MULTI-STIMULI EFFECTS ON THIN FILMS AND DEVICES

The Pennsylvania State University

The Graduate School

MULTI-STIMULI EFFECTS ON THIN FILMS AND

DEVICES

A Dissertation in

Mechanical Engineering

by

Md Zahabul Islam

Submitted in Partial Fulfillment

of the Requirements

for the Degree of

Doctor of Philosophy

August 2020

The dissertation of Md Zahabul Islam was reviewed and approved by the following:

Aman Haque

Professor

Department of Mechanical Engineering

Dissertation Advisor & Chair of Committee

Adri van Duin

Kenneth Kuan-Yun Kuo Early Career Professor

Professor of Mechanical Engineering

Professor of Chemical Engineering

Professor of Engineering Science and Mechanics

Director of the Materials Computation Center

Daudi Waryoba

Assistant Professor, Engineering

Coordinator, Applied Materials Option

Penn State DuBois

Douglas E. Wolfe

Professor

Department of Materials Science and Engineering

Department of Engineering Science and Mechanics

Department Head, Advanced Coatings at the Applied Research Laboratory

Karen A. Thole

Professor

Department of Mechanical Engineering

Head of the Department of Mechanical Engineering

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Abstract

Physical properties of materials are known to depend on microstructure and defects across

multiple length-scales. The structure-property scaling becomes enhanced at the micro and nano

scales, an example of which is the ‘smaller is stronger’ phenomenon. Size effects also render

nanoscale materials more sensitive to external stimuli such as stress, temperature, electrical

current, light compared to their bulk counterparts. Even more interesting is the observation of the

breakdown of classical physical laws at length scales (grain size, thickness) at or below

characteristic length scales for physical domains. For example, scaling of yield stress (known as

the Hall-Petch law) breaks down below ~25 nm, where a grain cannot accommodate statistically

significant number of dislocations to induce plasticity. Similar breakdown phenomena have been

observed for other (electrical, thermal) domains. Fundamentals of the mechanics and physics of

nanomaterials is a prerequisite for the development of nanotechnology, which makes the length

scale and external stimuli effect on materials behavior as an attractive field of research.

While extensive efforts are ongoing to explore nanoscale structure-properties relationship in

single domains, this dissertation is rooted in processing-structure aspects of materials by

exploiting pronounced coupling exists among physical domains. Since the core of materials

processing relates to the response of the material to external stimuli (such as temperature), our

approach is to explore size or confinement effects that could make materials more sensitive to

external stimuli compared to conventional bulk materials. This lays down our hypothesis, ‘size-

induced coupling of multiple domains is manifested in form of unprecedented synergy of

multiple stimuli, which can be exploited to tailor microstructure or defect density to achieve

control over physical properties’. This hypothesis is aligned to the ulterior goal of this

dissertation, which is to develop novel materials processing techniques that are faster, more

iv

effective and energy efficient compared to the conventional (high temperature) thermal

annealing. Accordingly, the objective of this dissertation is to validate the hypothesis and

demonstrate it on two classes of materials spanning nano to micro scales namely, (i) ~2 nm thick

two-dimensional (2D) and (ii) 100 nm to 100 micron thick metals and additive manufactured

alloys. For each of these cases, we have explored the multi-stimuli synergy to achieve control

over crystallinity, grain size and defect density.

Since nanoscale characterization is challenging even in single domains, a key hurdle for this

research is to develop an experimental setup, which can simultaneously apply multiple stimuli,

or conversely, characterize the materials in multiple domains. We achieved this with a Micro

Electro-Mechanical System (MEMS) based framework, where strain, temperature and electrical

current are simultaneously applied on the specimen inside a high-resolution microscope that

visualizes the microstructural changes in real time. The setup is small enough to fit inside a

transmission electron microscope (TEM), which can provide atomic resolution. We have

compared the magnitude of these stimuli for both cases (i) when they are applied simultaneously

and (ii) separately to quantify the synergistic effect of the stimuli. In addition, we have also

quantified the time rate or the dynamics of microstructural evolution. We have also analyzed the

stimuli magnitude and microstructural dynamics to demonstrate the efficiency of our proposed

multi-stimuli materials processing technique. Our findings shows the enhanced atomic and defect

mobility due to the electrical current. It also reveals the microstructural transformation of near-

amorphous material to nanocrystalline materials. This type of transformation is difficult or

energy extensive for conventional thermal annealing. We have also investigated the pronounced

effect of multi-stimuli (instead of single stimuli) to observe the microstructural changes.

v

In this research, at first we have chosen e-beam deposited thin films and additive

manufactured (AM) alloys as platform materials. Additive manufacturing is a highly non-

equilibrium manufacturing process where laser sintering/melting results in defects spanning

micro to nanoscales. While the pores and voids can be eliminated by conventional thermal

annealing, more challenging tasks are the nanoscale defects, such as sub-grain structures. Thus,

we have explored the effectiveness of the multi-stimuli processing on microstructure control of

additive manufactured alloys as well as thin films (zirconium and palladium, gold). We also used

ion irradiation to generate controlled defects in polycrystalline gold films and then investigated

the effectiveness of multi-stimuli on defects annihilation.

Another platform material is 2D materials for their extremely small length-scale (mono to

few atomic layers configuration). Our study on chemical vapor deposited (CVD) MoS2, with

few atomic layers, inside a TEM shows the effectiveness of the stimulus effects on defects

annihilation and microstructural changes at low temperatures. Later on, we extend our multi-

stimuli synergy on 2D material based back-gated field effect transistor (FET). External stimulus

such as electrical current generates both resistive heating i.e., Joule heating and atomic scale

force also known as electron wind force (EWF) in a material. Study shows that this external

stimulus can induce significant amount of momentum on the defective sites even at low

temperature due to the EWF. Study also reveals that this unique EWF accompanied at low

temperature can enhance device performance in a short period of time span, which indicates this

proposed technique will potentially lead to time and cost-effective post-processing of two-

dimensional materials and their devices.

The scientific contribution of this research will be experimental validation of the hypothesis

that simultaneously applied stimuli are more effective, energy efficient and faster in achieving

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control over defects and microstructure compared to conventional thermal annealing process.

The potential impact of successful validation is a novel material processing technique, whose

unprecedented atomic and defect mobility at lower temperatures will open a new horizon in

defect engineering to modulate physical properties to find applications from nanotechnology to

advanced manufacturing.

vii

TABLE OF CONTENTS

List of Figures

……………………………………………………………………………………

x

Acknowledgements

………………………………………………………………………………

………………………………………………………………………………….

xvii

Chapter 1 Introduction

……………………………………………………………………………

1

1.1 Central Theme of this Dissertation

1

1.2 Structure-Property Relationship in Materials

…………………………

2

1.3 Conventional Microstructural Control Approaches

5

1.4 Proposed Multi-stimuli Approach 7

1.5 Objectives and Impacts of this Research

……………………………………………..

12

Chapter 2 External Stimuli Sensitivity in Thin Films and Additive Manufactured

Alloys …..

19

2.1 Temperature-Electron Wind Force Synergy in Thin Films

……………………..

20

2.1.1 Objective and Motivation ………………………………………..

20

2.1.2 Materials and Methods ………………………………………….. 22

2.1.3 Results and Discussion ………………………………………….. 25

2.1.4 Conclusion ………………………………………………………. 29

2.2 Multi-Stimuli (Electron Wind Force and Mechanical Strain) Effects in

Thin Films ….………………………………………………………………

30


31



2.2.4 Conclusion ……………………………………………………… 45

2.3 Low Temperature Processing of Additive Manufactured Ti64 Alloy 46

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47



2.3.4 Conclusion ……………………………………………………… 56

Chapter 3 Stimuli Effects on Two-Dimensional (2D) Materials and Devices

……………..

58

3.1 Low Temperature-Electron Wind Force Synergy in Molybdenum Disulfide

…...

59


59



3.1.4 Conclusion ………………………………………………………. 70

3.2 Low Temperature Processing of 2D Material based Thin Film Transistors

…...

71


71



3.2.4 Conclusion ………………………………………………………. 83

Chapter 4 Synergy of Stimuli On Operation and Degradation of Nanoscale Devices

……..

85

4.1 On-state Degradatation of High Electron Mobility Transistor

……………….

86


86



4.1.4 Conclusion ………………………………………………………. 97

4.2 Off-state Failure of High Electron Mobility Transistor

………………………

97

ix


98



4.2.4 Conclusion ………………………………………………………. 107

Chapter 5 Ion Irradiation and External Stimuli Effect at Nanoscale …………………… 108

5.1 Irradiation Damage and Degradation in Nanoscale Transistor

………………….

109


109



5.1.4 Conclusion …………………………………………………….. 124

5.2 Low Temperature Synergy on Recovery of Irradiation Damage in Thin

Films…………………………………………………………………….......

124


125



5.2.4 Conclusion ………………………………………………………. 133

6. References ……………………………………………………………………………. 134

x

List of Figures

Figure 1.1. Passage of electrical current creates thermo-electro-mechanical effects that

can be tuned for microstructural control …………………………………………………

......................

1

Figure 1.2. Electrical Annealing showing defects annihilation …………………………. 2

Figure: 1.3 (a) Defect confinement leading to size effect in materials, and (b) The Hall-

Petch effect in materials …………………………………………………………………

3

Figure 1.4. (a-c) Defects spanning multiple length-scales in metallic materials, Defects

in 2D materials: (d) Vacancy, (e) dislocation, and (f) grain boundary structures in high-

resolution transmission electron microscope (TEM)……...………………………………

5

Figure 1.5. (a) Atomistic representation is showing uniform heating of materials during

conventional thermal annealing, and (b) Time and energy intensive conventional

thermal annealing exhibits grain growth of material ……………………………………

7

Figure 1.6. (a) Effect of electrical current on materials, (b) & (c) an alternative route

shows the localized heating at the vicinity of the grain boundaries (GBs) ……………..

8

Figure 1.7. Synergy of temperature, current and stress fields can create high atomic flux

and mobility along the grain boundaries, which are diffusion pathways in a metal. The

phenomenon is pronounced in areas with higher fraction of defects or disorder (such as

grain boundaries) ………………………………………………………………………..

9

Figure 1.8. (a) Schematic showing active temperature control using liquid N2, (b)

Gripping on massive heat sinks introduces two distinct temperature zone in the sample ..

11

Figure 1.9. MEMS device mounted on in-situ TEM holder, and (b) MEMS device with

heater, sensor, actuators and biasing capability ………………………………………….

13

xi

Figure 1.10. Mechanical straining: (a) Fabricated MEMS device using standard

nanofabrication techniques, (b) schematic showing sensors and samples integrated with

MEMS device, and (c) spring-equivalence of specimen-device system …………………

14

Figure 1.11 Flowchart showing EWF implemented in MD simulation using LAMMPS

package …………………………………………………………………………………...

15

Figure 2.1. (a) Scanning electron microscope (SEM) micrograph of the MEMS device

showing the current flow through the specimen. Inset shows diffraction pattern at 0

A/cm2 current density, (b) Atomistic model with grains oriented at different angles, and

(c) Electro-thermal simulation of sample with actual geometry, resistance and current

density …………………………………………………………………………………… 24

Figure 2.2 In-situ TEM study indicating grain growth as a function of dc electrical

current density ……………………………………………………………………………

26

Figure 2.3. Comparison between thermal and electrical thermal annealing: (a) TEM

bright field image after thermal loading (b) corresponding SAED pattern, (c) MD

simulation cell showing grain boundary reconstruction in limited locations indicated by

arrows, (d) TEM BF image after current loading, (e) corresponding SAED pattern, and

(f) MD simulation showing grain boundary reconstruction due to the electrical current

loading ……………………………………………………………………………………. 28

Figure 2.4. Time evolution of grain growth obtained from MD simulation trajectory: (a)

initial structure, (b) two triple points before electrical annealing, and (c) two triple points

after electrical annealing ………………………………………………………………… 29

Figure 2.5. In-situ TEM experimental and MD simulation setups: (a) Scanning electron

micrograph of the adopted MEMS device with actuators and electrodes including a

TEM bright field (BF) image and diffraction pattern of as-deposited specimen, (b)

temperature profile along the length of the sample obtained from electro-thermal

simulation mimicking the actual experimental conditions, and (c) atomistic simulation

cell with randomly oriented grains used to mimic the experiments …………………….. 34

Figure 2.6. (a) TEM bright field (BF) image of the as-deposited specimen after

prolonged exposure to the electron beam, and (b) corresponding SAED pattern ………. 37

xii

Figure 2.7. In-situ TEM BF and SAED evidence of grain growth in the specimen center

region at (a, d) 0 x105 A/cm2, (b, e) 6.5x105 A/cm2, and (c, f) 7x105 A/cm2 current

densities ………………………………………………………………………………….

