from these investigations is discussed in greater detail in review articles such as [31] and [32].

Figure 1.2 – The responsivity (R) of various layered materials to visible and UV light. R > 1 has been deemed necessary for practical applications involving 2DMs [33].

Additionally, tuning the properties of a single 2DM has been investigated using various doping schemes. These materials, despite their weak out of plane bonding, are highly sensitive to adsorbed species, which makes them ideal for gas sensors based on charge transfer [34] and allows for simple doping schemes by gas exposure [35]. These materials are often p-type doped due to adsorbed oxygen, which draws electrons out of 2DMs. Storage and annealing in vacuum would be needed to observe the intrinsic properties unless a passivation layer is applied. Charge transfer can also arise from substrate interactions with 2DMs, making the choice of substrate critical in pursuing electronic applications [36]. This charge transfer and changes in the doping

4 level are expected at the interfaces between 2DMs or between 2DMs and bulk solids in order to achieve equilibrium. Depending in the materials in question, this interface might be affected by

Fermi level pinning (FLP), a phenomena where the location of the Fermi level and the resulting band alignment are governed by charge transfer into interfacial states between the two materials rather than by the inherent properties of the two materials [27,37]. FLP is common in bulk solids, where it interferes with the formation of high quality electrical contacts and large dielectric band offsets. The insertion of graphene at the interface of a bulk metal and semiconductor has been shown to reduce FLP [38]. 2DMs such as graphene interfere with the formation of the interface states that lead to FLP while only introducing a small barrier to current flow, so a device would overall be improved. FLP has also been observed at the interfaces between 2DMs and bulk solids

[39,40]. One mechanism proposed [41,42] and successfully used [43] to reduce the impact of

FLP on 2DM-based devices is using 2D metals to contact 2D semiconductors. This logic may also be applied to the use of the 2D dielectric hBN as a gate dielectric material [20,44].

Traditionally, 2DM heterostructures are manufactured via the mechanical exfoliation of individual layers. This technique has been used to have control over the stacking order and layer number down to individual layers, which is excellent for fundamental research [45]. However, the choice for practical devices would be to grow each layer in the heterostructure using techniques such as powder vaporization [46], chemical vapor deposition (CVD) [47], metal- organic CVD [48], and molecular beam epitaxy [49]. These techniques have individually been successful in growing certain 2DMs in a controlled manner, and some have been used to generate heterostructures. However, precise control over stacking sequences and layers numbers using these growth techniques is still under development. As such, the work discussed later in

5 this thesis concerns only one 2DM – tungsten diselenide – and uses a combination of exfoliated and as-grown material in order to understand its behavior at various thicknesses.

1.2 Electronic Properties of Tungsten Diselenide (WSe2) Basic research on bulk crystals of tungsten diselenide (WSe2) was initiated decades ago due to its semiconducting properties, opening up possible applications in diodes, solar cells and thermoelectric generators [50–55]. The crystal structure and hexagonal symmetry of WSe2 is shown schematically in Figure 1.3. The thickness of a single layer is often experimentally estimated as 0.8 nm, and includes the van der Waals gap between the layer and the substrate

[56]. The thickness of the single layer alone is about 0.35 nm, and the in-plane lengths of the unit cell are ~0.33 nm [57,58]. Individual layers are composed of tungsten atoms that traditionally have trigonal prismatic coordination with selenium. This is referred to as the 2H phase, which is the stable and most commonly encountered phase. Octahedral coordination is rare, and has only been observed in the so-called 1T phase. The 1T phase has metallic conductivity, but must be induced by some chemical or physical means [59]. As the 2H phase will be the focus of the work done in this thesis, all further discussion will be limited to this phase.

Figure 1.3 – Larger view of WSe2 down the c-axis (top) and in-plane (bottom)

In its bulk form, WSe2 has a band gap of around 1.2 eV. As it is thinned to a single layer, the electrical band gap is found to monotonically increase to about 2.2 eV [23,60,61], while the optical band gap saturates at 1.65 eV [23,48]. This difference is due to the large exciton binding energy, which is the extra energy required to separate the electron and hole that are bound together in an exciton. Accurate knowledge of the electrical band gap is needed to evaluate the performance of devices based on single layers of WSe2 and compare it to other 2DMs and bulk solids, while the optical band gap determines its selectivity as a light sensor. This brief literature review will focus on electronic performance, and not optical. This material is capable of ambipolar conduction, making it suitable for either n-type conductivity or p-type conductivity

[62]. That said, WSe2 is commonly referred to as a p-type semiconductor in the literature, with its p-type characteristic arising from intrinsic defects [63].

A chart showcasing the basic characteristics (i.e. carrier mobility, ON/OFF ratio and subthreshold swing) of various WSe2-based transistors is shown in Table 1.1. Devices with large

ON/OFF ratio of > 108, hole mobility of ~300 cm2/V∙s, and ideal subthreshold swing of 60 mV/decade can be made based on single or few layers of this material, which is promising for developing this material for practical applications. Further improvements to the hole mobility are possible using p-type surface dopants such as octadecyltrichlorosilane [64]. More complex device structures such as inverters based on single-layer [65] and bi-layer [66] WSe2 have also been demonstrated due to the ambipolar conduction abilities of this material. However, most of these transistors and inverters were based on the mechanical exfoliation of individual layers. It would be preferable from a manufacturing point of view to controllably grow the materials and then contact them, using a technique such as metal-organic chemical vapor deposition [48] or

7 molecular beam epitaxy [23,61]. Device fabrication and evaluation based on as-grown material is still under development, as well as optimization of growth processes.

Table 1.1 – WSe2 device chart Ref. Number Fabrication Dielectric Contact Mobility ON/OFF Subthreshold of Method Metal (cm2/V∙s) ratio Slope Layers (mV/decade) 6 [67] 1 Mechanical Al2O3 In 142 > 10 exfoliation 8 [67] 1 Mechanical Al2O3 Ag ~200 > 10 300 exfoliation 6 [68] 1 Mechanical ZrO2 Pd (w/ ~250 > 10 ~60 exfoliation NO2) [69] 1 Powder Ion gel Au ~90 (p) > 105 (p) vaporization ~7 (n) > 104 (n) 4 6 [70] Pyramid Powder SiO2 Pd ~40 10 – 10 vaporization [56] 1 Powder SiO2 Ti 0.2 vaporization 6 [71] 12 Mechanical SiO2 Au ~302 ~10 ~250 exfoliation [72] ~3 Selenization SiO2 Cr ~0.2 [73] 5 Mechanical hBN/HfO2 ~300 exfoliation 6 [74] ~11 Mechanical SiO2 Au 64 > 10 exfoliation

1.3 The Near-Earth Space Radiation Environment The knowledge of radiation tolerance is needed mainly for specific applications such as space electronics, UV detectors, and neutron counters. Typical hazards in the space environment are high energy particles that may either be photons (i.e. hard UV light, X-rays, and gamma rays) or ions (i.e. protons, electrons, and ions). These forms of radiation are all ionizing, meaning that they have sufficient energy to break bonds and liberate electrons. Damage to semiconductors specifically will be further discussed in Section 1.4; the rest of this section will focus on a description of the near-earth space radiation environment.

Particle exposure in space comes from a limited number of sources: 1) energetic electrons, protons, and heavy ions located in the van Allen radiation belts; 2) energetic solar events such as solar flares; and 3) cosmic rays (more formally known as galactic cosmic radiation) that originate from outside the solar system but within our galaxy [75,76]. Solar particles and cosmic rays include X-rays and gamma rays as well as high energy charged particles. Particles from solar events and galactic cosmic radiation are transient, and increase both the dose and dose rate experienced by space electronics. On the other hand, particles in the van Allen radiation belts are trapped by the Earth’s magnetic field. The motion of trapped charged particles around the Earth is shown in Figure 1.4. Due to their opposite charge, electrons drift eastward while protons and ions drift westward.

Figure 1.4 – The motion of particles trapped in the van Allen radiation belt [77]

The number of trapped heavy ions is very small when compared to the numbers of trapped electrons and protons, and they tend to have low energy. Generally, protons closer to the Earth will have a higher average kinetic energy, due to concentrations of highly intense radiation that lie close to Earth. This is shown in Figure 1.5 and 1.6, where the energy and flux of mega- electron-volt (MeV) energy protons is plotted as a function of the McIlwain parameter (L), a dimensionless ratio of distance divided by the Earth’s radius as measured at the geomagnetic equator [78]. Lower energy protons (i.e. 0.1 MeV) are common throughout the near-earth space

9 environment, while higher energy protons (i.e. 400 MeV) are concentrated in the “inner zone” that extends to L = 2. Inner zone electrons tend to have energy of 5 MeV or less. The “outer zone” extends from L = 3 to L = 12. Proton energy and flux are generally monotonically decreasing in this zone, while the electron energy can extend up to 7 MeV and the flux can be an order of magnitude greater than in the inner zone. Between these two zones is the “slot”

(sometimes informally called the “safe zone”), which generally has very low particle density except in the event of solar activity [76].

Figure 1.5 – Energy distribution of trapped protons [76]

Figure 1.6 – Flux distribution of trapped protons [76]

These solar events increase the electron density in the slot by several orders of magnitude due to the resultant geomagnetic storms on Earth. Additionally, these events add large quantities of energetic particles to the near-Earth space environment. The majority of these particles are protons, and most of the remaining are alpha particles. The enhanced particle fluxes (i.e. up to

1000 cm2/s) of very high energy (i.e. greater than 10 MeV) protons are maintained for several days, but decay slowly over periods of several months. For most solar events, the emission of heavy ions is less than from galactic cosmic radiation and can be neglected. This is not the case in major events, where the abundance can increase by several orders of magnitude, leading to significant deleterious effects to space electronics (see Section 1.4 for further discussion).

Similar to solar event particles, cosmic rays consist primarily of very high energy protons; it has been estimated that galactic cosmic radiation is 85% protons, 14% alpha particles, and 1% heavier ions [75,76]. These protons generally have energy between 10 MeV and 10 GeV, with a peak in their distribution near Earth of about 1 GeV. It is possible, but extremely rare, for cosmic rays to have higher energy than this, and the range of physically possible energies is known to

11 extend up to 0.3 TeV (or 3 x 1020 eV). Interactions with lower energy cosmic rays tend to be suppressed due to interference from the solar wind and interplanetary magnetic field. As such, their average energy has been estimated as 300 MeV. Cosmic rays are considered to have a uniform distribution, and are a low flux (i.e. ~4 particles/cm2-s [76]) and omnidirectional source of radiation unless very close to the Earth. At those low altitude orbits, the Earth will create regions free from galactic cosmic radiation. For the case of both van Allen belt radiation and cosmic rays, some of these particles may actually be anti-matter, mostly anti-protons [79,80].