39

Figure 2.8. Specimen microstructure at 7x105 A/cm2 current density and room

temperature, (a) before, and (b) after application of 0.1% strain ……………………….. 40

Figure 2.9. MD simulation model on grain growth: (a) initial structure, (b) electrical

annealing prior to the application of strain, and (c) electrical annealing after strain

application, (Green color indicates face centered cubic (FCC) and red color indicates

hexagonal closed packed (HCP) phase of palladium) ……………………………………

43

Figure 2.10. (a) Strain energy increment during tensile straining of the system, and (b)

grain size as a function of applied tensile strain …………………………………………

44

Figure 2.11. Schematic showing experimental set-up with temperature controlled stage.. 49

Figure 2.12. Optical micrographs of Ti64 specimens in the (a) as-built, and (b) Low

temperature EWF processed conditions ………………………………………………….

52

Figure 2.13. Basal Schmid factor maps for the Ti64 specimens in the (a) as-built, and (b)

electric current processed sample, and (c) Calculated Taylor factors Twining during

electrical annealing: (d) as-built specimen and (e) after applying a current density of

5x103 A/cm2……………………………………………………………………………….

54

Figure 2.14. KAM maps for the Ti64 specimens in the (a) as-built and (b) electric

current processed conditions. A threshold of 5° was used to exclude well-defined grain

boundaries in the analysis ………………………………………………………………..

55

Figure 2.15. (a) Force-displacement plot obtained from nanoindentation experiment, (b)

calculated hardness and Young’s modulus ………………………………………………. 56

Figure 3.1. Schematic shows the transfer process of monolayer MoS2 on to a MEMS

device and subsequent experimental setup for in-situ TEM investigation ………………

63

Figure 3.2. (a) Observation of almost one order of magnitude reduction in electrical

resistance of MoS2 specimens during EWF annealing, and (b) spatial temperature

distribution for the highest current density ………………………………………………

64

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Figure 3.3. In-situ TEM electron wind force annealing results: (a) Bright-field image,

(b) SAED pattern of as-deposited nominally monolayer MoS2 specimen, and (c, d) The

same location after annealing at 9.5x105 A/cm2 current density ……………………….

66

Figure 3.4. (a)-(c) 6|8Mo type dislocation migration during the electrical annealing, (d)-

(e) Transformation of a 6|8 ring to 6|6|4 ring at the Grain boundary (GB) (Individual

grains are shown by different colors, smaller radius sphere indicates Sulphur atoms and

larger sphere indicates Molybdenum atoms) …………………………………………….

68

Figure 3.5. Transformation of vacancy defects at the GB: (a) initial sample, (b) vacancy

transforms to 6|8 S defects, (c) 6|8 S transforms to 6|6|4 ring defects, (d) formation of

6|8 S due to the dislocation motion, and (e) formation of 6|6|4 ring (Individual grains are

shown by different colors, smaller radius sphere indicates sulfur atoms and larger sphere

indicates molybdenum atoms) ……………………………………………………………

69

Figure 3.6. (a) Optical Microscope image of fabricated back-gated WSe2 FET transistor,

(b) schematic diagram of WSe2 transistor, (c) electro-thermal simulation model of WSe2

FET, and (d) temperature profile across the cross-section of the sample obtained from

model ……………………………………………………………………………………..

76

Figure 3.7. (a) Output characteristics of WSe2 FET after annealing at different drain

voltage, and (b) Improvement in drain current after annealing while FET surface was

maintained at 296K ……………………………………………………………………….

78

Figure 3.8. (a) Transfer characteristics showing WSe2 transistors performance after

annealing at different voltage, and (b) maximum drain current obtained after annealing ..

80

Figure 3.9. Monolayer WSe2: (a) prior to the annealing, and (b) after annealing ……… 81

Figure 3.10. (a)-(f) Failure at high biasing condition due to the electrical and thermal

field, and (g)-(j) side view of the sample showing void creation at the cathode side (left)

and mass accumulation at the anode side (right) (Color bar in Figure 3.12 shows the

atomic stress distribution in the sample) …………………………………………………

82

Figure 4.1. Comparison of properties among competing semiconductors ……………..... 87

Figure. 4.2. (a) Optical micrograph of GaN HEMT die, and (b) low magnification SEM

image of the die …………………………………………………………………………... 89

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Figure 4.3. Details of the GaN HEMT specimen preparation and transfer technique

using FIB for in-situ TEM reliability study …………………………………………........ 90

Figure 4.4. In-situ TEM reliability study of GaN HEMT: (a) electron transparent

specimen prior to the application of voltage stimulus, (b) device after failure, and (c)

output characteristic curve obtained from “on-state mode operation” …………………… 92

Figure 4.5. (a) Degradation of the passivation layer, (b) Evaporation of the buffer layer

due to the high thermal field, and (c) SAED indicating the transformation of GaN from

crystalline to amorphous state ……………………………………………………………. 94

Figure 4.6. GaN HEMT device degradation: (a) Hot electron induced failure at the

source side, (b) evaporation of the buffer layer and formation of small crystallites

(nanoparticles) due to the high thermal field, and (c) high-resolution image of a

spherical crystallite ………………………………………………………………………. 95

Figure 4.7. In-situ TEM reliability testing showing real-time device degradation ………. 96

Figure 4.8. (a) The experimental setup showing a MEMS chip on a TEM specimen

holder, (b) The specimen integrated with the MEMS chip, (c) SEM image of the

electron transparent GaN HEMT specimen, and (d) transfer characteristic of the HEMT

die, and (e) Off-state loading of the 100 nm thick HEMT sample ………………………. 100

Figure 4.9. Off-state characterization of (a) die-level transistor, and (b) 100 nm thick

HEMT during phase I loading …………………………………………………………… 101

Figure 4.10. Bright field TEM images acquired at drain voltages: (a) 0V, (b) 7.2V (c)

11.6V and (d) 23V, and (e) corresponding drain current vs. drain voltage data at gate

voltage -5V for a 100 nm thick GaN HEMT specimen …………………………………. 103

Figure 4.11. Real-time operation of GaN HEMT: (a) electron transparent HEMT

specimen prior to the loading, where arrows indicating pre-existing defects, (b)

magnified view of the arrows marked 2 and 3 at the on-set of source-drain leakage, (c)

molten drain metal pool at this instant the drain side at current density of 2000 mA/mm,

(d) metal diffusion through the GaN layer, (e) rapid breaching of the GaN-SiC interface,

105

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and (f) degradation of the SiC layer ………………………………………………………

Figure 4.12. EDS mapping of a failed HEMT specimen at the (a-d) gate and (e-h) drain

areas. (i, j) normalized weight percentage of the various elements in the gate and drain

area respectively ………………………………………………………………………….. 106

Figure 5.1. Experimental setup for in-situ TEM experiment of electron transparent

HEMTs: (a) GaN HEMT die, (b) a MEMS chip with the HEMT specimen mounted on

in-situ TEM electrical biasing holder, and (c) FIB lamella of the HEMT before

mounting on to the MEMS chip …………………………………………………………. 113

Figure 5.2. (a) Downward arrowhead in the schematic diagram of the GaN HEMT

showing irradiation direction, (b) displacement per atom (dpa) profile as a function of

depth for different doses of irradiation, (c) TEM image of a pristine HEMT showing

mostly bend contours, and (d) TEM image of an irradiated HEMT at 45 dpa showing

very high dislocation density …………………………………………………………….. 115

Figure 5.3. Die-level HEMTs specimens characterization curve as function of ion

irradiation damage in dpa: (a) transfer characteristics, and (b) output characteristics ....... 116

Figure 5.4. Electron transparent HEMT device: (a) before, and (b) after 2.8 MeV Au4+

ion irradiation for 60 minutes to a fluence of at 4x1014 ions/cm2. The rectangular dashed

box shows contrast change due to point defect accumulation, while the arrows indicate

dislocation activities at the GaN-SiC interface ………………………………………….. 117

Figure 5.5. (a) Drain current vs. drain voltage plot of electron transparent GaN HEMT

specimen during off-state operation inside the TEM. The data labels correspond to the

in-situ TEM images in Figure 5.6., and (b) Comparison between pristine and irradiated

conditions ………………………………………………………………………………… 119

Figure 5.6. TEM BF images showing source, gate and drain at the same time. Drain

voltage: (a) Vd= 0V, (b) TEM image of screw dislocations in buffer layer, (c) Vd= 5V,

(d) Vd= 7V, (e) Vd= 8.5V, and (f) Vd= 10.2V ……………………………………………. 120

Figure 5.7. (a) TEM BF image at the drain side of drain-gate region, (b) Dislocations in

the GaN layer, (c) High-resolution TEM (HRTEM) image of dislocations in GaN layer,

122

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(d, e) atomic strain mapping in the sample showing normal and shear strain field

associated with individual dislocations, and (f) simulated lattice fringes with the

dislocations ……………………………………………………………………………….

Figure 5.8. EDS mapping of GaN HEMT showing diffusion of chemical elements at

different drain bias: (a) Vd=0V, (b) onset of failure at Vd=10.2V, and (c) relative

changes in diffusion of chemical elements obtained from EDS at these two voltages ….. 123

Figure 5.9. (a) Displacement per atom (dpa) profile for irradiation dose of 6.5×1015

ion/cm2, (b) micro-electro-mechanical system (MEMS) device mounted on in-situ TEM

holder, (c) temperature profile obtained from electro-thermal simulation, (d) TEM BF

image showing irradiation damage, (e) BF image showing dislocation annihilation at

9.5×105 A/cm2, and (f) SAED pattern after EWF annealing …………………………….. 128

Figure 5.10. (a) Irradiated sample before processing, (b) dark field image of pre-

processed sample showing dislocation lines, (c) dislocation lines interaction during

EWF processing at 3.5×105 A/cm2, (d) partial annihilation (pink circle) of dislocation

lines, and migration towards grain boundary at 7×105 A/cm2, (e) migration of

dislocation lines towards GB and partial annihilation of dislocation lines at 9×105

A/cm2, and (f) complete annihilation of dislocations and defects at 9.8×105 A/cm2 …….. 130

Figure 5.11. Computational results on (a)-(c): Annihilation of dislocations and vacancy

clusters, and (d)-(f): SFT annihilation under EWF ………………………………………. 131

Figure 5.12. HRTEM images of low temperature processing of irradiated materials: (a)

defects in the irradiated sample before processing, (b) HRTEM image of SFT, (c)

surface induced annihilation of SFT at a current density of 7×105 A/cm2, (d) SFT-GB

interaction at 9×105 A/cm2, (e) after EWF annealing at 9.5×105 A/cm2, and (f) HRTEM

image after annealing at 9.5×105 A/cm2 …………………………………………………. 132

xvii

Acknowledgement

I would like to express my sincere gratitude to my advisor Dr. Aman Haque for his continuous

support, patience, and encouragement during my Ph.D. research. His resourceful guidance and

positive criticism helped me to accomplish my research goal. I could not have imagined having a

better advisor and mentor for my Ph.D. study.

Besides my advisor, I would like to thank the rest of my doctoral committee: Dr. Adri van Duin,

Dr. Daudi Waryoba and Dr. Douglas E. Wolfe for their time, and willingness to help.

My sincere thanks goes to Dr. Stephen Pearton and Dr. Fan Ren at University of Florida, Dr.

Khalid Hattar at Sandia National laboratories, Dr. Nicholas Glavin at Air Force Research

Laboratory, Dr. Joshua Robinson, and Dr. Saptarshi Das at Penn State, Dr. Huajian Gao at

Brown University, and Dr. Tim Rupert at University of California Irvine who provided me the

opportunity to work with them. Special thank goes to Dr. Daudi Waryoba for sharing his

profound knowledge with me on electron back-scattered diffraction (EBSD) technique.