1.4 Radiation Damage to Semiconducting Materials and Devices Ionizing radiation can interact with semiconductors through inelastic scattering interactions with electrons (electronic stopping) and elastic scattering interactions with nuclei (nuclear stopping).

After an electronic stopping event, electrons are ejected from the electron cloud, leading to ionization and charge buildup. In the case of a nuclear stopping event, the atom is ejected from its original position, leading to displacement and accumulation of point defects. Both are expected with high energy radiation exposure, and the cumulative effects are commonly separated to the total ionizing dose (TID) and displacement damage dose. TID is measured in units such as rad and Gray (100 rad = 1 Gy = 1 J/kg), Gray being the unit of absorbed radiation from the International System of Units, and is the primary focus of studies on radiation damage in semiconductors. Because different materials have different numbers of atoms within the same mass, the material should be specified (i.e. 100 rad(Si) or 1 Gy(Al)). The energy going into ionization is given by the stopping power, or the linear energy transfer (LET). LET is typically given in units such as keV/µm, which corresponds to the amount of energy transferred to the material being irradiated per unit thickness. The absorbed dose is then the integral over the energy range of the product of the particle spectrum and the stopping power. Alternatively, one

12 may track damage by only considering the particle spectrum, which tends to be common in the literature where the particle type, particle energy and length of exposure can be controlled with high precision.

The space dose rate is generally estimated to be between 0.0001 and 0.01 rad/s, but the considerable length of space missions can lead to large doses (i.e. up to 3 ∙ 105 rad each year)

[75]. To achieve such large total doses in a short time frame, the dose rate in component testing is generally much higher (i.e. 50 rad/s). However, it has been determined that some devices are actually more sensitive to radiation damage at the low rates typical of space than at the high rates used in testing [81]. This phenomenon is referred to as enhanced low-dose-rate sensitivity

(ELDRS), and is tested by exposing components at a rate of 0.01 rad/s. Components either have

ELDRS or they do not, and having ELDRS is not enough to make a component unsuitable for use in space; it depends on the application and tolerance for error. However, if ELDRS is present, it must be accounted for when determining the expected changes in device operation due to TID and displacement damage.

Electrons cause mainly ionization damage due to their light mass, while protons, alpha particles, and heavy ions can cause both ionization and displacement. Photons can only cause ionization. Very sensitive un-hardened silicon devices start to fail at a TID of 1 krad (10 Gy), while hardened microcircuits will start to fail at 10 Mrad (100 kGy) [82]. When ionization occurs, both electrons and holes will be produced; however, the lower mobility species will be trapped, leading to shifts in the threshold voltage, increases in leakage current, and changes in the timing of switches. Meanwhile, displacement damage leads to alteration of the band gap and band alignment of semiconductors, increases in carrier recombination and leakage current, and decreases in the carrier mobility [75]. This cumulative damage is inevitable, but can be mitigated

13 with shielding, derating, and conservative design. Care must be taken to ensure that the shielding does not significantly increase TID due to conversion of incoming high energy particles into low energy particles and bremsstrahlung [75,76].

For very high energy particles, large amounts of energy may be deposited by a single particle, leading to single event effects (SEE). An interaction can be considered an SEE, if a sufficiently large amount of charge can be induced by a single ionization event and collected all at once. This charge can be referred to as Qcrit, and increases in Qcrit correspond to a reduced rate of SEE [75].

In general, protons can start inducing SEEs if they have energy of 10 MeV or more, and the likelihood increases significantly if the proton energy meets or exceeds 30 MeV [83,84]. The mechanism for this is large amounts of ionization and displacement damage induced at once, along with nuclear reactions at higher energies. In general, there are two types of SEE: “soft”, or non-destructive events such as upsets, transients, and functional interrupts; and “hard”, or destructive events such as latchup, burnout, and gate rupture. SEE is not affected by shielding, but can be mitigated using other means such as redundant circuits. Alternatively, parts immune to SEE many used if available [75].

Silicon has historically been the focus of studies on the radiation tolerance of semiconductors due to its long use in industry and in space electronics [85]. However, the stability of wide band gap semiconductors such as silicon carbide [86] and gallium nitride [87] are also being studied.

As WBG materials, they are expected to have a higher tolerance to ionizing radiation when compared to silicon. More details on these specific materials are given in the cited review articles. Recently, the damage done to layered materials by ionizing radiation has also begun to be studied, and has also been summarized [32]. It is expected that based on the nano-scale dimensions of these materials, the probability of a highly damaging event is reduced, and so they

14 would be expected to be more tolerant against radiation. Studies of the layered semiconductor

WSe2 will be discussed in the next section.

1.5 Radiation Damage to WSe2 The knowledge of how radiation impacts WSe2 and WSe2-based devices is needed to qualify its use in radiation-rich environments such as space and determine its ability to act as a radiation detector. However, the literature in this area remains limited, given that most radiation studies on

2DMs are focused on graphene [32]. Only a few studies of radiation damage to WSe2 have been published [53,74,88]. Using focused beams of high energy electrons in a TEM (i.e. 60 kV and

100 kV), large numbers of vacancies can be generated to etch holes inside the 2D layers and fabricate nanoribbons. As the ribbons are thinned to a critical thickness of ~5Å, they transform in nanowires with a different crystal structure, a chemical formula of WSe, and metallic conductivity [88]. Helium ions can be used to preferentially remove selenium atoms, leading to a transition from intrinsic or p-type conductivity to n-type conductivity in irradiated devices [74].

16 2 The WSe2 transitions into an n-type insulator until a total exposure of 10 ions/cm , after which point it becomes metallic due to the large number of defects (see Figure 1.9).

Argon ions with energy of 5 to 10 kV have been used to study how the surface of bulk WSe2 is altered by energetic particle bombardment [53,89]. At low doses (i.e. 5 ∙ 1011 ions/cm2), individual impact craters can be resolved using scanning tunneling microscopy (STM; see Figure

1.7). The crater diameter is ~1.5 nm, and they are surrounded by walls with an outer diameter of

5 to 6 nm and height of ~0.3 nm. Areas not impacted by ions retain their original crystal structure. As the dose increases, the distance between craters decrease and they then begin to overlap. Beyond a dose of at least 1015 ions/cm2, individual craters cannot be resolved. Instead, the surface becomes corrugated and nanoscale hillocks can be observed in STM (see Figure 1.8).

The corrugation varies from 0.1 to 0.5 nm on a length scale of 1 to 5 nm. The original crystal structure appears to be maintained, but the surface layers become significantly disordered. This corrugated and deformed surface is observed to be stable at room temperature for weeks at a time, and the damage can be undone by annealing at 900K in ultra-high vacuum for 40 minutes

[89].

Impact crater

Figure 1.7 – Individual impact craters on the surface of WSe2 by bombardment via 5 keV argon ions and imaged via STM at an image size of 45 nm by 45 nm [53]

Corrugated surface

Figure 1.8 – Atomically-resolved STM image (8 nm by 8 nm) of a corrugated WSe2 surface [53]

Figure 1.9 – The change in resistance of a few-layer WSe2-based device on a SiO2/Si substrate due to helium ion bombardment [74]

The goal of this thesis is to advance the literature by further understanding the radiation tolerance of WSe2 against ionizing photons and charged particles. In particular, the response of this material to MeV energy particles needs to be known to determine its suitability for space electronics.

Chapter 2 – Experimental Procedure

2.1 Mechanical exfoliation of WSe2 WSe2 crystals were placed onto silicon dioxide and silicon carbide substrates by mechanical exfoliation. An optical image of the exfoliated flakes is shown in Figure 2.1. Bulk WSe2 crystals were deposited onto a piece of tape, and thinned by using the tape to peel the crystal apart. After thinning the crystals, the WSe2 flakes were transferred onto substrates such as silicon dioxide and silicon carbide by pressing the piece of tape onto the substrate and peeling it away, leaving thin flakes behind. In order to increase the number of flakes transferred by mechanical exfoliation, a thermal treatment [90] was used in certain experiments (see Section 2.5). After placing the tape on the substrate, the substrate was set on a hot plate set to 100°C in air. After 5 minutes of heating, the sample was removed from the hot plate and the tape was removed. The thickness of the exfoliated flakes was measured using an optical profilometer. The surface chemistry of the samples was analyzed using X-ray photoelectron spectroscopy (XPS). Using

XPS, it was determined that the heat treatment does not oxidize the WSe2. This data is shown in

Figure 2.2

SiC

WSe2

Figure 2.1 – Optical image of exfoliated WSe2 flakes

Figure 2.2 – XPS W 4f spectra of exfoliated WSe2 without and with the heat treatment. Only one doublet is observed, corresponding to WSe2. Intensity for all XPS measurements is in units of counts per second (CPS).

2.2 Metal-organic chemical vapor deposition (MOCVD) of WSe2 Nanoscale WSe2 domains and thin films were grown on various substrates (e.g. silicon carbide, sapphire, and graphite paper) by metal-organic chemical vapor deposition (MOCVD). The full details are given elsewhere [48], and a schematic of the reactor is shown in Figure 2.3. Briefly, gaseous pre-cursors of tungsten (e.g. tungsten hexacarbonyl – W(CO)6) and selenium (e.g. dimethyl-selenium – DMSe or hydrogen selenide – H2Se) are allowed to react at elevated temperatures (i.e. 600°C or 800°C) to form atomically-thick WSe2 domains and films. Hydrogen is used as a carrier gas for the vaporized W(CO)6 and DMSe. The domain size and surface coverage of the WSe2 crystals is dependent of the flow rate of the two pre-cursors and the carrier gas, the growth temperature, and the total growth time. For all growth runs, DMSe and H2Se are

o at room temperature and W(CO)6 is at 30 C. The specific parameters used for growth in each experiment will be discussed in their respective experimental sections.

W(CO)6 + DMSe / H2Se + H2

Figure 2.3 – Schematic MOCVD reactor

Prior to growth, substrates are cleaned by: sonication in the common solvents acetone and isopropyl alcohol individually for ten minutes; rinsing in de-ionized water; cleaning in 80°C

NanoStrip™ for 20 minutes; rinsing in de-ionized water; quick drying with a nitrogen gun.

Substrates are inspected with an optical microscope prior to growth. The domain size and surface coverage of MOCVD-grown WSe2 samples are characterized using a combination of atomic force microscopy (AFM) and scanning electron microscopy (SEM). AFM was also used to measure the height of individual domains. The thickness of the thin films was measured using primarily AFM, with optical profilometry used as a supplement, and surface chemistry was analyzed using XPS.

2.3 Stability of MOCVD-grown WSe2 in the XPS environment Samples of MOCVD-grown WSe2 were exposed to soft X-rays and low energy electrons for 12 hours to determine the sensitivity of this material to low energy radiation. WSe2 was grown on sapphire, alpha silicon carbide, and graphite paper using MOCVD. The sapphire substrate was annealed at 500oC for 5 minutes, and then the growth was conducted at 800oC for 20 minutes.

The flow rate of hydrogen gas, H2Se, and vaporized W(CO)6 were 196 sccm, 4 sccm, and 0.0006 sccm, respectively. The silicon carbide substrate was annealed at 500°C for 5 minutes, and then

o the growth was conducted at 600 C for 20 minutes. The flow rate of hydrogen gas, H2Se, and vaporized W(CO)6 were 225 sccm, 4 sccm, and 0.0006 sccm, respectively. The graphite paper substrate was annealed at 800oC for ten minutes, and then the growth was conducted at 800oC for

30 minutes. The flow rate of hydrogen gas, DMSe, and vaporized W(CO)6 were 420 sccm, 4.51 sccm, and 0.003 sccm, respectively.

Monochromatic aluminum k-alpha radiation with a photon energy of 1486.6 eV was utilized in a Kratos Axis Ultra spectrometer at an estimated flux [91] of 2.16 ∙ 1011 photons/mm2/s

(shown in Figure 2.4). A low energy electron flood gun was active for all samples, and the electrons are magnetically confined to the vicinity of the sample. Survey spectra are collected at the beginning and the end of X-ray exposure. The position of the valence band maximum is measured by collecting the valence band spectra and applying a linear fit to the leading, low binding energy, edge. Core level spectra of individual elements are collected during X-ray exposure. The analysis chamber is kept at ultra-high vacuum conditions (i.e. 10-9 Torr) during all measurements. Charge referencing is done with respect to the C 1s peak, which has a binding energy of 284.8 eV for the adventitious carbon on sapphire and silicon carbide and a binding energy of 284.5 eV for the graphite in graphite paper. To supplement the experimental analysis,

21 calculations of X-ray absorption and attenuation are done using XCOM: Photon Cross Sections

Database [92] and Geant4 Monte Carlo code [93].

Figure 2.4 – Kratos Axis Ultra spectrometer

2.4 Oxygen plasma treatment of WSe2 Plasma treatment of WSe2 was carried out using an M4L plasma etch tool (shown in Figure 2.5).

Samples were generated by mechanical exfoliation and MOCVD growth as detailed in Section

2.1 and Section 2.2. Exfoliated flakes were deposited onto silicon dioxide and silicon carbide substrates. MOCVD-grown flakes and films were grown on epitaxial graphene, sapphire, and silicon carbide substrates. For the epitaxial graphene and sapphire substrates, an anneal was done at 500oC for 5 minutes, and then the growth was conducted at 800oC for 20 minutes. The flow rate of hydrogen gas, DMSe, and vaporized W(CO)6 were 222 sccm, 3.73 sccm, and 0.0006 sccm, respectively. Multiple growth runs on SiC were carried out. For the first run, the substrate was annealed at 500oC for 5 minutes, and then the growth was conducted at 800oC for 20 minutes. The flow rate of hydrogen gas, H2Se, and vaporized W(CO)6 were 227 sccm, 7 sccm, and 0.0006 sccm, respectively. For the second run, anneal and growth conditions were the same.

The flow rate of hydrogen gas, H2Se, and vaporized W(CO)6 were 257 sccm, 7 sccm, and 0.0006 sccm, respectively. For the third run, anneal conditions were the same, and the growth was

o conducted at 800 C for 30 minutes. The flow rate of hydrogen gas, H2Se, and vaporized W(CO)6 were 257 sccm, 7 sccm, and 0.00036 sccm, respectively.

Figure 2.5 – M4L plasma asher tool. Image copyright PVA TePla America.

The standard plasma treatment condition was as follows: pressure of 550 mTorr; oxygen flow rate of 150 sccm; helium flow rate of 50 sccm; 100 W of RF power; 30 seconds of plasma exposure. To determine which aspects of oxygen plasma treatment have the greatest impact on

WSe2, the pressure, oxygen flow rate, RF power and exposure time were varied. Pressure was varied from 150 mTorr to 2 Torr, which were found to be the limits of the plasma etch tool.

Oxygen flow rate was varied from 0 sccm to 150 sccm. RF power was either at 100 W or 500 W.

Exposure time was varied from 3 seconds to 90 seconds. Helium gas flow rate is kept constant for all experiments. Changes in surface chemistry and oxidation were analyzed using a Kratos

Axis Ultra XPS analyzer with monochromatic aluminum k-alpha X-rays. Charge referencing is done with respect to the C 1s peak.

2.5 Proton and heavy ion irradiation of exfoliated WSe2

WSe2 flakes were mechanically exfoliated onto silicon carbide substrates as described in Section

2.1. The initial surface chemistry of these samples, as well as of bulk WSe2 crystals and bare SiC substrates, are analyzed using the Kratos Axis Ultra and PHI VersaProbe II XPS machines.

Samples are then mounted onto large (i.e. 20 cm by 20 cm) pieces of oxidized silicon from a silicon wafer. Mounting is achieved by coating the backside of the samples with polymethylmethacrylate (PMMA) resin, pressing them into the large silicon piece, and then curing at 80°C in air for 5 minutes. After curing the samples, they are stored in a stainless steel vacuum chamber than can achieve medium vacuum (i.e. < 100 mTorr) and transported to the

Michigan Ion Beam Laboratory for irradiation. The cured WSe2/SiC samples and storage chamber are shown in Figure 2.6.

Figure 2.6 – WSe2/SiC samples mounted onto SiO2 pieces and stored in a vacuum chamber

Irradiation with 2 mega-electron-volt (MeV) protons was done at various fluences – 1014 protons/cm2, 1015 protons/cm2, 1016 protons/cm2, and 1017 protons/cm2 – that are typical in the lifetime of space electronics [94]. Irradiation with 1 MeV, 200 keV, and 40 keV protons was also done at a fluence of 1016 protons/cm2. All heavy ion experiments are done at a fluence of 1016 ions/cm2, using 2.5 MeV Fe ions, 5 MeV Fe ions and 4 MeV Ag ions. Beams were raster scanned over an area of ~6 mm by ~6 mm at a 7° angle to avoid channeling effects. The beam current was generally in the range of 300 to 500 nA and the sample temperature was kept between 50°C and 100°C (a representative thermal image is shown in Figure 2.7). Beam experiments are done in high vacuum (i.e. <10-8 Torr).

Figure 2.7 – Thermal monitoring of proton and ion exposed samples

XPS analysis of proton bombarded samples was done using both XPS systems. The initial experiment used the Kratos Axis Ultra, while subsequent experiments used the PHI VersaProbe

II. XPS analysis of ion bombarded samples used only the PHI VersaProbe II system, which is shown in Figure 2.8. Charge referencing is done with respect to the C 1s peak regardless of XPS system. The thermal treatment (see Section 2.1) was used for all exfoliated samples that were analyzed using the PHI VersaProbe II. Potential changes in the optical properties of the WSe2 and SiC were recorded using a Perkin-Elmer Lambda 950 spectrophotometer. The transmittance of light with wavelengths ranging from 200 nm to 2000 nm was measured, and the converted to

25 light absorbance as a function of photon energy. Simulation of ion beam damage was done using

SRIM (The Stopping and Range of Ions in Matter) software [95].

Figure 2.8 – PHI VersaProbe II spectrometer

Chapter 3 – Stability of MOCVD WSe2 in the XPS environment

3.1 X-ray attenuation and energy absorption in WSe2 In general, when a beam of high energy photons such as an X-rays encounter a material, the beam is attenuated and gradually loses intensity. Attenuation proceeds through various phenomena, as shown in Figure 3.1 for MoS2 and WSe2. The primary attenuation mechanism is photoelectric absorption, where the absorbed energy of the photon breaks a bond in the material, producing an electron with a fixed energy and a positive ion with a vacancy. This mechanism dominates for soft X-rays, being several orders of magnitude more probable than either coherent or incoherent scattering, the other two major attenuation mechanisms for soft X-rays. After a photoelectric absorption event, an electron from a higher energy core shell may fill the vacancy while releasing an Auger electron. This de-excitation process leads to cascading electron ejection and a build-up of positive charge, which must be compensated in order to prevent sample damage [96,97]. For metallic samples, positive charging is compensated by electrons from surrounding unexposed regions. For semiconducting and insulating samples, the charging cannot be effectively compensated by surrounding areas. An electron flood gun is used in XPS analysis due to this issue (see Section 3.2 for further discussion).

Figure 3.1 – X-ray attenuation by the layered semiconductors MoS2 and WSe2.

The sensitivity of a bulk material to X-rays is quantified by its mass attenuation coefficient

(MAC; cm2/g), as shown in Figure 3.1, which is defined as the ratio of the attenuation coefficient of the material to its density. Knowledge of MAC is needed to understand differences in energy absorption and radiation damage between different bulk materials, and will depend on the photon energy, as well as the constituent elements and crystal structure of the irradiated material.