I thank my former colleagues Dr. Baoming Wang, Dr. Raghu Pulavarthy, Dr. Tun Wang,

Niranjana Sunderasan, James Kidd and Angela Paoletta for sharing their invaluable experience

with me. I am grateful to my friends for their support and encouragement. I would also like to

thank MRI staff at Penn State Nanofab: Guy Lavallee, Shane Miller, Kathy Gehoski, Michael

Labella, Andy Fitzgerald, Chad Eichfeld, Ted Gehoski, Bill Drawl, Bill Mahoney, Bangzhi Liu,

Beth Jones; Penn State Material Characterization Lab: Trevor Clark, Jennifer Gray, Tim Tighe,

Wes Auker, Max Wetherington, Nichole Wonderling, Maria Dicola, Julie Anderson, Ke Wang,

Haiying Wang, Mike Norrell; Penn State Electrical Characterization Lab: Jeffrey Long and Steve

xviii

Perini for helping me with tools’ training regarding the nanofabrication process, microscopy,

and characterization techniques.

I gratefully acknowledge the support from the Division of Civil, Mechanical, & Manufacturing

Innovation (Nanomanufacturing program) of the National Science Foundation through award #

1760931 and US National Science Foundation (DMR 1609060). The views expressed in the

dissertation do not necessarily represent the views of the US National Science Foundation or the

United States Government.

Last but not the least, I would like to thank my parents for their unconditional love and supports

throughout my life. I would also like to thank the rest of my family members for supporting me

throughout my Ph.D. journey and my life in general.

1

Chapter 1 Introduction

1.1. Central Theme of this Dissertation

Mechanical properties such as yield strength and fracture toughness are governed by

microstructure (grain size, crystallinity, phases and precipitates, defects) at all length-scales [1-

6]. Similarly, electrical and thermal properties are also microstructure dependent. This is evident

from a vast literature in mechanics and physics of materials dedicated to quantitative and

qualitative mapping of structure-property relationship [1, 7-12]. Therefore, significant efforts

have been made to develop materials synthesis and processing techniques with the core objective

of achieving control over the materials properties and functionality. The state of the art

microstructural processing techniques predominantly use temperature as the stimulant while

often experimenting with alloying elements, defects and interfaces (such as grain boundaries) to

control microstructure. Even after extensive research, microstructural control has remained

elusive, which is the major motivation for this research.

Figure 1.1. Passage of electrical current creates thermo-electro-mechanical effects that can be

tuned for microstructural control.

2

This dissertation explores a new route where synergy of thermal, electrical and

mechanical stimuli is exploited for faster, energy efficient materials processing. Figure 1.1 shows

experimental evidence of electrical current driven microstructural/defect control in additive

manufactured 316 stainless steel. Passage of electrical current in metallic materials compound

thermal, mechanical and electrical effects that can contribute to the atomic or defect mobility. A

unique aspect of this dissertation is to consider athermal effects by eliminating the thermal field

generated from the Joule heating. This is shown in the shaded region in Figure 1.2. Another

important aspect is the role of any external mechanical (stress) or thermal (temperature)

stimulant superimposed on the electrical current density effects. Localization effects of these

stimuli is proposed to be the predominant factor behind the synergy. The resulting mechanics

and physics involve an intricate relationship between solid mechanics, heat transfer, electro-

migration and materials science, which we aimed to explore in this dissertation.

Figure 1.2. Proposed multi-stimuli synergy for materials processing.

1.2. Structure-Property Relationship in Materials

Figure 1.3 illustrates how microstructure governs mechanical properties at all length-

scales. Here, we define physical size as the thickness or width of a specimen material.

3

Microstructural size, on the other hand, can be grain size (D), diameter of dislocation loop (d),

obstacle spacing (L), spacing between partial dislocations (w) among others [2]. While the

strength does not differ for a steel rod of two different cross-sections, it can significant vary if the

grain sizes are different even for the same cross section. For metals and alloys, the classical Hall-

Petch formulation is shown in Figure 1.3b, which demonstrates the influence of microstructure

over yield strength. It is essentially a scaling law for strength, which is applicable for grain size

as small as 100 nm. Below this length-scale, the Hall-Petch scaling law becomes less reliable. At

the same time, when the physical size becomes comparable with this length-scale, remarkable

compounding effect on properties is observed. For even smaller size, either physical or

microstructural, the classical scaling law breaks down and can reverse. The inverse Hall-Petch

relationship has been observed experimentally which indicates yield strength decreases or

remains constant at or below 100 nm grain size [3]. Thus, we define ‘characteristic’ length-scale

as where classical scaling law breaks down or deviates appreciably.

Figure: 1.3 (a) Defect confinement leading to size effect in materials [2]. (b) The Hall-Petch

effect in materials [4].

4

Microstructural defects in materials spanning from atomic vacancy to three-dimensional

pores/voids pervade over the length-scales (Figure 1.3) and significantly affect physical

properties [1-3, 6, 8, 10, 12-14]. For example, grain size in metals is one of the most influential

parameter that controls physical properties of nanocrystalline materials [5]. Nanocrystalline

materials exhibit pronounced grain size effects (10-100nm) as manifested through the breakdown

of the classical Hall-Petch relation [6] (shown in Figure 1.3b). Likewise, in a two-dimensional

(2D) material, if physical size such as grain size (1-20nm) overlaps characteristics length scale

i.e., mean free path in thermal or electrical domain, then a remarkable effects are observed on

transport properties [15-19]. Hence, electrical, thermal and optical properties are influenced by

grain size. Electrical conductivity decreases with grain size [7], which is prominent around

length-scales (film thickness, surface roughness and/or grain size) comparable to the electron

mean free path [20]. With decreasing interconnect size, electrical conductivity is critical for

electronic devices. In metals, heat is carried by electrons and a similar grain size dependence of

thermal conductivity is observed [8]. Similarly, phonon scattering is also grain size dependent

[21]. In addition to the grain size, other microstructural defects such as interstitials, vacancies,

dislocations and stacking faults [22, 23] also affect physical properties such as mechanical [13,

24-26] , electrical [14, 27, 28], thermal [29-31], and optical properties [32-34], to name a few.

Figure 1.4 represents commonly observed different types of defects in metals and two-

dimensional (2D) materials. Grain size control, elimination of defects and residual stress is

critical from both performance and reliability perspective for real-world applications [35].

Therefore, high temperature annealing is a common post-processing step in fabrication.

However, high temperature annealing does not always result in perfect microstructure.

Depending upon the systems, it can create thermal stress to stagnate the defect diffusion or even

5

may create new defects. In the following sections, we will discuss how we can gain control over

microstructure and defects. To answer this question, we will investigate the available material

processing techniques and their advantages as well as limitations, and how we can overcome

these limitations.

Figure 1.4. (a-c) Defects spanning multiple length-scales in metallic materials [36-38], Defects in

2D materials: (d) Vacancy [39], (e) dislocation [40], and (f) grain boundary structures [41] in

high-resolution transmission electron microscope (TEM).

1.3 Conventional Microstructural Control Approaches

Materials science has evolved around ‘processing-microstructure-property’ relationship,

where the processing or synthesis component is particularly highlighted. Materials processing,

6

for example annealing, requires mobility of defects and atoms, which has been conventionally

imparted by high temperature. Thermal processes such as conventional thermal annealing induce

random diffusion, which is a very slow mechanism [42] as shown in Figure 1.5b. Also, they heat

both defective and crystalline regions uniformly. This is schematically shown in Figure 1.5a. In a

conventional annealing thermal stimulus is used for controlling grain size, phases and defect

density. Conventional thermal annealing is performed by raising the temperature uniformly to a

very high value (homologous temperature exceeding 0.5) at certain rate, and then hold it and

eventually cool it down to room temperature at rates depending on the material [43-45]. The role

of temperature here is to increase atomic and defect mobility to facilitate microstructural

reorganization. However, during the uniform temperature raise of the material most of the energy

we spend to heat up the existing crystalline area (grain interior), hence we consider the uniform

temperature raise of the whole grain as an energy inefficient technique. For example, Figure 1.5b

shows how temperature and time influence the strength and grain size respectively. In addition,

during the thermal annealing of materials and devices with complex interfaces, requirement of

high temperature itself can degrade the materials by creating more defects or stagnating the

existing ones due to the residual thermal stress between materials’ interfaces. Thus, this high

temperature processing might not work for flexible electronics or temperature sensitive devices.

These conventional processing limitations inspire us to investigate an alternative pathway for

active controlling of microstructure and defects. We therefore conclude that the current art of

single (thermal) stimuli microstructural control is not only energy intensive, but also time

consuming, and could be detrimental for flexible or temperature sensitive devices. Thus, in the

following sections we will probe in detail on our newly developed low temperature and time

7

efficient technique, which exploits multi-stimuli synergy to process and control microstructure in

an active manner.

Figure 1.5. (a) Atomistic representation is showing uniform heating of materials during

conventional thermal annealing, and (b) Time and energy intensive conventional thermal

annealing exhibits grain growth of material [42].

1.4 Proposed Multi-Stimuli Approach

While thermal processing has been around since the Bronze Age, this dissertation

considers a new route, where synergy of electrical current, mechanical strain and temperature is

focal point of research. A potential outcome could be non-thermal enhancement of defect or

atomic mobility that could help us to achieve microstructural control. We hypothesize that at

micro to nanoscales (relevant to microstructure but not necessarily the physical dimension) these

stimuli are synergistic, i.e., their effects are not merely additive but compounding. A corollary of

our hypothesis is that significantly higher atomic and defect mobility can be achieved even at

lower temperatures due to the electrical current passage, and the effects will be more pronounced

8

if it is accompanied by mechanical strain. Thus, in our present study, we developed a non-

thermal route to conventional thermal annealing. Figure 1.6 illustrates approach where we pass

electrical current through metallic or semiconducting materials. The literature identifies this

process as ‘electrical current annealing’ (or more commonly electro-pulsing [46-50]), which

induces Joule heating (Figure 1.6a) to raise the temperature. It also develops ‘electron wind

force’ (EWF), which is nothing but the momentum transfer from electrons to defects. Thus, EWF

is inherently a mechanical force that is highly localized around defects where the momentum

transfer (or electron-defect scattering) takes place. This is analogous to a gentle wind blowing

the leaves (defects) in a tree while the limbs (lattice) remain unperturbed. A major contribution

of this dissertation is the investigation of EWF alone on the acceleration of defect mobility. In a

stark contrast to electro-pulsing, we achieve this by actively removing the Joule heating. This not

only keeps the material at near-room or low temperatures, but also triggers mobility ‘just in

location’ or where needed. In other words, our proposed isolation and application of EWF

specifically targets defective atoms, leaving the crystalline lattice alone.

Figure 1.6. (a) Effect of electrical current on materials, (b) & (c) an alternative route shows the

localized heating at the vicinity of the grain boundaries (GBs).

9

To investigate our hypothesis, we start with very small current density (30-50% of the

electromigration failure limit). Electrical current is a very effective stimulus to generate [51, 52]

and flow vacancies. This is due to the electron wind force arising from the momentum transfer

between conduction electrons and metal ions. The most important aspect is that the current

density impacts the defects and grain boundaries significantly more than the crystalline grain

interior. Inside the grains, the uniform lattice structure means there is less momentum transfer.

The opposite is true at the grain boundaries, where most of the vacancy generation and motion

takes place [53]. Figure 1.7a shows this phenomenon, where blue arrows indicates electron flow

and red arrows shows the impact of electron on grain boundary. Thus, electrical annealing is

highly localized to defects around the grain boundary, where we need the atomic mobility and

not the entire grain area. This enhances the energy efficiency.

Figure 1.7. Synergy of temperature, current and stress fields can create high atomic flux and

mobility along the grain boundaries, which are diffusion pathways in a metal. The phenomenon

is pronounced in areas with higher fraction of defects or disorder (such as grain boundaries) [54].

10

Mechanical stress field also has a synergistic effect on the electrical annealing process

[55]. The vacancy flux due to electrical annealing decreases as soon as internal stresses build up

in the metal due to vacancy transport from cathode end to anode end. The role of an external

mechanical field is to counter the effects of the internal stress and the net effect will be higher

vacancy flux and mobility [56]. The most important aspect of stress field is that stresses are

highly localized at the grain boundaries (Figure 1.7c). This means only a small amount of stress

is required as an external stimulus, because it becomes amplified at the defects and grain

boundaries. This effect is very similar to current density, and therefore it also specifically targets

defects and not the material uniformly, which makes it both time and energy efficient. To

summarize, we can express individual contribution from each stimulus such as EWF (JEWF),

temperature (JT) and mechanical stress (Jσ) on the atomic flux mobility as follows [57]:

𝐽𝐸𝑊𝐹 =𝑁𝑒𝑍∗𝐷𝜌

𝑘𝑇𝑗; 𝐽𝑇 = −

𝑁𝐷𝑄𝜌

𝑘𝑇2 ∇𝑇; 𝐽𝜎 =𝑁𝐷𝛺

𝑘𝑇∇𝜎 (1.1)

where, N is vacancy concentration, k is Boltzmann’s constant, T is absolute temperature, e is the

elementary charge, Z* is effective charge number, D is diffusivity, ρ is resistivity, j is current

density, Q is the heat of transport, Ω is the atomic volume, and σ is the hydrostatic stress.

Since our goal is to develop non-thermal low temperature microstructural processing

using multi-stimuli, a distinct feature of our study is that we remove the heat generated by

current to separate the effects of temperature. Thus, to ensure low temperature processing we

took two distinct approaches. In the first approach, we actively control the sample temperature

by passing liquid N2 through the sample stage as shown by the schematic diagram in Figure 1.8a.

In the second approach, we gripped freestanding thin film specimen with massive heatsinks. For

a freestanding specimen, the current flow heats up the mid-point region the most, while the two

https://www.sciencedirect.com/topics/engineering/atomic-volume

https://www.sciencedirect.com/topics/engineering/hydrostatic-stress

11

end regions (anode and cathode) are constrained to lower temperature as shown in Figure 1.8b.

This is because the end regions are more massive and act as heat sinks. Suitable design of the

specimen and the two ends makes it possible to have the ends at ambient condition. The

uniqueness of this setup is that we are able to decouple the effect of Joule heating and the

electron wind force in the same specimen. The mid-point of the specimen shows the effect of

EWF and temperature, while the two end regions show the effects of EWF only.

Figure 1.8. (a) Schematic showing active temperature control using liquid N2, (b) Gripping on

massive heat sinks introduces two distinct temperature zone in the sample.

To validate the low temperature processing, we pass electrical current through the

specimen. The electrons are minimally scattered by the lattice, but they transfer their momentum

wherever they meet defects and grain boundaries [58]. This gives rise to highly defect-specific

atomic scale force EWF. By combining this EWF with other stimuli such as mechanical strain,

we can achieve active microstructural control across multiple length scale, namely bulk-, micro-,

and nano-scale. The uniqueness of this technique is the non-thermal processing of microstructure

and defects using multiple stimuli i.e., mechanical strain and EWF, whereas the latter one is

intense around the defects and non-existent in defect-free lattices.

12

1.5 Objectives and Impacts of this Research

The objective of the present research is to validate the hypothesis that at micro to nanoscales

multi-stimuli synergy can be exploited to achieve time and energy efficient microstructural

processing and control at low temperature to gain control over microstructure and defects. Thus,

impact of this hypothesis is the potential for tailoring defects and microstructures for desired

performance. Our developed multi-stimuli materials processing technique can potentially

outperform the existing approaches in both speed and energy requirement, thereby profoundly

impacting the next generation materials processing technology. The impact of this research can

be realized for the design and development of materials spanning from nano to microscale. For

example, these unique multi-stimuli synergy can be exploited to enhance performance of next

generation advanced manufacturing including but not limited to nanoelectronics, sensors as well

as additive manufactured alloys. Our scalability study of multi-stimuli synergy on additive

manufactured metallic alloys shows potential promise for active microstructural control, which

will be realized by the large potential of metallic applications in civil and military infrastructure,

automotive, aerospace, heavy machinery and cutting tools, oil and gas, protective coatings, bio-

medical implants and devices, to name a few [59-61]. Additionally, discovery of graphene[62]

has brought the concept of single layer atomic device and components such as back-gated field

effect transistor (FET) closer to the real life applications. Thus, 2D materials such as graphene,

h-BN, MoS2, WSe2, phosphorene, silicene etc., have gained significant interest due to their

outstanding electrical, optical, chemical, and thermal properties [63-70]. Potential applications of

2D materials encompasses electronic devices[71], sensors[72], catalysts[73], energy conversion,

storage devices[74], to name a few. Our onboard low temperature rapid processing might be an

13

alternative viable path to enhance device performance without introducing residual thermal

stress.

To accomplish this objective, we have developed a versatile micro-electromechanical

system (MEMS) based experimental setup with sensors, heater and actuators integrated with

micro-, nano-scale specimens. The uniqueness of this set-up allows measuring multi-domain

properties under controlled exposure to multiple stimuli as well as measuring the resulting

properties at high resolution. These fabricated MEMS devices (as shown in Figure 1.9) have

been also tested with electron, infrared, thermoreflectance, and Raman microscopes to access the

various stimuli and or domains. Figure 1.9b shows the SEM micrograph of the fabricated MEMS

device with sensors, actuator and heaters. In order to accomplish electrical annealing, we

supplied electrical current through the electrode pad B1 and C. Mechanical strain can be applied

on the sample by passing current through the actuator pad A1 and A2. Combination of these

stimulus (e.g multi-stimuli effect) can be achived at the same time by powering up the elctrode

pads simultaneously. Figure 1.10 shows our devices’ capability to study multi-stimuli synergy on

microstructural control.

Figure 1.9. MEMS device mounted on in-situ TEM holder, and (b) MEMS device with heater,

sensor, actuators and biasing capability.

14

Figure 1.10. Mechanical straining: (a) Fabricated MEMS device using standard nanofabrication

techniques, (b) schematic showing sensors and samples integrated with MEMS device, and (c)

spring-equivalence of specimen-device system.

To understand the multi-stimuli synergy on microstructure control, we have also developed

molecular dynamics simulation code using LAMMPS [75] package. We started with single

stimulus effect (e.g either mechanical strain, thermal or electrical current) and later on, we have

extended our study to observe the multi-stimuli effects on metallic thin films as well as 2D

materials. The goal of the multi-stimuli study using computational techniques such as molecular

dynamics is to assess the compounding effect on microstructural changes, grain growth and

defect annihilation not quantitatively but qualitatively. To mimic the EWF effects in our MD

simulation, we apply additional force on individual atoms during the simulation. Thus, the net

force on each atom is a combination of force due to the interatomic potential and imposed

electron wind force. A flowchart to express EWF implementation in MD simulation is shown as

follows:

15

Figure. 1.11. Flowchart showing EWF implemented in MD simulation using LAMMPS [75]

package.

Where, r is the position, v is the velocity, U is the potential energy, m is the atomic mass, Δt

is the time step, F is the force, 𝑍∗is the effective valence number, e is the electron charge, 𝜌 is the

16

specific resistivity, j is the current density, and a is the acceleration. Similar approach has been

taken to study electromigration failure at higher current density in metal interconnects using MD

simulation [76, 77]. In our study, we have implemented embedded atom model (EAM) potential

to simulate metal and alloys. EAM potentials are well calibrated and has been used to study

metals and alloys [78, 79]. Functional form of an EAM potential for pure metals requires at least

cohesive energy, lattice parameter and elastic properties calibration. In addition to these

parameters, alloys require heat of fusion and energy differences among phases for the calibration

process. EAM potentials used in our study are well calibrated and reported in literatures [78, 79].

We have also used reactive empirical bond-order (REBO) [80] and Stillinger-Weber (SW) [81]

potentials for molybdenum disulfide (MoS2) and tungsten diselenide (WSe2) samples

respectively. Both REBO [80] and SW [81] potentials have been calibrated and used to study

mechanical and thermal properties of MoS2 and WSe2 sample respectively. Additionally, recently

developed ReaxFF [82] potential has been implemented to study mechanical properties of MoS2.

The results obtained from the ReaxFF [82, 83] are in good agreement with experimental data,

and this agreement could be attributed to the inclusion of higher number of parameters and

probably more accurate functional form of ReaxFF potentials. However, in this dissertation we

choose REBO and SW potentials to study 2D materials such as MoS2 and WSe2 to perform

simulations within affordable computational time with large simulation cells containing grain

boundaries (GBs).

In this dissertation, we have investigated the effects of various external stimuli on the

microstructure control and properties of different material system across multiple length scale,

and chapters’ summarization are as follows:

17

In chapter 2, we have addressed multi-stimuli effects such as electrical current, electron wind

force (EWF), Joule heating, and mechanical strain effects at nanoscale. To accomplish this goal

we have investigated electrical current flow effects below the electromigration failure limit in

nanocrystalline 100nm thin films. We observe one order of magnitude higher grain growth

indicating concurrent effects of electron wind force and Joule heating specifically target the grain

boundaries, producing much higher grain boundary mobility compared to high temperature

annealing alone. Afterwards, we have investigated multi-stimuli synergy i.e., a combined effect

of electron wind force, joule heating and mechanical strain effects. Our in-situ TEM study

reveals that application of mechanical strain even in the elastic range (about 0.1%) dramatically

increased the grain size almost instantaneously at the low temperature regions. This is

unprecedented because the annealing and recrystallization literature is essentially founded on the

concept of plastic strain initiated defects that act as nuclei of recrystallization. In the subsequent

section, we have explored multi-stimuli synergy on micro- and bulk-scale samples considering

scaling up issues. Our studies show that the accelerated atomic and defect mobility induced by

multi-stimuli can be exploited for microstructural control of additive manufactured alloys.

In chapter 3, we have considered electrical current effect on non-thermal microstructure

controlling process at extreme length scale i.e., <10nm. To assesses the effectiveness of low

temperature processing at this extreme length scale, we choose two-dimensional (2D) material

with few layers thin (2~3 layers) nanocrystalline molybdenum disulfide (MoS2). Our study

reveals that moderate current density gives rise to atomic scale mechanical force whenever the

electrons encounter defects in the lattice or grain boundaries. The effectiveness of the defect

annihilation is reflected on the physical properties improvement of 2D materials after processing.

Later on, we implemented this low temperature processing on back-gated tungsten diselenide

18

(WSe2) field effect transistor (FET) for real world application. Our study indicates that EWF

driven low temperature processing can effectively annihilates defects and results in two order of

magnitude higher output current.

In Chapter 4, we shifted our material system to thin films transistor with more complex

geometry and shape. We chose GaN based high electron mobility transistor (HEMT). Due to the

specimen geometry and delicate interface, thin film HEMT inherently possesses residual stress.

Again, depending on the mode of operation external stimuli such as voltage biasing can induce

high thermal, mechanical and electrical field in HEMT devices. Thus, HEMT devices are ideal

candidates to assess the effectiveness of the multi-stimuli effects on thin films, devices and

interfaces. This study paves the path for the real-time operation of HEMT inside a high-

resolution electron microscope. Though this study is not directly related to the micro-structural

control process, it explains the degradation mechanism of complex material system under the

combined effects of multiple stimuli.

In chapter 5, we have considered real-world application of our newly developed low

temperature processing technique to eradicate irradiation-induced defects. The objective is to

recover or enhance the physical properties by actively controlling defects and microstructure.

Irradiation is considered as an energetic process where highly energetic particles (ions, neutrons,

electrons) collide with atoms in materials, energizing and displacing them from their original

lattice positions, thereby generating various types of defects. In the first section of this study, we

explore the effects of irradiation on HEMT and their subsequent failure mechanism under the

effects of external stimuli. In the next section, we explore the effectiveness of electron wind

force in annihilating defects originating from irradiation damage. Study reveals that EWF can

efficiently eliminate defects in irradiated materials even at low temperature.

19

Chapter 2

External Stimuli Sensitivity in Thin Films and Additive Manufactured Alloys

Contents of this chapter are based on the following journal articles:

Zahabul Islam, Baoming Wang, Aman Haque, Current density effects on the microstructure of

zirconium thin films, Scripta Materialia,Volume 144, Pages 1359-6462, 2018.

Author of this dissertation designed the experiment, performed the sample preparation,

device design and fabrication, experimentation, data analysis, computational modeling as

well as computational data analysis and manuscript writing. Baoming wang assisted on

experiment design and manuscript writing. Aman Haque guided on experiment design,

and involved in data analysis as well as manuscript preparation.