However, the utility of MAC, which is a bulk constant, for 2DMs is not obvious. Shown in

Figure 3.2 are the differences in X-ray attenuation and energy absorption between WSe2 and the three crystalline substrates investigated here: sapphire, silicon carbide (6H-SiC) and graphite paper (GP). These differences are calculated using the MACs of each material, which are listed in Table 3.1 along with the density and attenuation coefficient (in cm-1) based on Geant4 calculations and the band gap. Based solely on MAC, WSe2 would be expected to be more sensitive to X-rays than any of the underlying substrates used in these experiments. This sensitivity can be attributed mainly to the selenium atoms, which have a photon absorption edge near the aluminum k-alpha X-ray photon energy [92]. However, the atomic thickness of the as- grown TMDs limits the amount of X-ray absorption to a value estimated as ~0.3%. This is shown on the insets in Figure 3.2. Due to the limited absorption of X-rays within the as-grown

TMDs, it was hypothesized that X-ray damage to WSe2 would be dependent on the sensitivity of the substrate due to substrate-TMD interactions. As such, combinations of WSe2 and substrates that are more conductive (such as graphite paper) and have a lower MAC (such as 6H-SiC) should see the least damage. This is found not to be case for WSe2 (as discussed in Section 3.2 and 3.3) nor for MoS2, which was grown for comparative purposes (as discussed in Section 3.4).

Figure 3.2 – X-ray attenuation and absorption by WSe2 and the three substrates.

Table 3.1 – Table of mass attenuation coefficients, densities, attenuation coefficients and band gap values for WSe2 and the three substrates Mass Attenuation Density Attenuation Band Gap Material Coefficient (cm2/g) (g/cm3) Coefficient (1/cm) (eV) Sapphire 965 3.97 3831.2 ~8.5

6H-SiC 598.91 3.21 1922.5 ~3

WSe2 3319.57 9.35 31038 ~2 to ~1.2 Graphite 718.5 2 1437 0 Paper

3.2 Low energy electron flood gun utility in XPS In general, a low energy electron flood gun is active during XPS analysis for all samples to account for surface electron depletion and ensure identical analysis conditions. The electrons are confined to the vicinity of the samples by a magnetic lens. For a semiconducting or insulating sample, the extent of charging from X-rays may vary from spot to spot due to surface roughness.

For layered materials such as MoS2 and WSe2, there can also be differential surface charging due to differences in the number of layers. The effect of such differential charging is peak broadening. There may also be steps in potential between materials due to charge accumulation at interfaces. As such, the separation between peaks may be larger or smaller than expected by

29 several volts. The presence of low energy electrons prevents these effects by maintaining a uniform surface potential across the sample [98,99]. The low energy electrons from the flood gun may also induce a negative potential that compensates the surface charging induced by the X- rays. The driving force for gaining a fixed negative charge would be compensation of a pre- existing positive potential, either induced by X-rays or pre-existing due to defect or interface states. This negative surface potential would have the opposite effect on the measured binding energy in semiconductors, while still not affecting metallic materials. In an XPS measurement, the surface potential induced by X-rays and by the flood gun should effectively cancel out [98].

The goal of this study is to see if this remains the case over longer periods of time with more X- ray exposure and more electron ejection from the 2DMs and substrates.

3.3 Charging in WSe2 from soft X-rays and low energy electrons The combined impact of soft X-rays and low energy electrons on as-grown WSe2 is limited to small amounts of charging. Charging can be measured in XPS by measuring subtle changes in the binding energy of elemental peaks and the valence band maximum (VBM). The resolution of our XPS system is estimated as 0.1 eV, and so shifts of lesser magnitude than this can be considered insignificant. As said previously, it was expected that the least charging should be observed for the growth on graphite paper based on the substrate conductivity, with progressively more charging in 6H-SiC and sapphire based on their band gaps (3 eV for SiC

[100,101] and 8.5 eV for sapphire [102]). The XPS data shown in Figure 3.3 reveals that this is not what occurs. While there is more charging for WSe2 on sapphire than for WSe2 on SiC, the most charging for the WSe2 is observed for the growth on the graphite paper. Additionally, the charging on SiC should be considered insignificant. The charging of the W 4f peaks and the Se

3d peaks correspond well to each other, as shown in Figure 3.4, so the metal peaks will be the focus of this analysis.

(a) WSe2 on Graphite Paper (b) WSe2 on Sapphire (c) WSe2 on Silicon Carbide

Figure 3.3 – Shifts in binding energy of the VBM and W 4f electrons in WSe2

Figure 3.4 – Comparison between metal (W 4f) and chalcogen (Se 3d) peak shifts

By comparing the charging between the core level peaks and the VBM, it can be determined if a homogeneous surface charge is created [103]. A positive surface potential is generated in the

WSe2 grown on sapphire, but there is potentially some differential charging based on the different binding energy shifts of the W 4f (54 meV) and the VBM (125 meV). A positive surface potential of ~125 meV was also generated in the WSe2 sample grown on GP, and the

31 differential charging there was minimal. The graphite itself does not charge, so this shift can be fully attributed to charging in the WSe2. As said previously, there is essentially no charging for the WSe2 grown on SiC, making it the most stable in terms of X-ray induced charging.

The differences in WSe2 charging are attributed to differences in WSe2 domain size and morphology. Growth on sapphire produced isolated domains that cannot effectively dissipate charge. Growth on 6H-SiC and GP produced continuous films, but the growth on GP has fins that are characteristic of vertical growth (see Figure 3.5). Atoms can be preferentially removed from these exposed edges [104], leading to charged defect states in the 2DM. As such, the simplest means of controlling WSe2 sensitivity to low energy ionizing radiation appears to be tuning the growth rather than changing substrates.

(a) (b) (c)

WSe2 Al2O3 (d) MoS (e) 2 SiC

Al O MoS2 2 3 Figure 3.5 – Scanning electron microscopy images of the following samples: a) WSe2 on graphite paper; b) WSe2 on Al2O3 c) WSe2 on SiC; d) MoS2 on Al2O3; e) MoS2 on SiC

3.3 Comparison between WSe2 and MoS2 For comparative purposes, MoS2 samples were grown using powder vaporization on sapphire and 6H-SiC substrates [105], and also exposed to soft X-rays for the 24 hour period. Due to the elemental composition of MoS2, it only absorbs 30% of the dose that WSe2 would absorb due to its lower MAC of 1226 cm2/g. Since more X-ray energy would be absorbed in the substrate in this case than in the case of WSe2, the substrate would initially be expected to play more of a

32 role. However, this is still not observed to be the case. Both samples exhibit significant charging, with a magnitude greater than the WSe2 samples, and this is clearly shown by comparing the shifts in the metal binding energy between all samples (see Figure 3.6). However, more charging is observed for growth on SiC than for growth on sapphire.

Figure 3.6 – Binding energy shifts of the metal peaks of all studied 2DMs

Furthermore, the charging is towards lower binding energy, which indicates that the MoS2 is gaining negative charge from the low energy electrons. As such, it appears that the soft X-rays in general have minimal long-term impact on the 2D layers of WSe2 and MoS2 when low energy electrons are present. MoS2 is generally an n-type material, which means that it should have a free electron concentration and a fixed positive charge from its dopant species [63]. Most of the negative fixed charge gained from accumulating low energy electrons is accumulated up to a

15 2 15 fluence of 3.9x10 photons/cm for MoS2 on sapphire, and up to a fluence of 9.35x10

2 photons/cm for MoS2 on SiC. It is possible that these electrons are compensating for the pre- existing positive fixed charge that is present due to its doping type. A reduction of the VBM for

MoS2 on sapphire was observed, corresponding to the compensation of n-type doping (see Figure

3.7).

Figure 3.7 – Binding energy shift of the VBM for MoS2 on sapphire

Differences in the X-ray sensitivity between MoS2 samples grown on different substrates are best explained by differences in the domain size and surface coverage of the as-grown material as well. SEM images of the samples used in this study are shown in Figure 3.5. MoS2 domains in sapphire often have a domain size of at least 1 µm, and are often much greater than this (i.e. 10

µm or more). On the other hand, MoS2 domains on silicon carbide generally have a domain size between 500 nm and 800 nm. This greater sensitivity should be attributed to the reduced coupling between 2DMs grown via powder vaporization and the substrate. Greater coupling between graphene and the substrate is known to stabilize it against X-rays, and this effect has now been shown to potentially be universal to all 2DMs [106,107]. MoS2 samples grown on SiC exhibit more charging than MoS2 samples grown on sapphire, due to the reduced domain size and greater number of exposed edges. As said previously, the charging is due to defects at the edges of the as-grown domains. They appear to be more likely to form with reduced coupling

34 between the 2DM and the substrate and with smaller as-grown domains. As such, the growth of continuous films is best for reducing X-ray damage, especially if grown via MOCVD or a similar method with strong coupling.

Chapter 4 – Stability of exfoliated and MOCVD WSe2 in air and rapid oxidation with low energy plasma treatments

4.1 Stability of exfoliated and MOCVD WSe2 in air and medium vacuum Layered materials such as 2DMs are traditionally considered to be inert to ambient oxidation due to their layered structure and lack of dangling bonds. However, steps or other variations in layer number lead to exposed edges at which various oxide compounds can form [108]. Similarly, point defects such as selenium vacancies can allow for the formation of tungsten oxide [61]. As such, the stability of WSe2 against atmospheric oxygen must be determined prior to any radiation studies where transportation and storage of the samples in oxygen-containing environments is required. This was done to subtract out any impact that air exposure and storage at medium vacuum might have over the timeline of the proton and heavy ion experiments (a few days).

Ideally, the samples would not oxidize at all in this timeframe, but completing these studies would ensure that any oxidation that might occur can be factored into our experimental results.

As was expected, the samples are found not to oxidize over the short timeframe of three days.

Shown in Figure 4.1 are the tungsten 4f spectra of exfoliated WSe2 on 6H-SiC before and after exposure to oxygen either from air or from storage in medium vacuum. In all cases, the only

7/2 major feature observed is the doublet corresponding to WSe2. The binding energy of the W 4f peak is initially at 32.34 ± 0.03 and shifts slightly upwards to 32.39 ± 0.02 eV, remaining consistent with literature values [69,109]. A similar study was done using WSe2 grown on 6H-

SiC substrates via MOCVD, and the resulting XPS spectra for that study are shown in Figure

4.2. The spectral shape is constant here as well, indicating no oxidation. When peak shifting does occur, it is minor and consists of a downwards shift (~50 meV), except for one sample stored in air that a more significant shift of ~290 meV.

Figure 4.1 – Changes in the W 4f spectra of exfoliated WSe2 due to storage in medium vacuum or air exposure for three days. The small contribution of 5p3/2 core shell electrons is not shown.

3/2 Figure 4.2 – Changes in the W 4f and 5p spectra of MOCVD-grown WSe2 due to storage in medium vacuum or air exposure for three days.