Zahabul Islam, Huajian Gao, Aman Haque, Synergy of elastic strain energy and electron wind

force on thin film grain growth at room temperature, Materials Characterization, Volume

152,Pages 85-93, 2019.

Author of this dissertation designed the experiment, performed the sample and device

preparation, experimentation, data analysis, computational modeling as well as

computational data analysis and manuscript writing. Huajian Gao guided on

computational data analysis, discussion and manuscript preparation. Aman Haque guided

on experiment design, and involved in data analysis as well as manuscript preparation.

Daudi Waryoba, Zahabul Islam, Ted Reutzel, Aman Haque, Electro-Strengthening of the

Additively Manufactured Ti-6Al-4V Alloy, (submitted)

20


experimentation, computational modeling as well as computational data analysis and

manuscript writing. Daudi Waryoba also conducted the experiment, data analysis and

manuscript writing. Ted Reutzel was involved in additive manufacturing of the Ti64

alloy and manuscript writing. Aman Haque guided on experiment design, and involved in

data analysis as well as manuscript preparation.

2.1 Temperature-Electron Wind Force Synergy in Thin Films

In this section, we investigate the multi-stimuli effects such as electrical current flow in

nanocrystalline zirconium thin films using in-situ Transmission Electron Microscope (in-situ

TEM) and molecular dynamics (MD) simulation. We observed at least one order of magnitude

higher grain growth at current density of 8.5x105 A/cm2 (Joule heating temperature 710 K) in 15

minutes compared to conventional thermal annealing at 873 K for 120 minutes. Both experiment

and simulation results support our hypothesis that the concurrent effects of electron wind force

and Joule heating can produce much higher grain boundary mobility compared to high

temperature annealing alone, and lead to grain growth.

2.1.1 Objective and Motivation

Nanomaterials show strong grain size dependence of their physical properties [4, 5]

across length scales [2], which has motivated the pursuit for microstructural optimization and

control. In a conventional thermal annealing temperature is used as the stimulus for controlling

grain size, phases and defect density. For most of the metallic materials, this temperature is in the

range of 0.3-0.4 𝑇𝑚 (where 𝑇𝑚 corresponds to homologous temperature). However, the applied

temperature field is uniform, targeting both crystalline and defective regions, thus making the

21

process time consuming and energy inefficient. Present study proposes that electrical current

could potentially achieve similar or higher grain boundary and defect mobility at lower energy

and time input. This is due to the pronounced scattering at the grain boundaries and defect sites

[51, 52], effectively enhancing atomic mobility exactly where it is needed for grain growth, and

not uniformly across the sample.

High current density effects are typically considered to be a cause for degradation of

microelectronic interconnects through electromigration [52, 84]. Beyond a critical density, mass

transport takes place due to the electron momentum transfer, particularly intensified at the

defective areas such as grain boundaries (GBs). Other current density studies have focused on

electro-plasticity [85], a phenomenon where electrical current flow induces plasticity in materials

that are otherwise very hard and brittle. To study the fundamentals of electrical current density

effects on microstructures, we adopted a combined experiment-simulation approach. The

experiments were performed inside a Tecnai-LaB6 Transmission Electron Microscope (TEM).

The high resolution imaging and selected area diffraction (SAD) modes make TEM first choice

in visualization and characterization of microstructural changes [86]. However, the challenge in

this technique is the very small work envelope of the TEM chamber, typically accommodating 3

mm diameter grids for specimens [9]. Molecular dynamic (MD) simulation code has been widely

implemented to study mechanical properties [87-89] and electro-migration failure [90]. The

primary modelling challenges are incorporating electron-matter interaction during transport

directly. To overcome this barrier, in our MD modeling approach, we represent the effect of the

electrical current by applying an equivalent electron wind force (EWF) and observe the resulting

atomic/defect migration. The discrepancy in time and length-scales between experiment and

modeling makes it impossible to reach a quantitative agreement. We therefore pursue qualitative

22

and mechanistic contributions from the computational modelling efforts to interpret the

experimental observations.

2.1.2 Materials and Methods

In this present study, we investigate the grain growth mechanism due to the electrical

current flow in zirconium thin films. Zirconium is a transition metal with a hexagonal closed

pack (hcp) lattice structure with high melting point (2128 K), biocompatibility, good corrosion

and radiation resistance, making it a popular choice in nuclear, aviation and surgical implant

applications. We used physical vapor deposition (PVD) technique to deposit about 140 nm thick,

99.97% pure zirconium films on silicon-on-insulator (SOI) substrates. The as-deposited films

showed near- amorphous structure. We used SOI substrate to co-fabricate a micro-electro-

mechanical (MEMS) device with the specimen. The specimen dimension was about 100 microns

long, 5 microns wide. We used standard photo-lithography, lift-off and deep reactive ion etching

on a 100mm (i.e., 4 inch) wafer so that the actuator and heater structures were co-fabricated with

the specimen which ensures perfect specimen alignment and gripping. The chip fabrication

process started with photolithography to transfer a lift-off pattern of the device structure. The

pattern features 100 microns (μm) long and 5 μm wide dog-bone shaped specimens, electro-

thermal actuators and reference points to measure the applied strain. The SOI wafers with 20 μm

device and 2 μm buried oxide (BOX) layers allowed these features to be electrically insolated, so

the mechanical grips also worked as electrodes. The zirconium thin film was evaporated using e-

beam on the lift-off patterns. The wafer was then dry etched with deep reactive ion etching,

which realized the actuators and electrodes with vertical side walls in the device layer. The wafer

was then patterned and the entire handle layer is subsequently etched from the backside. When

the BOX layer was dry etched, all the micromachined silicon structures became freestanding. We

23

then used a physical shadow mask to isotropically etch the silicon beneath the specimen gauge

section, rendering it freestanding. Details of the device design and fabrication are given

elsewhere [91]. Figure 2.1a shows the freestanding zirconium thin film on MEMS device, where

the heavily doped silicon structures act as electrodes. The device fits in a TEM specimen holder

with electrical biasing capability. In our experiment, we passed electrical dc current below

electromigration failure limit (i.e., <106 A/cm2) through the electrodes A and B as shown in

Figure 2.1a to conduct the electrical annealing without damaging the specimen. The inset of

Figure. 2.1a shows the selected area electron diffraction (SAED) pattern of a specimen before

passing electrical current, where the completely diffused rings suggest near-amorphous (< 5 nm

grain size) microstructure without any porosity (Fig. 2.3a). MD simulation model (Figure 2.1b)

was prepared with ten grains with similar size but oriented at different angle. We also performed

electro-thermal simulation of Joule heating using COMSOL® to determine the temperature field

along the sample length (Figure 2.1c). In our in-situ TEM experiments, we passed dc current

through the specimen in a stepwise fashion. Since TEM cannot measure temperature field

directly, this information was obtained from multiphysics simulation of the specimen with actual

geometry, current density and resistance. Simulated temperature profile along the specimen at a

current density of 8.5x105 A/cm2 under vacuum condition mimicking the TEM chamber is

shown in Figure 2.1c. The highest temperature developed at the mid-section of the sample and

was about 710 K.

The effect of electrical current on grain growth mechanism in zirconium thin film was

studied using classical MD simulation conducted by LAMMPS [75] software using Embedded

atom method (EAM) [78] potential. Voronoi tessellation-based models of hcp zirconium were

built with 10 numbers of grains with an average size of 8nm. These grain sizes were chosen to

24

Figure 2.1. (a) Scanning electron microscope (SEM) micrograph of the MEMS device showing

the current flow through the specimen. Inset shows diffraction pattern at 0 A/cm2 current density,

(b) Atomistic model with grains oriented at different angles, and (c) Electro-thermal simulation

of sample with actual geometry, resistance and current density [92].

mimic the as-deposited specimen in the earlier phases of electrical annealing. In our model we

orient the grain at different angle namely 0˚, 5˚, 10˚, 15˚, 30˚and 45˚ as shown in Figure 2.1b,

whereas 0˚ angle lies along [1210] direction and [0001] direction corresponds to film normal i.e

c-axis. We checked the model for overlapping of atoms at the grain boundaries. At first, we

performed energy minimization using conjugate-gradient (CG) method followed by NPT

dynamics for several thousand steps in LAMMPS. We used Huntington-Grone [93] ballistic

model to apply equivalent EWF on each atom. The EWF on each atom is calculated using the

following equations [94]:

𝐹𝑤𝑖𝑛𝑑 = 𝑍∗ × 𝑒 × 𝑗 × 𝜌 (2.1a)

25

Where, 𝑍∗is effective valence number, e is electron charge, j is the current density and 𝜌 is the

specific resistivity of zirconium. In our present simulation, we consider 𝑍∗ as 3.4 [95] and 𝜌 as

421 nΩ.m[96]. During our simulation, we applied periodic boundary conditions in all directions.

Verlet algorithm was employed for time integration during the NPT dynamics with a time step of

0.5fs. EWF was applied on individual atom followed by energy minimization and NPT dynamics

run. We set the simulation temperature at 710 K obtained from electro-thermal simulation by

considering Joule heating effect during the current flow through the sample.

2.1.3 Results and Discussion

Figure 2.2 shows the experimentally observed grain growth during the dc current passage

through the specimen inside a Tecnai LaB6 TEM. We allowed 5 minutes delay between two

consecutive current increments. We continuously monitored grain growth and took TEM bright

field (BF) and selected area electron diffraction (SAED) to probe the grain growth. TEM BF and

associated SAED images indicating microstructural evolution are shown in Figure 2.2a-2.2c. In

our experiment, we observed very fast grain growth dynamics at a current density of 8.5x105

A/cm2 (Figure 2.2b), where the microstructural changes were discernible within few minutes.

Vigorous grain growth was observed at an accelerated current density loading of 1.1x106 A/cm2,

discernible in few seconds. Inset micrograph in Figure 2.2a and 2.2d represent TEM diffraction

patterns for the initial and final conditions in only 15 minutes time span.

We also performed thermal annealing on specimens to assess the effectiveness of

electrical current annealing. In case of thermal annealing, we had to anneal the specimen at 873

K to see any appreciable growth, which is higher than the Joule heating temperature due to the

current flow. The process was very slow, taking 8 times as much time as allowed in the current

26

Figure 2.2 In-situ TEM study indicating grain growth as a function of dc electrical current

density [92].

annealing experiment. Figure 2.3d shows electrical current annealing to produce at least one

order of magnitude larger grain size compared to thermal annealing (Fig. 2.3a). This is also

reflected by the more resolvable spots in Figure 2.3e compared to Figure 2.3b, where the

diffraction pattern of thermally annealed specimen shows only diffused ring patterns. To explain

the observed phenomena, we hypothesize that current annealing efficiently eliminates defects

and dislocations localized around the defective regions such as grain boundaries (GBs). It is well

known that electrical current annealing involves both EWF and Joule heating. Conventional

thermal annealing involves bulk heating of materials whereas electrical current induced EWF

initiates defects annihilation at the targeted locations such as GBs. Thus, thermal annealing is

more energy intensive compared to the electrical current annealing due to the heating up of the

entire sample. We also did not observe any damage in Figure 2.3d, which confirms that electrical

annealing below electro-migration failure limit (i.e., current density < 106 A/cm2) can potentially

be a time and energy efficient path towards active microstructure control.

27

We achieve a qualitative validation of our hypothesis by performing MD simulation that

indirectly captures current flow effect by imposing EWF on individual atoms. Figure 2.3c and

2.3f show MD simulation cell after thermal and electrical annealing respectively. The electrical

annealing was conducted for 25 ps followed by energy minimization and 50 ps NPT dynamics

run. While total thermal annealing takes 1.1 ns, which included first stage heating from 300K to

873K with a temperature ramp rate of 0.012 K/fs, second stage annealing by holding temperature

at 873K for 1.0ns and final stage cooling from 873K to 300 K at a cooling rate of 0.012 K/fs. We

then equilibrated the system for 100 ps at room temperatures i.e 300 K. Thermal annealing

occasionally led to grain boundary (GB) reconstruction (indicated by arrows in Figure 2.3c),

while other grains remain mostly intact. Due to the presence of defects, GBs are at higher energy

state compared to the interior crystalline regions. Thus, any external driving force such as

temperature, strain or electrical current will increase grain size by reducing the GBs area.