4.2 Oxidation of exfoliated and MOCVD WSe2 when exposed to low energy plasma

On the other hand, the controlled oxidation of WSe2 can be useful for specific applications.

Considering the example of silicon, the ability to reliably grow an insulating oxide on top of a semiconducting material has allowed it to become a staple of electronics manufacturing and research. As discussed in Section 1.2, the band gap of WSe2 is highly dependent on the number

37 of layers as the material approaches single layer thicknesses, going from 1.2 eV as multilayer to

~2.2 eV as single layer. Stoichiometric tungsten trioxide (WO3) is a wide band-gap semiconductor (Eg ~ 3 eV [110,111]), that often has n-type properties due to oxygen vacancies

[112] and has a large work function generally in excess of 5 or 6 eV [113]. Based on these properties, tungsten oxide is being explored for conductive oxide applications rather for insulating purposes. This material is attractive as a contact for p-type semiconductors in its metallic sub-stoichiometric form of WOx [114], and forming in-plane heterostructures between

WOx and WSe2 has been shown to boost the device mobility [115].

The advantage of being able to grow a conductive oxide layer directly on WSe2 cannot be understated, due to the issues of insufficiently large metal work functions for low-resistance ohmic contacts to ultrathin WSe2 and the presence of Fermi level pinning between metal contacts and 2DMs [39,41]. Ideally, this oxidation process should be controllable, uniform, and rapid for integration with device fabrication. Previously explored processes such as ozone treatment [109] or thermal oxidation [116] are controllable, but not rapid. The ozone treatment also results in layer-by-layer conversion of WSe2 into an oxide, so uniform layers can be generated over time.

A process that also has controllable layer-by-layer conversion and is rapid would be ideal.

For example, oxygen plasma treatment is a rapid means of cleaning [117] and functionalizing

[118] surfaces. The plasma oxidation process can be completed in a few minutes or even seconds due to the higher activity of oxygen ions in the plasma [117], and presence of the helium ions that can break additional bonds. Oxidation of WSe2 occurs primarily during plasma oxidation.

When the plasma is generated, low energy oxygen and helium ions are formed by stripping electrons from the oxygen and helium gas present in the chamber. These ions can then interact with the WSe2. Helium can only break bonds because of its inert nature. However, oxygen ions

38 can both break bonds in WSe2 and react with it, forming W-O and Se-O bonds. The compounds that form are WOx and SeOx. Figure 4.3 and 4.4 show the changes in the XPS spectra of W 4f (in

Figure 4.3) and Se 3d (in Figure 4.4) in the exfoliated and MOCVD WSe2 grown on 6H-SiC that indicate oxidation due to the standard plasma treatment discussed in Section 2.4.

Figure 4.3 – W 4f spectra of exfoliated and as-grown WSe2 before and after the standard plasma treatment. Units on the vertical axis are counts per second (CPS) normalized by the intensity of the major W 4f peak for WSe2.

Initially, the tungsten spectrum shows a tall doublet (W 4f) accompanied by a single, short

3/2 peak (W 5p ). This is characteristic of the tungsten in as-grown WSe2 with no detectable surface oxidation. Meanwhile, the selenium spectrum shows a single doublet (Se 3d), also characteristic of WSe2. After the plasma treatment, an additional doublet appears in the tungsten spectra, indicative of the oxide. In the selenium spectra, a single degenerate peak with a binding energy ~5 eV greater than the binding energy of Se in WSe2 appears that is indicative of the oxide. SeOx is observed to be stable after plasma treatment as it is a room temperature process; however, it decomposes into selenium and oxygen vapors at elevated temperatures [119], while the tungsten oxide remains. It should be noted that this peak was not observed on samples exfoliated onto SiO2, which is the typical substrate used in literature. In addition to this oxide peak, a second selenium doublet that nearly overlaps with the original doublet is observed in

MOCVD-grown samples with an as-of yet unknown identification. One possibility is elemental selenium clusters stabilized at defects, but this has yet to be verified. It is also not yet clear why such features are observed only in the MOCVD-grown WSe2. Control samples which are not exposed to the plasma show no changes in the extent of oxidation (see Figure 4.5 for W 4f and

4.6 for Se 3d).

Figure 4.5 – W 4f spectra of the control samples of exfoliated and as-grown WSe2. Units on the vertical axis are CPS normalized by the intensity of the major W 4f peak for WSe2. Note that the as-grown material starts out with some oxide.

Figure 4.6 – Se 3d spectra of the control samples of exfoliated and as-grown WSe2. Units on the vertical axis are CPS normalized by the intensity of the major Se 3d peak for WSe2.

The plasma treatment process is also observed to uniformly convert the top layer of the

MOCVD-grown WSe2 into an oxide after 30 seconds. The AFM scans used to make this determination are shown in Figure 4.7. Initially, the samples consist of individual triangles

(Figure 4.7a and 4.7c). This remains the case after treatment (see Figure 4.7b and 4.7d); however, there is a uniform increase in the sample height as measured in AFM by ~1.5 nm, corresponding to about two unit cells of tungsten oxide, similar to what was observed using the ozone treatment [109].

(a) (b)

WSe2

Oxide- coated

WSe2 (c) (d)

WSe2

Figure 4.7 – AFM scans of the following samples: (a) MOCVD WSe2 grown on epitaxial graphene (EG) prior to treatment; (b) MOCVD WSe2 on EG after treatment; (c) MOCVD WSe2 grown on 6H-SiC prior to treatment; (d) MOCVD WSe2 on SiC after treatment. The treatment used here is the standard treatment discussed in Section 2.4.

As stated previously, the ions in the plasma will break bonds in the WSe2 layers. This reactive ion etch process is highly directional, and so the ions will interact with the sample primarily in the vertical direction. That is, the ions impact the WSe2 from the top layers down, and so oxidation will occur layer-by-layer, again like ozone. However, due to their low energy, the ions do not appear to penetrate underneath the first layer. As such, oxidation would be expected to stop there. This etch process also impacts the entire top layer at once, making it uniformly defective. As such, oxidation proceeds all at once, rather than spreading out from initial defect

42 sites and edges on the top-layer [109]. The MOCVD-growth process is highly anisotropic and leads to 100nm to 1+ µm domains of single to few layer (i.e. < 10 nm thick) WSe2, while the exfoliation process leads to thicker flakes (i.e. 1 µm on average; see Section 6.2 for more details) with less surface coverage. As such, more oxidation would be expected for the MOCVD-grown samples, which is what is observed based on Figures 4.3 and 4.4.

4.3 Impact of plasma parameters and exposure time on the oxidation process

Tuning plasma parameters would be expected to alter the strength of the plasma, and thus impact the oxidation process. However, investigations regarding the impact of oxygen flow rate, exposure time, chamber pressure and RF power on the oxidation of WSe2 reveal that this is generally not the case. Changing the plasma exposure time appears to be the only important factor in tuning the amount of tungsten oxide that forms, and only for the exfoliated material. All others have no obvious impact. The same is true for selenium oxide. For these experiments, only one parameter was changed at a time; the rest were kept constant. Samples not subjected to plasma exposure have no significant changes in oxidation for either the exfoliated or as-grown material (as shown in Figures 4.5 and 4.6).

As discussed previously, the plasma can only damage and oxidize the topmost layer of WSe2.

Based on this, it would be expected that the amount of tungsten oxide generated with remain constant after a certain amount of time with plasma exposure. Shown in Figure 4.8 is a plot showing how the percentage of tungsten atoms bonded to oxygen and the percentage of selenium atoms bonded to oxygen vary by varying the plasma exposure time. It is apparent that the expected saturation occurs for the exfoliated material, but not for the MOCVD-grown material.

In the exfoliated material, there is slightly less tungsten oxidation after 10 seconds than there is

43 after 30 or 90 seconds. This suggests that the oxidation of the exfoliated WSe2 due to plasma exposure is self-limiting, as discussed previously. For the exfoliated sample on SiO2, there is no detectable SeOx. For the exfoliated sample in SiC, the amount of SeOx generated after 10 seconds and 30 seconds are comparable and less than the amount generated after 90 seconds. It is proposed that the SiC substrate plays a role in stabilizing the selenium oxide, but determining its exact contribution requires further investigation.

Figure 4.8 – Changes in the concentration of tungsten oxide and selenium oxide due to low energy plasma exposure when the treatment time is varied.

For the MOCVD-grown material on SiC, there is no clear trend between treatment time and oxidation. The MOCVD-grown films and domains behave the same as the exfoliated flakes for

10 seconds or more, just with greater oxidation. However, the most tungsten oxidation is actually after 3 seconds of exposure. The most selenium oxidation is after 90 seconds, but the least amount is also after 10 seconds. The reason for this is not obvious. The sample that was exposed for only 3 seconds and the one exposed at 90 seconds start with the most initial tungsten oxidation (9.21% vs. 9.84%), and are comparable after the treatment (80.76% vs. 77.56%). It appears that the highly anisotropic few-layer samples of MOCVD-grown WSe2 are so much

44 more sensitive to induced damage and oxidation from the reactive ions that the differences in plasma exposure time are not significant, and any exposure will lead to large amounts of oxidation. In stark contrast, storage of a plasma-treated sample in air by two weeks does not increase the extent of oxidation by much at all (see Figure 4.9). As stated earlier in Section 4.1, the WSe2 is relatively stable against ambient oxidation, and the plasma treatment only de- stabilizes the remaining material by a small amount due to the layer-by-layer oxidation process.

This suggests that plasma treatments can be used to make air-stable oxide coatings for WSe2.

Figure 4.9 – Changes in XPS spectra of MOCVD-grown WSe2 on SiC due to plasma treatment and subsequent storage in air. There is no significant change in the extent of oxidation with air exposure (72.55% WOx after treatment vs. 73.55% WOx after storage in air).

Shown in Figure 4.10 is a plot showing how the percentage of tungsten atoms bonded to oxygen and the percentage of selenium atoms bonded to oxygen vary by varying the oxygen flow rate. Exfoliated samples start of pristine with no oxide contributions, while the samples grown via MOCVD start off with some tungsten oxide. For exfoliated samples, there appears to be a dependence on the substrate used. Samples exfoliated onto SiO2 see more tungsten oxide

45 formation with an increase in the oxygen flow rate (R2 = 0.98). No selenium oxide formation is observed. For samples exfoliated onto SiC, there is no correlation between tungsten oxidation and oxygen flow rate. However, there is a correlation between reducing the oxygen flow rate and increasing the amount of selenium oxide (R2 = 0.97). There is no correlation between increased oxygen flow rate and increased tungsten oxidation or selenium oxidation for MOCVD-grown samples. Additionally, there is significant oxidation even when no oxygen is in the chamber.