External stimuli such as temperature increases the kinetic energy of atoms, which also increases

atomic vibrations at the GBs. Above recrystallization temperature, the thermal driving force

minimizes the GB energy by annihilating defects at the GBs. However, thermal annealing

requires uniform heating of entire material. On contrary, electrical current loading involves both

EWF and Joule heating [97, 98]. EWF accompanied by Joule heating generates driving force that

eliminates GBs defects. Our simulation results as shown in Figure 2.3f shows defect annihilation

and high atomic diffusion due to the pronounced scattering at the GBs during electrical

annealing, which yields larger grain size.

28

Figure 2.3. Comparison between thermal and electrical current annealing: (a) TEM bright field

image after thermal loading (b) corresponding SAED pattern, (c) MD simulation cell showing

grain boundary reconstruction in limited locations indicated by arrows, (d) TEM BF image after

current loading, (e) corresponding SAED pattern, and (f) MD simulation showing grain

boundary reconstruction due to the electrical current loading [92].

Atomistic simulation qualitatively validates our hypothesis. For example, it is evident

from Figure 2.3f that atomic re-orientation and diffusion are dominant at the GBs during the

electrical current flow. Merging of the GBs are seen to be a result of diffusional motion of the

defects under the impetus of the EWF. GBs experience localized stress field as shown in Figure

2.4a, which could be attributed to the local disorder in atomic position, orientation and defect

density at GBs. The localized stresses in the GB regions also indicate higher potential energy

29

states compared to the grain interior atoms. The mechanical stress field around the defects is

another reason behind the localized and targeted enhancement of atomic mobility at the GBs.

Figure 2.4b shows two triple points with an initial orientation at 10°, 15°, 30° and 45° as

indicated by arrows. After electrical annealing, we observed disappearance of GBs as shown in

Figure 2.4c. The reconstruction of original hcp-crystalline structure of zirconium from different

grain sites (as shown in Figure 2.4c) clearly indicates that the grains grow in zirconium thin film

due to the synergy of EWF and Joule heating at the grain boundaries. We also noticed inter

granular diffusion, which provides an evidence of higher mobility of atom due to the pronounced

electron scattering at the GBs. Additionally, we also noticed that all the GBs are oriented along

the same direction (~0°) as shown by arrows in Figure 2.4c after electrical annealing. These

evidence indicate that electrical current could significantly eliminate defects and increase the

grain size in thin films.

Figure 2.4. Time evolution of grain growth obtained from MD simulation trajectory: (a) initial

structure, (b) two triple points before electrical annealing, and (c) two triple points after electrical

annealing [92].

2.1.4 Conclusion

To summarize, we have performed in-situ TEM experiments on near-amorphous

30

zirconium thin films to gain fundamental understanding of external stimuli such as electrical

current density effects on grain growth. We noticed almost two orders of magnitude grain growth

in less than 15 minutes at a current density of 8.5x105 A/cm2. This was about one order of

magnitude higher than what we observed in similar specimens but under convention thermal

annealing at 873 K for 120 minutes. MD simulation results show that the effect of EWF is to

impart very high atomic mobility that is localized to the defects and GBs, which makes the

electrical annealing more energy efficient compared to the conventional thermal annealing where

the crystalline grain interiors are heated to the same temperature as the defective areas. Localized

stress fields around the defects also increase the atomic mobility under EWF thus increase grain

size by annihilating defects at the GBs. The findings of this study may play vital role in

developing novel energy and time efficient techniques for active microstructural reorganization

and control in the near future.

2.2 Multi-Stimuli (Electron Wind Force and Mechanical Strain) Effects in Thin Films

In this section, we will investigate low temperature grain growth using a synergy of

electron wind force (EWF) and mechanical strain. It is well known that thermal annealing is

commonly used for defect elimination and grain growth in polycrystalline materials. Here, we

propose an alternate route through a synergy of electrical current and tensile strain at or near

room temperature. Our experiments involve flow of electrical current below the electromigration

limit in 100 nm thick freestanding palladium films with approximately 5 nm initial grain size.

Electro-thermal simulation shows that Joule heating increases the temperature up to 470 K

(homologous temperature of 0.25) at the middle section of the specimen, whereas, the massive

heat sinks at the two ends of the specimen constrain them to remain at room temperature. In-situ

transmission electron microscopy (in-situ TEM) study shows that more than two orders of

31

magnitude grain growth in the high temperature regions (middle section of the sample) and little

growth at the room temperature regions. However, application of elastic strain (about 0.1%)

dramatically altered the scenario and increased the grain size by more than 10 times in a few

seconds near the room temperature regions. Our finding indicates that the multi-stimuli synergy

of elastic strain energy and electrical current density may achieve grain growth in metallic

materials even at room temperature. Molecular dynamics (MD) simulation of this phenomenon

reveals that the externally applied strain is localized at the grain boundaries (GBs) in

nanocrystalline metals, which promotes the effects of electron wind force on the GB’s atoms.

Thus, we conclude that synergy of two or more stimuli can achieve grain growth at room or even

lower temperatures.


In this section, we wish to study the role of externally applied strain in microstructural

control. For example, grain size is considered to be one of the most influential parameters

controlling mechanical properties of polycrystalline materials [5]. Due to the high fraction of

grain boundaries (GBs) nanocrystalline materials exhibit pronounced grain size effects as

manifested through the breakdown of the classical Hall-Petch relation [6]. Grain size also

influences electrical, thermal and optical properties. [7] [8] [21]. It is well known that GB effects

are usually stronger than those of line or point defects. Additionally, elimination of residual

stress in a sample is critical for both performance and reliability perspective of interconnects

[35]. Therefore, high temperature annealing is a common post-processing step in fabrication.

Grain size and residual stress in interconnects depend on parameters related to the

deposition process, temperature, pressure, substrate etc., and are difficult to control during

32

deposition. Additional parameters such as intrinsic stress, micro-texture, defect density and

surface roughness need to be considered. To overcome these issues, thermal annealing is widely

used, where temperature is raised (> 0.3 homologous temperature) to promote atomic diffusion

or defect mobility. This thermal stimulus activates the microstructural evolution process that is

governed by the minimization of surface, interface and mechanical strain energy [99]. Grain

growth is a mechanism for such energy minimization and is not a monotonic function of

temperature. Rather, the initial grain size and texture, residual stress, annealing environment and

film thickness are vital to grain growth. However, for complex multi-layer systems it is not

possible to raise temperature arbitrarily without introducing thermal stresses.

Thus, motivation for this study comes from the above-mentioned challenges in grain size

and residual stress control. In particular, we propose a non-thermal route towards grain size

control using multi-stimuli such as electron wind force (EWF), temperature and mechanical

strain. Our basic premise is that, acceleration of grain growth kinetics may be achievable even at

lower temperatures with the synergy of stimuli such as electrical current and mechanical strain.

The effect of electrical current is to (a) raise the temperature (i.e., Joule heating) and (b) apply

EWF at the defective areas (such as GBs and triple junction) due to the transfer of the electron

momentum. EWF creates directed-diffusion of atoms that can take different paths, such as grain

interior, boundary and external surface [100], which decreases atomic migration energy barrier

[101], and facilitates atomic rearrangement for energy minimization. Since our goal is to develop

non-thermal microstructural processing, a unique feature of this study is to remove the heat

generated by current in order to separate the effects of temperature. We achieved this through

gripping freestanding thin film specimen with massive heatsinks. Thus, when we pass current,

the center region remains heated, while the edge regions are constrained to the ambient

33

temperature. By doing this we decouple the effect of Joule heating and the EWF in the same

specimen.

Without temperature effects, it is expected that the electron wind force may not have the

enough atomic mobility needed for annealing or grain growth. We therefore attempted to utilize

mechanical strain as a promoting stimulus. For nanocrystalline metals stress assisted grain

growth is observed due to the higher GB mobility and low activation energy [102]. These have

been also reported by literatures at temperatures as low as ambient [103-105]. However, stress-

assisted grain growth requires relatively high levels of stresses and dependent on types of

material. Nevertheless, the focus of this study is the synergy of mechanical strain and EWF,

which we hypothesize to have the potential to provide the atomic mobility even at lower

temperatures.


To investigate the synergy of electric current and mechanical strain on grain growth

kinetics we choose 100nm thick palladium films. Palladium melts at 1555 ºC and is used in

catalysis [106] and electronics applications [107]. The experimental setup consists of 5mm x

3mm silicon-on-insulator (SOI) chips, each containing 100 nm thick palladium film specimens

which are integrated with electrodes, micro-heaters and mechanical actuators as mentioned in

section 2.1.2. Figure 2.5a shows a zoomed-in view of the chip, which is designed to fit in an in-

situ transmission electron microscope (TEM) specimen holder. We also develop molecular

dynamics (MD) models mimicking the in-situ TEM experiments to understand the fundamental

mechanisms behind the multi-stimuli annealing process qualitatively.

34

Figure 2.5a shows a TEM image and diffraction pattern of the as-deposited specimen.

Figure 2.5b shows a COMSOL® multi-physics model for a simulation of the temperature

distribution at current density of 7x105 A/cm2. The relatively massive silicon grips act as heat

sinks, bringing down the temperature to the ambient. This unique boundary condition allows us

to isolate the temperature effects from EWF.

Figure 2.5. In-situ TEM experimental and MD simulation setups: (a) Scanning electron

micrograph of the adopted MEMS device with actuators and electrodes including a TEM bright

field (BF) image and diffraction pattern of as-deposited specimen, (b) temperature profile along

the length of the sample obtained from electro-thermal simulation mimicking the actual

experimental conditions, and (c) atomistic simulation cell with randomly oriented grains used to

mimic the experiments [108].

35

The in-situ TEM experiments were performed inside a FEITM Talos F200X

scanning/transmission electron microscope (S/TEM) with a resolution of 0.12 nm using field

emission gun and 200 kV acceleration voltage. In-situ TEM enables real-time high-resolution

imaging as well as selected area electron diffraction (SAED). In a typical experiment, we first

bias the specimen while measuring the current. About 5 minutes of time gap is allowed between

the current increments to discern the ensuing microstructural changes. We increased the bias as

soon as the relative magnitude of grain growth per current increment decreases. At this point, we

activated the thermal actuators to apply a strain on the palladium thin film. The resulting strain is

measured by comparing the extension of the fixed end of the specimen with respect to the

moving end. We observe the two ends as well as the middle section of the specimen during the

straining.

We prepared MD simulation (Figure 2.5c) cell with 22 randomly oriented grains with

approximate grain size of 5nm. We performed electro-thermal simulations of electric current

flow and subsequent temperature field due to Joule heating using COMSOL®. During this multi-

physics simulation, we considered actual geometry, resistance and current density of the sample

and later on, we fed temperature in the MD simulation model. Figure 2.5b shows the temperature

profile along the specimen at a current density of 7x105 A/cm2 mimicking the TEM chamber.

The highest temperature 470 K, which is < 0.25Tm (where Tm is 1828 K for palladium) was

observed at the middle section of the sample. Additionally, we employed classical MD

simulation to investigate electrical current induced grain growth mechanism in palladium thin

film using LAMMPS [75] package. In our present simulation, we chose a time step of 0.5 fs.

Voronoi tessellation-based simulation cells of face-centered cubic (fcc) Palladium were built as

36

shown in Figure 2.5c. This grain size was chosen to mimic the as-deposited specimen as well as

grain size distribution in the earlier phases of electrical annealing.

During our simulation, periodic boundary condition was applied along the longitudinal

direction (x-direction) while keeping free surface boundary conditions along transverse and film

normal directions (y and z directions respectively). Energy minimization was carried out using

conjugate-gradient (CG) method followed by NVT dynamics at 470K for several thousand steps.

EWF was applied on individual atoms followed by energy minimization and NVT dynamics run.

To describe the interatomic forces among atoms in a polycrystalline palladium thin film we

employ the embedded atom method (EAM) [79] potential. In a classical MD simulation it is

difficult to directly employ electron effects for a large system, thus we first quantify the

equivalent wind force on each atom using the Huntington-Grone [93] model using the following

equations [94]:

𝐹𝑤𝑖𝑛𝑑 = 𝑍∗ × 𝑒 × 𝑗 × 𝜌 (2.2a)

where 𝑍∗is the effective valence number, e the electron charge, j the current density and 𝜌 the

specific resistivity of palladium. In our present simulation, we consider 𝑍∗ as -5.1 [95]. Thus

total force on an individual atom during the electrical current flow can be written as follows:

𝐹𝑇𝑜𝑡𝑎𝑙 = 𝐹𝐸𝐴𝑀 + 𝐹𝑤𝑖𝑛𝑑 (2.2b)

In order to mimic the experiment and investigate the effect of strain on grain growth

under EWF, we apply tensile strain on the palladium polycrystalline simulation cell along its

longitudinal direction (i.e. the x-direction) at a strain rate of 5×108𝑠−1. Visualization of the grain

37

growth was conducted on OVITO [109] software. During the simulation, we also monitor

potential energy and stress distribution on individual atoms.


Figure 2.6a shows TEM bright-field image of a 100 nm thick physical vapor deposited

(PVD) palladium films. SAED pattern indicates that as-deposited specimens initially had a near-

amorphous (<5 nm grain size) structure (Figure 2.6b). To investigate electron beam irradiation

effect, we kept the sample for prolonged period (i.e., 20 hours) under the electron beam and we

did not observe any discernible grain growth. Apparently, electron beam has no/insignificant

effect on grain growth or microstructural changes as shown in Figure 2.6a. Experimentation

outside TEM (i.e ex-situ) also yields similar order of magnitude grain growth, which indicates e-

beam does not have any significant effect on grain growth.

Figure 2.6. (a) TEM bright field (BF) image of the as-deposited specimen after prolonged

exposure to the electron beam, and (b) corresponding SAED pattern [108].

38

Figure 2.7 shows the experimental results on electrical current induced grain growth at

the middle section of the palladium specimen. We noticed grain growth is a non-linear function

of current density and we did not increase current density beyond 7 × 105 𝐴/𝑐𝑚2 to avoid any

electromigration damage. Theory indicates that Joule heating is a nonlinear function of current

density whereas EWF is a linear function of current density. Thus, 7% increment of current

density can increase the Joule heating by 15% as shown in Figure 2.7b and 2.7c. After a current

density of 6.5x105 A/cm2, nonlinear increment of Joule heating accompanied by electron wind

force induced significant grain growth as shown in Figure 2.7b and 2.7c. The average grain size

after electrical annealing was 550 nm, which is more than two order of magnitude higher than

the as-deposited grain size. Grain growth is mostly uniform, however the SAED pattern in Figure

2.7f indicates significant change in texture. Electrical current induced grain growth kinetics is an

order of magnitude faster than thermal annealing, and takes only 5 minutes at each current

values.

It is important to note that electrical current effect in conductive materials has been

consistently demonstrated to change microstructure of metallic materials. Thus, electrical current

accompanied by temperature field is known to play a dominant role in enhancing atomic

mobility. In the previous section, we have argued that the role of EWF is pronounced when the

specimen has appreciable volume fraction of defects [110] (such as grain boundaries), where

momentum transfer is immensely amplified. Thus, the decoupling of thermal contribution will

allow us to access fundamental understanding of electrical flow related phenomena in metals

[111]. In this respect, configuration of our freestanding thin film specimen allows us to study the

role of temperature (in the middle section) and electron wind (at the edges, connected to the grips

that act as massive heat sinks) separately in the same specimen. The electro-thermal simulation

39

as shown in Figure 2.5b shows the temperature profile along the specimen length due to the

Joule heating during the electrical current flow. From the TEM BF it is clear that we have very

large grain growth at the middle region of the sample, while very little effect is observed due to

the near room temperature at the edges.

Figure 2.7. In-situ TEM BF and SAED evidence of grain growth in the specimen center region at

(a, d) 0 x105 A/cm2, (b, e) 6.5x105 A/cm2, and (c, f) 7x105 A/cm2 current densities [108].

Figure 2.8a shows the microstructure at the anode end of the specimen, where the current

density is 6.7x105 A/cm2, but the temperature remains at the ambient. After annealing, the

average grain size is about 10 nm. This is significantly lower compared to the center region,

suggesting that the EWF alone cannot induce significant amount of grain growth at room

temperature. This observation motivates us to study the synergy of mechanical strain and

40

electrical current on grain growth at room temperature. While it is well known that surface and

strain energies are primary driving force for the annealing and recrystallization processes.

However, most of the studies in the literature involve very large, plastic strain [112] in the form

of cold work or field-assisted sintering. Here, we explore the effect of strain energy in the elastic

regime, which is yet to be studied. We hypothesize a synergy between elastic strain and EWF

[113] might enhance grain growth. Experimental finding on synergy study is shown in Figure

2.8b, where even as small as 0.1% mechanical strain exhibits profound impact on the grain size

at room temperature.

Figure 2.8. Specimen microstructure at 7x105 A/cm2 current density and room temperature, (a)

before, and (b) after application of 0.1% strain [108].

To study the effect of elastic strain energy, we applied electrical bias on the on-chip

actuator pad. To stay in the elastic region, we keep the applied strain in the specimen at about

0.1%. Strain was measured by tracking the moving end of the specimen with respect to a fixed

reference frame on MEMs device. While tracking a region near the edge (and hence around the

41

room temperature), we applied this strain in a single step. We almost instantaneously observed a

dramatic increase in the grain size as shown in Figure 2.8. The significant (>10x) increase in

grain growth suggests that even at very low temperature (i.e., room temperature), there exists a

strong synergistic effect of elastic strain and electrical current density in promoting grain growth

dynamics. Stress induced grain growth at high deformation have been reported [114, 115].

However, in our present study, instead of applying high stress, we coupled lower strain (in the

elastic regime) and EWF. This coupling yields significant amount of grain growth even at room

temperature. This grain growth is attributed to the pronounced interaction between multi-stimuli

(strain and current density) and pre-existing high-volume fraction of defects at the GBs. Defects

(which could be an outcome of plastic deformation) are required for recrystallization. Due to the

high-volume fraction of GBs in nanocrystalline materials, they contain large number of defects

and stored elastic energy in terms of defects and disorders compared to their bulk

counterpart. Additionally, in a nanocrystalline material stresses are highly localized at the grain

boundaries (GBs). Due to this localized stress, a small amount of applied strain could amplify its

effects at the GBs, which indicates lower external strain could specifically targets the grain

boundary atoms to enhance atomic mobility. Due to the pronounced scattering of conducting

electrons at the GBs, EWF is also higher at the GBs. Thus, pronounced effects of strain and

electrical current on the GBs atoms could recrystallize the nanocrystalline materials even at

lower strain and temperature.

The literature indicates that plastic strain energy is the core requirement for nucleation of

recrystallization. For example, high density of dislocations in a sub-grain structure possess

higher plastic strain energy compared to the lattice. Therefore, a sub-grain is the first region to

start recrystallization. Therefore, there is no evidence in the literature that elastic strain energy

42

could drive grain growth. We propose a conceptual model and study its effectiveness using a MD

simulation model. Here, the fundamental premise is that the grain boundaries are not just high-

volume fraction of surface defects in the nanocrystalline thin films; they also experience

localized, higher mechanical stress compared to the grain interior. The effect of EWF is

pronounced at the grain boundaries, where the electrons impart their momentum to enhance the

mobility of the atoms. However, an additional driving force is still required for the atoms to

migrate towards minimizing their energy at low temperature. We propose that small amount of

external elastic strain energy is sufficient and the synergy can be strong enough to induce

significant grain growth at room temperature. It is important to note that while the strain energy

is crucial for the process, it has little or no effect without EWF. This indicates that at room

temperature single stimulus might not induce significant atomic mobility for the grain growth.

To comprehend the synergy of electrical current and strain on grain growth mechanism in

palladium thin films, we performed MD simulation as shown in Figure 2.9. Figure 2.9 shows

interior crystalline region of individual grains by hiding the GB atoms. Figure 2.9b and 2.9c

show 1nm wide fixed electrode area at the two edges colored by red bands. As mentioned

earlier, initial sample contains randomly oriented 22 grains with an average grain size of 5nm

prior to the annealing. After the first stage annealing, number of grains reduces to 11 due to the

grain growth. Once we apply the strain and anneal the sample, the grains grow further, and the

total number of grains reduces to 7. GBs contain defective atoms that are at a higher energy state

compared to their interior crystalline regions. Due to the high energy state of GBs, external

stimuli such as electrical current, temperature or mechanical strain could potentially annihilate

some of these defects and increase the grain size.

43

Figure 2.9. MD simulation model on grain growth: (a) initial structure, (b) electrical annealing

prior to the application of strain, and (c) electrical annealing after strain application, (Green color

indicates face centered cubic (FCC) and red color indicates hexagonal closed packed (HCP)

phase of palladium) [108].

44

Recent studies show that [97, 116] electrical current loading could annihilate GBs’

defects efficiently and thus lead to the grain growth. Electrical current which introduces both

EWF and Joule heating [97, 98] could create pronounced electron scattering at the GBs’

defective sites which further reduces GBs’ energy by annihilating randomness at the GBs. As

mentioned earlier, EWF is effective at the GBs, where they could exchange their momentum

with disordered atoms. Due to the imparted EWF and Joule heating effects we notice significant

amount of grain growth compared to the pre-strained sample. The localized stresses around the

grain boundary regions are also an indicator to the higher energy states of the system. The initial

higher stress around the GBs enhances the atomic mobility during the electrical current flow and

allows defective atoms to relax and rearrange at the GBs.

Figure 2.10. Computational study: (a) Strain energy increment during tensile straining of the

system, and (b) grain size as a function of applied tensile strain [108].

Figure 2.10a shows the strain energy increment in the sample during tensile loading of

the sample. However, in our present simulation we applied 1% strain (0.01 nm/nm; Figure 2.10a)

as shown by dotted blue color vertical line in Figure 2.10a. This small amount of strain confirms

45

that the applied strain is within the elastic regime. Due to the time and length scale effects in MD

simulation we applied higher strain (1%) compared to the experimental strain (0.1%). With

several thousand of atoms, small time step and complex forcefield, MD simulations could be

highly time intensive. Thus, we adopted an exaggerated computational strain to capture the

physical phenomenon within affordable computational time. Study shows that elastic strain

contributes ~1meV strain energy per atom into the system. Figure 2.10b shows the synergistic

effect of the electron wind force and mechanical strain on grain growth, and represents the grain

size as a function of applied tensile strain (for the given experimental value of applied EWF). We

notice grain size increment with the applied strain accompanied by EWF. Thus, our study

indicates there exists a strong synergy between EWF and mechanical strain even at room

temperature.

2.2.4 Conclusion

Electrical current induced phenomena such as electro-plasticity and electro-pulsing are

receiving increasing attention due to the rapidness in materials processing compared to the

conventional heat treatment processing. Prior to the commercialization, the fundamentals of the

contributions from the current, temperature and strain must be understood. Unfortunately,

decoupling of these effects is often arduous. To overcome this issue, we developed a unique

experimental setup, where the same specimen could be divided into two distinct regions. The

center region experiences a high temperature while the two edges of the freestanding thin film

specimens are subject to the same current density, but remain near room temperature. In this

case, insignificant amount of grain growth near the edges of the sample suggests the dominant

role of temperature on grain boundary mobility.

46

Present study reports the experimental evidence of multi-stimuli synergy i.e., mechanical

strain and EWF effects on grain growth even at lower temperature. It is well known that surface

and strain energy can act as driving force for grain growth. However, very little is known about

elastic strain driven grain growth. This is because fundamental electro-plasticity or dynamic

electropulsing studies introduce significant amount of mechanical strain during the passage of

current. On the other hand, in this study we propose a new direction where a relatively small

(0.1%) elastic strain is applied on the specimen in conjunction with EWF, resulting in

remarkable grain growth. Interestingly, the EWF manifests at the grain boundaries due to the

higher localized stress at the GBs. The experimental evidence as well as the computational

modeling indicates that the synergy of EWF and mechanical strain can be exploited to achieve

grain growth at near room temperatures.

2.3 Low Temperature Processing of Additive Manufactured Ti64 Alloy

In this section, we will examine scaling up feasibility of electron wind force (EWF)

induced low temperature processing. To accomplish this goal we passed electrical current

through the additive manufactured Ti64 alloy to annihilate defects and reduction of residual

stress without sacrificing mechanical properties. Our results show that both grain size and nano

hardness increased by 15% and 16% respectively after EWF processing. This is attributed to the

pronounced dislocation interactions as well as defect healing during low temperature processing.