This oxidation can come from two potential sources: residual oxygen in the plasma chamber, and oxygen from the ambient when exposed immediately after plasma treatment. The changes in the oxygen flow rate would change the relative concentrations of helium to oxygen in the plasma, and it can be said that changes in the concentration can only tune the amount of oxidation in specific cases.

Figure 4.10 – Changes in the concentration of tungsten oxide and selenium oxide due to low energy plasma exposure when the flow rate of oxygen gas is varied.

Shown in Figure 4.11 and 4.12 are plots showing how the percentage of tungsten atoms bonded to oxygen and the percentage of selenium atoms bonded to oxygen vary by varying the chamber pressure and the RF power. Within the range of pressures that can be sustained by this

46 system, changes in the chamber pressure are not effective in tuning the amount of tungsten oxide. It can tune the amount of selenium oxide, but the samples shown have different trends.

When the pressure is increased, there are two competing effects. More oxygen atoms are present that can be converted into reactive ions, and so the amount of oxidation is expected to increase.

More helium ions would aid in this. On the other hand, increasing the pressure also reduces the mean free path of the ions, and the reduced number of interactions between the ions and WSe2 makes oxidation less likely [120]. The samples grown on the two different faces of the SiC are differently influenced by this competition, though it remains to be seen if there is a fundamental cause behind this different response. It should be noted that the tunability is about 20%. For the impact of RF power, only preliminary studies were done. It seems as though increasing the plasma power may have some impact, but this is considered small.

Figure 4.11 – Changes in the concentration of tungsten oxide and selenium oxide due to low energy plasma exposure when the chamber pressure is varied. Control samples are at zero.

Figure 4.12 – Changes in the concentration of tungsten oxide and selenium oxide due to low energy plasma exposure when the RF power is varied. Control samples are at zero.

Chapter 5 – Impact of Proton Irradiation on Exfoliated WSe2

5.1 Effect of 2 MeV protons at various fluence levels The 2 MeV protons are calculated to penetrate the WSe2/SiC samples down to a depth of ~32 µm using SRIM [95]. These particles transfer energy to the sample via inelastic scattering off electron clouds or by elastically scattering off atomic nuclei. Inelastic scattering off electrons, or electronic stopping, is dominant at the start of the proton track, and leads to ionization and charging. Elastic scattering off nuclei, or nuclear stopping, is dominant near the end of the proton track, and leads to atomic displacement and vacancy generation. Ionization is thus the dominant damage mechanism that would be expected in the WSe2 surface layer, while displacement damage would be prominent deep in the SiC. Charging effects can be detected in XPS by shifts in the binding energy of the core level electrons in the WSe2, due to the sensitivity of the measured binding energy to local electric fields.

Shown in Figure 5.1 are the shifts due to 2 MeV proton bombardment of the W 4f and Se 3d core level peak binding energies relative to their initial position. The numerical values are listed in Table 5.1 for convenience. As stated in Section 3.2, the resolution of our XPS systems is 0.1 eV. We thus define a significant peak shift here to be 0.2 eV or more. Control samples were prepared and have no significant peak shifts, as expected. Samples exposed to radiant fluences of

1014 protons/cm2 and 1015 protons/cm2 also have peak shifts that are not significant. Samples exposed to radiant fluences of 1016 protons/cm2 and 1017 protons/cm2 are observed to have significant increases in the measured binding energy of core shell electrons. The exact numerical values are observed to have a slight variation between the different experiments. These differences should be attributed to spatial inhomogeneity in the samples, as they are qualitatively similar.

Figure 5.1 – Binding energy shifts in W 4f and Se 3d core shell levels due to 2 MeV protons at various fluence levels

Table 5.1 – Tabulated binding energy shift values used for Figure 5.1 Fluence (protons/cm2) W 4f binding energy shift (eV) Se 3d binding energy shift (eV) 0 0.0217 0.0271 0 0.0624 0.0754 0 -0.0991 -0.0333 0 0.0413 0.0705 0 -0.0958 -0.1447 0 -0.0736 -0.0204 1014 0.0377 0.0349 1015 -0.035 -0.0605 1015 0.1031 0.1046 1016 0.3465 0.3487 1016 0.4442 0.4198 1016 0.6224 0.6093 1017 0.2911 0.2949 1017 0.8206 0.8546

Shifts of the W 4f peaks and Se 3d peaks correspond well to each other, and indicate that the charge is uniformly distributed in the WSe2 layers. The W 4f spectra of a representative sample that was exposed to 2 MeV protons at 1016 protons/cm2 and analyzed using the PHI VersaProbe

II XPS system are shown in Figure 5.2. No oxidation was observed for any sample under any level or proton beam exposure. This lack of oxidation is expected, based on the light mass of protons, their deep penetration depth, and the dominance of ionization over displacement at the

50 sample surface. It is clear that the extent of charging depends non-linearly on the fluence, and requires the proton fluence to surpass a threshold value before it impacts the WSe2 properties.

This charging is attributed to a combination of ionization in the WSe2 and charge transfer from the SiC into the WSe2. Due to the small concentrations of WSe2 on the SiC surface (i.e. < 5%), the charging should primarily be caused by damage in the SiC that is discussed more in Section

5.4. The explains the stability of the built-up charge in the WSe2, as it has yet to dissipate after two weeks (see Figure 5.2)

Figure 5.2 – W 4f and 5p3/2 core shell electron spectra from before, after, and two weeks after exposure to 1016 protons/cm2 with an energy of 2 MeV

The threshold dose of 1016 protons/cm2 measured here is noted to be high when compared to studies of proton damage to MoS2 devices, which are significantly impacted by doses as low as

1013 protons/cm2 [121]. In both cases, this damage is not from direct damage to the 2DM surface layers, but due to damage in the substrate that has a secondary impact on the surface material. As such, the difference should be attributed to the difference in radiation hardness between the two substrates, suggesting that SiC is much more desirable as a substrate for space electronics than

SiO2. However, SiC cannot be used as a gate dielectric to WSe2 due to its band gap of 3 eV, which is not large enough to minimize leakage currents by maintaining large band offsets of 1 eV or more [27,122]. As such, SiC can be used as a semi-insulating substrate, but a radiation

51 tolerant top-gate dielectric that is compatible with WSe2 must still be found. The details of band offsets between WSe2 and SiC and consequences for future devices are discussed in more detail in Section 5.5.

5.2 Effect of 1 MeV, 200 keV, and 40 keV protons at a fluence of 1016 protons/cm2 The effect of 1 MeV protons is almost the same as that of the 2 MeV protons, as shown by changes in the binding energy of core level peaks in the W 4f and Se 3d spectra (see Figure 5.3 and Table 5.2). The peak shifts are slightly larger than in the case of 2 MeV protons, but not significantly so. The similarity of these results would be expected based on the expected penetration depth of these protons of ~11 µm. On the other hand, the charging induced by 200 keV at the same fluence is judged to not be significant, and the charging induced by 40 keV protons is just barely so (see Figure 5.3 and Table 5.2). These protons have expected penetration depths of 1.2 µm and 0.258 µm, and would be expected to transfer more energy directly to the

WSe2. However, the reduced dimensions of the WSe2 flakes minimize the amount of energy that can be transferred to them. Additionally, less energy overall is transferred into the SiC substrate due to the reduced proton energy. As such, the impact of 200 keV and 40 keV protons at a fluence of 1016 protons/cm2 is similar to the impact of 2 MeV protons at a fluence of 1015 protons/cm2 or less – that is; not significant.

Figure 5.3 – Binding energy shifts in W 4f and Se 3d core shell levels from proton exposure at various energies at a fluence of 1016 protons/cm2

Table 5.2 – Tabulated binding energy shift values used for Figure 5.3 Proton energy (MeV) W 4f binding energy shift (eV) Se 3d binding energy shift (eV) 0.04 0.2394 0.2137 0.2 0.1046 0.0721 1 0.4665 0.4333 1 0.6732 0.6715 2 0.3465 0.3487 2 0.4442 0.4198 2 0.6224 0.6093

5.3 Proton impact on the SiC substrate and its optical properties As stated previously, there is charging in the SiC due to proton-induced ionization. However, much of this charge that was near the surface appears to be transferred into the WSe2. As such, it is difficult to determine if shifts in the binding energy of the core level peaks of the SiC are significant or not. These shifts are shown in Figure 5.4 and 5.5 for the silicon (Si 2p) and the carbon (C 1s) in the substrate. Samples without any WSe2 were also exposed to 200 keV and 2

MeV protons at 1016 protons/cm2, and that data is shown in Figures 5.4 and 5.5. Consistently significant shifts are only seen at a fluence of 1017 protons/cm2 for the 2 MeV protons, and at a fluence of 1016 protons/cm2 for the 200 keV protons.

Figure 5.4 – Binding energy shifts in Si 2p and C 1s core shell levels due to 2 MeV protons at various fluence levels

Figure 5.5 – Binding energy shifts for Si 2p and C 1s core shell levels from proton exposure at various energies at a fluence of 1016 protons/cm2

More consistent than the binding energy shifts, are changes to the optical properties of the

SiC. Samples exposed to a fluence of 1016 protons/cm2 are observed to darken and turn translucent, and samples exposed to a fluence of 1017 protons/cm2 are observed to turn black and opaque. In contrast, control samples and those exposed to lower doses remain transparent due to the wide band gap of SiC and nanoscale dimensions of the 2DMs (see Figure 5.6). This color change corresponds to a change in the light absorbance properties of the SiC, as shown in Figure

5.7 for a sample exposed to 1016 protons/cm2. Initially, there is only one absorption edge observed in the SiC – that of the band gap, with a value around 3.1 eV. After irradiation at 1016 protons/cm2, this edge is observed to decrease slightly to 3 eV. More significant is the shoulder that emerges at low energy, with an absorption edge around 1.1 eV. Visible light consists of photons with wavelengths ranging from 390 nm to 700 nm, which is and energy range of ~3.2 eV to ~1.77 eV. As such, there is much stronger absorbance of visible light photons in the SiC due to the physical damage, and so its changes color from transparent to black.

Figure 5.6 – (a) Camera picture of samples prior to irradiation; (b) camera picture of samples after irradiation. Samples were exposed to 2 MeV protons unless stated otherwise.

(a) (b) (c)

Figure 5.7 – Absorbance spectra of the SiC samples with and without irradiation

This change is attributed to displacement damage deep in the SiC that occurs when protons interact with carbon and silicon atoms via nuclear stopping. Proton damage to SiC generally

55 results in point defects such as silicon vacancies, carbon anti-sites and carbon-silicon di- vacancies [123,124]. Of these, silicon vacancies are known to be dominant in the semi-insulating

SiC used in this work. Induced at the same time is a deep state in the band gap located at about

1.2 eV below the conduction band edge [123,125], which is noted to be a deep acceptor [126].

This deep level, often referred to as the R center in the literature, is attributed to these vacancies

[123]. As such, enhanced light absorption due to a region deep in the SiC containing large clusters of vacancies is the mechanism responsible for the black color of the irradiated material.

5.4 Proton impact on WSe2 and SiC valence band offset In addition to the samples of exfoliated WSe2 on SiC and bare SiC, samples of bulk WSe2 were prepared and exposed to the proton beam. Samples were either not exposed to the beam or were exposed to 200 keV or 2 MeV protons at a fluence of 1016 protons/cm2. There is no oxidation in these samples due to beam exposure, and the measured binding energy shifts of the core level peaks are listed in Table 5.3. Due to the wide variance in the peak shifts, no conclusions were drawn using this data.

Table 5.3 – Tabulated peak shifts for bulk single crystal WSe2 Proton energy (MeV) W 4f binding energy shift (eV) Se 3d binding energy shift (eV) Control -0.2842 -0.2877 Control -0.5713 -0.5671 0.2 -0.5423 -0.5484 2 -0.3768 -0.3675 2 0.183 0.1899

Using all the available XPS data, the valence band offset (VBO) between WSe2 and SiC can be measured to determine how the band alignment is influenced by proton bombardment. The

VBO is measured using the following equation [60,127,128]:

푊푆푒2 푊푆푒2 6퐻−푆푖퐶 6퐻−푆푖퐶 ∆퐸푣 = ∆퐸퐶퐿(푖) + (퐸푊 4푓 − 퐸푉퐵푀 ) − (퐸푆푖 2푝 − 퐸푉퐵푀 )

Here, ΔECL(i) refers to the energy difference between the W 4f for WSe2 and Si 2p for 6H-SiC at the interface between the exfoliated WSe2 and the SiC substrate. All other energies refer to the binding energies measured using the bulk material. The VBM was measured using the linear fit approach mentioned in Section 2.3. Determination of the VBO requires six measurements, each of which has an estimated uncertainty of 0.1 eV. As such, the overall uncertainty in the VBO calculation is determined to be 0.24 eV. The initial VBO between WSe2 and SiC and changes to it due to proton beam exposure are tabulated in Table 5.4. These VBOs and the changes to them are also plotted in Figure 5.8 for convenience. The binding energy separation data used to calculate the VBOs is tabulated in Table 5.5. In general, the initial value of the VBO is between

1.2 eV and 1.4 eV. Using the band gaps of multilayer WSe2 (1.2 eV) and 6H-SiC (3.1 eV), it can be determined that these two materials have type I band alignment, and that the experimental value of the conduction band offset (CBO) is between 0.5 eV and 0.7 eV. Given electron affinity values of ~4 eV for WSe2 and ~3.3 eV for SiC, the ideal CBO and VBO are ~0.7 eV and ~1.2 eV, which are close to our measured values.

Table 5.4 – Tabulated initial valence band offsets between WSe2 and SiC and changes induced to the VBO by proton exposure at 1016 protons/cm2 Valence Band Offset (eV) [initial] Proton energy (MeV) Valence Band Offset (eV) [Final] 1.22 2 0.71 1.29 2 1.06 1.21 0.2 1.44 1.40 Control 1.50

Figure 5.8 – Plot of initial valence band offsets between WSe2 and SiC and changes induced to the VBO by proton exposure at 1016 protons/cm2

Table 5.5 – Core level separation data used to calculate the VBO between WSe2 and SiC State Energy (MeV) W 4f 7/2 – VBM (eV) ΔE (core levels) (eV) Si 2p 3/2 – VBM (eV) Initial 2 31.7436 68.4043 98.9273 Initial 2 31.7014 68.383 98.7989 Initial 0.2 31.7211 68.221 98.7371 Initial Control 31.7683 68.3337 98.7035 Final 2 31.8038 67.8764 98.969 Final 2 31.7597 68.1295 98.8301 Final 0.2 31.7793 68.4055 98.7415 Final Control 31.7812 68.3983 98.6771

As expected, the control samples do not have significant changes to the VBO, which is observed to increase by ~0.1 eV. Exposure of 1016 protons/cm2 to 200 keV and 2 MeV protons appears to have opposing effects on the VBO, where the 200 keV protons increase the VBO and the 2 MeV protons reduce it. The changes appear to be roughly stable appear two weeks of storage in medium vacuum, implying that the alteration of the VBO is permanent. In all cases, the primary cause of the shift in VBO is due to changes in the core level peak separation between

WSe2 and SiC. As mentioned in Sections 5.2 – 5.4, significant core level binding energy shifts

16 are observed in the WSe2/SiC heterostructure for WSe2 with 2 MeV proton exposure at 10 protons/cm2 and for SiC with 200 keV proton exposure at 1016 protons/cm2. The other material does not have a significant shift. When the bulk materials see peak shifts, the VBM sees a similar

58 shift due to the formation of uniform surface potentials. As such, the energy separation between the core level electrons and the VBM in the bulk material is rather constant.

The changes in VBO have consequences for the expected electronic properties of the

WSe2/SiC heterostructure. Based on the changes in band alignment and their likely origin, changes in the doping levels of either the WSe2 or SiC are expected. For exposure to 2 MeV

16 2 protons at 10 protons/cm , it is expected that the WSe2 becomes significantly more n-type relative to its original doping level. As such, the WSe2 may be slightly compensated, fully compensated, or transition into an n-type material, depending on the original position of the

Fermi level. For exposure to 200 keV protons at 1016 protons/cm2, it is expected that the SiC becomes significantly more n-type relative to its original doping level. This could lead to excess conduction through the SiC that would increase the total device current and lead to extra heat generation. In both cases, the other material should be affected, but not significantly so. In either case, the device would be negatively impacted, and permanently so. It is expected based on the results discussed in Section 5.3 that 40 keV protons would not impact the VBO significantly while 1 MeV protons would be the same as 2 MeV protons.

If SiC were used as a gate dielectric in a device, we would expect that 200 keV proton exposure would lead to an increase in the electron leakage current between the WSe2 and SiC, while the 2 MeV proton exposure would lead to an increase in the hole leakage current. These increases would likely be permanent, although the magnitude may fluctuate with time. However, doing so is not recommended based on the band alignment and changes to it. Using SiC as a substrate over other materials such as SiO2 is recommended, but a superior top-gate dielectric with similar or better radiation tolerance must be found through investigations like this one.

5.5 Comparisons between WSe2 and MoS2 For comparative purposes, MoS2 samples were also prepared by mechanical exfoliation onto silicon carbide substrates and were exposed to proton beams in the following conditions: 2 MeV and 1016 protons/cm2, 200 keV and 1016 protons/cm2, and 40 keV and 1016 protons/cm2. The measured shifts in the binding energy of the Mo 3d and S 2p core shell electrons are shown in

Figure 5.9 and listed in Table 5.6. The corresponding binding energy shifts in the Si 2p and C1s peaks for these samples is shown in Figure 5.10. Unlike WSe2, there are no consistent trends in the charging behavior of MoS2. Regardless of exposure to the proton beam, significant binding energy shifts – and thus significant amounts of charging – can be observed in the exfoliated

MoS2. Furthermore, some of these shifts are negative; whereas for WSe2, all shifts were positive.

As such, the impact of proton exposure to exfoliated MoS2 on SiC requires further study.

Figure 5.9 – Binding energy shifts for Mo 3d and S 2p core shell levels from proton exposure at various energies at a fluence of 1016 protons/cm2

Table 5.6 – Tabulated binding energy shifts for exfoliated MoS2 Proton energy (MeV) Mo 3d binding energy shift (eV) S 2p binding energy shift (eV) 0 -0.2316 -0.2335 0 0.1002 0.0726 0 0.2481 0.2544 0.04 -0.2015 -0.2001 0.2 -0.1019 -0.084 2 0.165 0.1682 2 0.4592 0.4988

Figure 5.10 – Binding energy shifts for Si 2p and C 1s core shell levels from proton exposure at 16 2 various energies at a fluence of 10 protons/cm from the MoS2/SiC samples

Using the available XPS data and a bulk sample of MoS2, the VBO between MoS2 and SiC can also be calculated, and is measured as 0.67 eV. Given a multilayer band gap of 1.29 eV for

MoS2 [129], we determine that these materials have a type I band alignment and the CBO can be extracted as 1.04 eV. However, MoS2 and WSe2 have approximately the same electron affinity of ~4 eV [130,131], meaning the ideal CBO should be around 0.7 eV. As such, there is some charge transfer occurring at this interface that is likely influencing the charge transfer that would be induced by the protons. Further study of MoS2 is highly recommended because of this charge transfer effect.

Chapter 6 – Impact of Heavy Ion Irradiation on Exfoliated WSe2

6.1 Heavy metal ion damage to exfoliated WSe2 Exposure of materials to MeV energy ions damages them through a combination of electronic stopping (inelastic collisions with electrons) and nuclear stopping (elastic collisions with nuclei).

The energy transfer from these two mechanisms leads to ionization and displacement that damages semiconductors, and is known to depend on ion species and energy. The expected damage can be characterized by the displacements per atom (dpa), which can be calculated using

SRIM. Based on the dpa of an ion beam with a fluence of 1016 ions/cm2, it was hypothesized that the most damage to WSe2 should be caused by the 4 MeV Ag ions (dpa = 24.4). About half as much displacement damage should be caused by 2.5 MeV Fe ions (dpa = 11.7), which should cause approximately twice the damage of the 5 MeV Fe ions (dpa = 6.9). This displacement damage would correspond to changes in the XPS spectra of the WSe2 such as selenium desorption and tungsten oxidation.

However, no such trend in the XPS data was observed. Figure 6.1 shows the tungsten 4f spectra of exfoliated WSe2 before and after exposure to the three different ion beams. The three spectra have qualitatively similar features: initially, they show features corresponding only to

WSe2; after exposure, features corresponding to WOx can be observed. The ratio of W-O bonds generated to W-Se bonds is measured as 0.51 ± 0.05, indicating that the damage induced by the three ion beams is very similar. This value increases after two weeks of storage in the vacuum chamber to 0.64 ± 0.05, a 25% change. This indicates that the damaged WSe2 is now much more sensitive to ambient oxygen than the pristine material, which did not oxidize in this time frame based on analyzing control samples. The amount of selenium desorption is also similar in the three cases, and the changes to the selenium 3d spectrum are shown in Figure 6.2. The measured

62 ratio of Se:W in the samples starts out as 1.96 ± 0.05, indicating the good quality and stoichiometry of the exfoliated material. After irradiation, the Se:W ratio is degraded to 0.68 ±

0.05, due to the increased volatility and reduced mass of selenium when compared to tungsten.

The ratio degrades further with storage in vacuum to 0.59 ± 0.02. The WSe2 would be expected to fully convert to a tungsten oxide phase over much longer time scales based on this oxidation rate.

Figure 6.1 – W 4f spectra of as-exfoliated samples (black) and after irradiation (red)

Figure 6.2 – Se 3d spectra of as-exfoliated samples (black) and after irradiation (red)

The differences between the three samples are shown in Figure 6.3. The differences between the samples are not as stark as would be suggested by the differences in dpa. It is proposed that at the high fluence used in this experiment, the differences in dpa between the three ions are masked by the large number of ions bombarding the WSe2. Similar structural damage results in three cases, leading to similar oxidation once exposed to ambient after the irradiation. Future experiments with varying dose levels are proposed to determine if trends between dose, dpa, and desorption and oxidation exist.

Figure 6.3 – Changes in C:Si ratio, Se:W ratio, SiOx:SiC ratio, and WOx:WSe2 ratio due to heavy metal ion exposure and storage in medium vacuum for two weeks.

6.2 Heavy metal ion damage to SiC substrates In addition to damaging the WSe2, the SiC substrate is also damaged by the heavy ions. These ions have a calculated penetration depth of about 1 µm for 2.5 MeV Fe and 4 MeV Ag, and 2 µm for 5 MeV Fe ions. These expected ranges mean that the ions are expected to penetrate through many of the exfoliated WSe2 flakes, which have an estimated average thickness of 1 µm based on optical profilometry (see Figure 6.4), as well as impact the uncovered areas of SiC. The calculated dpa for the selected heavy ions in SiC is 2.4 for the 5 MeV Fe ions, 3.8 for the 2.5

MeV Fe ions, and 8.5 for the 4 MeV Ag ions. Much as in the case of WSe2, the damage done to

SiC does not reflect these differences in dpa. The proposed cause of this is the same one posited for the exfoliated WSe2. Damage done to SiC is reflected in changes to both the silicon 2p and carbon 1s spectra in XPS.

(a) (b)

Figure 6.4 – Optical image (a) and height profile (b) of exfoliated WSe2 flakes on SiC

The major change in the silicon spectra (Figure 6.5) comes from the formation of a silicon- oxygen compound. Initially, only one chemical state is present for silicon – a nearly-degenerate doublet corresponding to the carbide. During the irradiation process, silicon-carbon bonds are broken and atoms of both species are ablated from the substrate, but the SiC nature of the substrate is still maintained. Upon exposure to the ambient, some silicon atoms react and form a silicon oxide with an expected silicon valence of 2+ or 3+ [132], indicative of a substoichiometric surface oxide. This oxide appears as a second nearly-degenerate doublet.

Figure 6.5 – Si 2p spectra before (black) and after (red) heavy metal ion irradiation

In contrast, no new peaks appear in the carbon spectra (Figure 6.6). From low binding energy to high binding energy, the carbon 1s spectrum for all samples is composed of silicon carbide, the sp3 carbon-carbon bonds in surface carbon, and the carbon-oxygen bonds in surface carbon. 65

All three states are remained after irradiation. However, the proportions are altered, with more carbon belonging to the two surface states rather than to SiC. Furthermore, the binding energy of the carbon in SiC increases from ~282.6 eV to ~283.4 eV. This increase in the binding energy is attributed to the full amorphization of the SiC, which has been previously observed in the literature [133] and is expected based on the dpa values of the ions used in this experiment [134].

Figure 6.6 – C 1s spectra before (black) and after (red) irradiation

A change in the SiC composition is also observed, as shown in Figure 6.3. Initially, the ratio of C to Si is measured as 0.84 ± 0.01. After irradiation, this ratio is now 0.47 ± 0.05, indicating that the lighter carbon atoms are preferentially sputtered from the SiC. Differences between samples are not reflective of the differences in dpa between the heavy ions, as discussed previously. The ratio after storage is 0.46 ± 0.02, indicating that the SiC retains its stability in air.

Similarly, the average ratio of Si-O to Si-C bonds is 0.79 ± 0.21 after irradiation and 0.80 ± 0.25 after two weeks in the vacuum chamber. While there are more pronounced differences between samples exposed to different ions, these differences are still not reflective of the differences in dpa. The most oxidation occurred on the sample exposed to the 2.5 MeV Fe ion, which has an intermediate dpa. The cause of this deviation from the expected trend is currently not known.

6.3 Heavy metal ion modification of WSe2 and SiC valence band spectra In addition to changes in the elemental spectra of WSe2 and SiC, changes in the valence band spectra are observed. The valence band spectra observed in XPS and shown in Figure 6.7 are a convolution of the valence bands of SiC and WSe2. The valence band maximum (VBM), as measured in the XPS software, is initially at 0.1918 eV, 0.0942 eV, and 0.5320 eV for the samples that would subsequently be bombarded by 5 MeV Fe ions, 2.5 MeV Fe ions, and 4 MeV

Ag ions. VBM values in this range are typical for this heterostructure, and result from the p-type doping and small band gap (~1.2 eV) of the exfoliated WSe2. After irradiation, the VBM of the three samples decreases to values of -0.1729 eV, -0.3424 eV, and -0.0353 eV, respectively. As all binding energies are with respect to the Fermi level, it is expected that heavy ion bombardment at a total fluence of 1016 /cm2 leads to a semiconductor-to-metal transition in the heterostructure due to excessive p-type doping. Such doping is attributed to the damage and oxidation of the WSe2. There is a positive trend between the extent of the binding energy shift and the dpa of the heavy ions. However, it is still the case that doubling the dpa of the ion does not double the extent of the doping. Determining the cause of this difference will require further investigation as suggested previously.

Figure 6.7 – Valence band spectra of the WSe2/SiC heterostructure before (black) and after (red) heavy metal ion irradiation

In addition to these shifts, changes in the shape of the valence band are also observed. The intensity peak observed around 9.5 eV is attributed to the crystalline SiC substrate, and

67 disappears after irradiation. The intensity drop here corresponds both to the loss of crystallinity and the ablation of SiC from the samples. Increases in intensity at binding energies near 8 eV and near 12 eV are attributed to the presence of tungsten oxide [135] and silicon oxide [136], respectively.

Chapter 7 – Conclusions and Future Work

In this work, the response of WSe2 to various forms of ionizing radiation has been explored through various investigations. MOCVD-grown WSe2 was found to be more stable against the combined stimuli of 1.486 keV X-rays and low energy electrons than the PV-grown MoS2 due to a greater coupling with the substrate. It is essentially insensitive to these stimuli when grown as a continuous thin film with no vertical features or discontinuities. The MOCVD-grown films are also found to be stable in air and medium vacuum over time. Low energy oxygen plasma can be used to convert the top layers of WSe2 into a mixture of WOx and SeOx. The MOCVD-grown material is more sensitive to oxidation than material exfoliated from a bulk crystal. This is likely due to a higher defect density in the as-grown material, as well as the reduced thickness.

16 The response of WSe2 to high energy proton exposure is complex at fluences of 10 protons/cm2 or more. There does not appear to be physical damage to the exfoliated flakes, but there is significant charging in the WSe2, both from direct ionization and from charge transfer into it from the SiC substrate. This surface charge leads to changes in the binding energy of the core shell electrons, but there are no macroscopic effects such as oxidation. The damage to the

SiC substrate causes it to change color from transparent to black, and a new absorption edge with an energy of ~1.2 eV was found, corresponding to the R center that is typically generated by high energy protons. The surface charge in the SiC is not as large as that of the WSe2, leading to apparently permanent changes in the band alignment that would be expected to be deleterious to a device using SiC as a gate dielectric. However, devices based on this heterostructure would also be more stable against radiation than WSe2-based devices on SiO2. As such, we can recommend SiC as a substrate for space-based WSe2 devices, but a suitable gate dielectric still needs to be found.

Heavy metal ion damage is more straightforward in comparison. Exposure to high energy

16 2 heavy metal ions such as Fe and Ag at a fluence of 10 ions/cm de-stabilizes the WSe2 via selenium volatilization. Upon air exposure, tungsten oxide is formed and can be detected in XPS.

Similarly, the ions readily break bonds in the SiC leading to increases in surface carbon, formation of silicon oxide, and amorphization of the remaining SiC. The samples also underwent a semiconducting to metallic transition due to the extensive damage. These changes are all stable over two weeks and appear to be permanent. As such, this exposure level is likely to destroy a space-based device based on the WSe2/SiC heterostructure.

Much further work is suggested based on the results obtained from these various studies. The stability of TMDs in the XPS environment as a function of the domain size and layer thickness is still to be determined. Studies of WSe2-based devices are suggested. For low energy oxygen plasma, complementary studies that more fundamentally investigate the role of chamber pressure and RF power on the oxidative potential of the plasma are encouraged, as well as further explorations of the differences between as-grown material and exfoliated samples. Determination of the proton sensitivity of a dielectric material known to be compatible as a top-gate to WSe2 and the sensitivity of the as-grown material are recommended to support device studies, and two such dielectric materials might be hBN [137] and HfO2 [138,139]. Further work with MoS2 is also encouraged to determine the actual impact of MeV protons on this 2DM. For heavy metal ion damage, determination of the damage threshold to WSe2, the sensitivity of the as-grown material, and the device response are yet to be known.

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