Reduction in the residual strain and an increase in the intrinsic strength obtained from electron

back-scattered diffraction (EBSD) study corroborate the effectiveness of the EWF annealing at

microscale sample.

47


Microstructures in an additive manufacturing (AM) process are governed by a complex

combination of processing parameters, which include the scan rate of input energy source, power

of the energy source, deposition rate, and the dwell time between individual layers [1,2]. Thus,

mechanical properties [3,4] of AM parts show large scatter and inconsistency due to the

numerous process parameters, which further results into a wide range of microstructures. Among

many other materials, titanium alloys especially Ti-6Al-4V (i.e., Ti64) alloy has received

considerable interest among material science community due to their applications in aerospace

and biomedical industries [15-17]. Commercially available Ti64 alloy possesses high corrosion

resistance and high specific strength [18]. Owing to the superior properties such as corrosion

resistance, tensile strength, elongation, Young's modulus and fatigue properties of commercially

available Ti64 alloy; AM fabricated Ti64 alloy contains different types of internal defects

[19,20] which negatively affect their properties. Typical AM fabricated Ti64 materials contains

columnar prior -grains that are aligned with the build direction [5-7]. Within these columnar

prior -grains, a mixture of colony -laths, ʹ-martensite, and retained phases is present [5-7].

Both the size and fraction/proportion of the -laths, ʹ-martensite, and size of grains determine

the mechanical properties. Aside from these microstructural parameters, internal macro- and

microscopic defects have also shown to significantly affect the mechanical properties.

Microscopic defects include micropores, residual/transformation stresses, and microstructural

defects such as grain boundary (GB), cracking and non-equilibrium microstructures due to

intrinsically high-solidification rate of the process [8,9]. Likewise, the spatial distribution of

these defects creates unwanted anisotropy in tensile properties [21,22]. Thus, in order to ensure

high quality of AM Ti64 alloy, microstructural defects further need to be reduced.

48

To overcome this issue post-processing methods such as hot-isostatic pressing (HIP) or

heat treatment are often employed to reduce defects and residual stress in AM parts. However,

HIP and other heat treatment operations often result in inferior mechanical properties. For

example, porosity is a common defect in metal AM parts and currently, the most acceptable post-

processing technique to mitigate its effects is to apply hot-isostatic pressing (HIP). HIP is often

applied to AM structures in order to cure process-related voids and pores within the builds.

However, the application of this post-processing treatment has been shown to coarsen the

microstructure and decrease the strength of AM materials [12]. Heat-treatment is another post-

processing method that is used for stress relieving and microstructural modification/enhancement

of AM parts [2,13,14]. However, high temperature annealing that is typically used for optimizing

the microstructure of AM parts results in mechanical properties degradation due to thermally

activated processes such as recovery, recrystallization, and grain growth.

Unlike conventional heat treatment or HIP, in this section, we propose a novel technique

to eliminate microstructural defects at room temperature by employing electric current annealing.

This method uses EWF to promote an enhanced atomic mobility near the defect cores to

annihilate microstructural defects without compromising the mechanical properties. The

investigation was performed on powder bed fusion additive manufactured Ti64 alloy using

optical microscopy, electron backscatter diffraction (EBSD), and nanoindentation

characterization techniques to assess the effectiveness of the method.


The Ti64 sample was manufactured at Penn State’s Center for Innovation Material

Processing thru Direct Digital Deposition (CIMP-3D). The samples were fabricated using an

49

EOS M280 Laser-Powder Bed Fusion machine. The material feedstock for the build was EOS

Titanium Ti64 (which conforms to ASTM B348 Grade 23 powder material specifications), with

particle size ranging between 15 μm to 45 μm. The maximum current density during the low

temperature processing was 5 x105 A/cm2. A custom stage with electrical biasing capability was

built to perform this experiment (as shown in Figure 2.11). To ensure the sample remains at low

temperature processing, we passed liquid nitrogen (N2) through the stage attached with the

sample as shown in Figure 2.11. We controlled the sample temperature at 20°C. During the

experiment, we gradually increased the current by 0.1A per step while continuously monitored

the resistance and microstructure of the sample. Each biasing step was held for 5 minutes. The

sample mounted inside scanning electron microscope (SEM) was checked after every current

step to observe any discernible microstructural change.

Figure 2.11. Schematic showing experimental set-up with temperature controlled stage.

Mechanical polishing followed by ion milling was employed to prepare the sample for in-

situ EBSD study to characterize microstructures. Sample was polished with both SiC sandpaper

up to a P2400 grit size, and with 3 µm diamond slurry. Final polishing was done on a vibratory

50

polisher for at least 2 hours with a chemical-mechanical polishing slurry consisting of 0.05 µm

alumina-silica. Additionally, etching was performed using micro-etchant solution consisting of

1.5 mL HF, 4 mL KNO3, and 94 mL of H2O for optical microscope imaging. Average Grain size

was calculated using ASTM E112–13, Standard Test Methods. We acquired EBSD data using

Helios 661 NanoLab FEI dual beam field emission scanning electron microscope (FE SEM)

equipped with Aztec EBSD Nanoanalysis software, ver. 4.2 (Oxford Instruments). Scanning was

conducted at a step size of 0.1 μm, and a minimum of 3 randomly selected scans were acquired

for each specimen. All EBSD analyses were performed using CHANNEL 5 software. Grain and

-lath boundaries were marked using a threshold of 10°. We also performed nanoindentation

experiment of the as-received and electric current processed specimens using Bruker

Nanoindenter with the Berkovich diamond tip [23] based on the Oliver and Pharr approach [24]

and ISO 14577 standard for measurements.

2.3.3. Results and Discussions

It is well known that during passage of electrical current through a conductive material, it

generates both resistive heating also known as Joule heating and electron wind force (EWF)

effect. As mentioned earlier, electrical current density effect is considered to have negative

impact on materials specifically above a threshold value (i.e., 106 A/cm2) [117, 118]. Thus,

during electrical current annealing one needs to take precaution regarding electrical current

density. In our present study, we keep the current density below this threshold value and actively

cool the sample to investigate the effects of EWF at low temperature. Recently this technique has

been shown to yield comparable microstructure and properties to that of conventional thermal

annealing at a very short duration and a relatively low temperature [34-36]. This has been

attributed largely to a non-thermal effect, rather than the thermal component (Joule heating).

51

EWF induced momentum transfer promotes intensive vacancy diffusion, atomic flux mobility at

the defects and grain boundaries [37,38]. This enhanced atomic mobility has the potential to

annihilate defects. Unlike previous sections, in the following section we will investigate EWF

effects on micro-scale sample to validate the scaling up feasibility of the low temperature EWF

processing.

In our previous section we have calculated EWF using Huntingdon-Grone model [93].

However, EWF can be also calculated form quantum dislocation theory [119]. Based on

quantum dislocation theory EWF develops due to the interactions between drifting free electrons

and defects such as dislocations. From quantum dislocation mechanics, the magnitude of the

EWF depends on the orientation between the Burgers vector and the direction of the current

vector. In addition, based on the principle of virtual work, the direction of the electron wind

force should be normal to the shear slip vector at the point of interest. Thus, considering specific

electrical resistivity due to dislocations, the electron wind force per unit length of dislocation can

be simplified as follows:

Fwind =ρd

Ndejne (2.3.1)

Where, j is the current density vector, ρd is electrical resistivity, Nd is the mobile dislocation

density, e is electron charge, and ne is electron concentration. Titanium has a density

of 4.50 g/cm3, and an atomic weight of 47.867 g/mol, electron density/concentration is 𝑛𝑒 =

2.265 × 1029 m−3, effective valence number Z*=4, electrical resistivity of 1.78μΩ.m [120], and

dislocation density is of the order of 1014 m-2 [121]. Thus, calculated EWF is 3.23N/m. On the

other hand, if we considers slip in the basal and prism systems, these systems have <a>

dislocations with a Burgers vector 𝒃 =1

3< 1120 >, and magnitude of |𝒃| = 𝑎, where a is the

52

lattice parameter for hcp Ti64, and c/a=1.589. Then we could estimate corresponding critical

shear drag force required per unit length of dislocation using the following equations:

𝐹𝐶𝑆𝐹 = 𝜏𝐶𝑅𝑆𝑆𝑏 (2.3.2)

Where, 𝜏𝐶𝑅𝑆𝑆 is critical resolved shear stress for Ti64 is 300~400MPa [122], and b is burger

vector. Thus, calculated required force to active dislocation is 0.118N/m, which is one order of

magnitude lower compared to EWF generated due to electron wind. Thus, we could conclude

that even in the absence of thermal field EWF alone can activate dislocations in the specimen.

Figure 2.12 represents the -lath microstructures of the as-built and electric current

processed specimens. The average calculated grain sizes were 143.9 ± 41.4 µm and 165.8 ± 33.4

µm, for the as-built and electric current processed specimens, respectively. EWF processed

sample shows 15% grain size increment compared to the as-received sample. In our present

study, we also performed EBSD to assess microstructural changes. To investigate the mechanical

properties, we performed a detail analysis of the Taylor factor/Schmid factors on specimens

Figure 2.12. Optical micrographs of Ti64 specimens in the (a) as-built, and (b) Low temperature

EWF processed conditions.

53

before and after low temperature processing. Among three primary slip systems, there is a

consensus that prismatic slip is easier to activate compared to basal and pyramidal slip

[25,27,28]. Pyramidal slip system is very resistant to slip, and its critical resolved shear stress

(CRSS) is reported to be about 7 times higher than that of basal slip system [25,29]. It can also

be argued that he <c + a> slips have Burgers vectors that are significantly longer than for the <a>

slips. Thus, in our present study we excluded pyramidal slip system analysis from our study.

Figure 2.13 shows calculated Schmid factor maps for basal slip system of pre- and post-

processed sample. It is apparent that both microstructures exhibit a basketweave -lath structure.

The widths of the -laths were determined to be 0.46 ± 0.41 µm and 0.53 ± 0.47 µm for the as-

received and EWF processed specimens, respectively. Interestingly, we notice same magnitude

of increment in -lath structure width and grain size i.e., 15%. The Schmid factor maps for basal

slip system show a distinctive change in Schmid factors from high in the as-built specimen to

low in electric current processed specimen. The corresponding Taylor factors are presented in

Figure 2.13c, and it exhibits higher values of Taylor factors for electric current processed

specimen compared to the as-built specimen. Figures 2.13d and 2.13e show TEM micrographs of

the microstructure of Ti64 alloy in the as-built condition and after electric current processing,

respectively. We have compared the same area of interest before and after the application of

electric current as shown in Figures. 2.13d and 2.13e by cyan color dotted circle. As-built

specimen does not contain any twins as shown in Figure. 2.13d. However, upon the application

of the electric current we noticed twin formation due to the electron wind force (EWF) effects as

shown by yellow color arrow in Figure. 2.13e. This twining formation might contribute to the

enhancement of the hardness of the sample after annealing.

54

Figure 2.13. Basal Schmid factor maps for the Ti64 specimens in the (a) as-built, and (b) electric

current processed sample, (c) Calculated Taylor factors; Twining during electrical annealing: (d)

as-built specimen and (e) after applying a current density of 5x103 A/cm2.

In this study, we have also analyzed residual strain in the pre-, and post processed sample

through measurement of local variations/spread in lattice orientations or Kernel Average

Misorientation (KAM) [53,54]. This analysis is done for every pixel in a kernel with a

predefined threshold value of 5°. The threshold is used to exclude well-defined grain boundaries

in the analysis. Thus, KAM study allows us to identify with high residual strain [53,54]. The

higher the residual strain, the higher the KAM value, and vice versa for lower residual strain.

Figure 2.14 clearly indicates significant reduction of residual strain in a post-processed sample.

55

It is well known that AM parts inherently possess significant residual stress due to the highly

non-equilibrium process. Hence, our study suggests that EWF processing has the ability to

eliminate this residual stress from AM parts.

Figure 2.14. KAM maps for the Ti64 specimens in the (a) as-built and (b) electric current

processed conditions. A threshold of 5° was used to exclude well-defined grain boundaries in the

analysis.

To assess the effectiveness of the low temperature processing we measured mechanical

properties of pre- and post-processed samples using nanoindentation test as shown in Figure

2.15. We notice 16% increment in hardness value i.e., from about 5 GPa to ~6 GPa after

processing. While harness value increased significantly, we did not notice any significant change

in elastic modulus calculated from Oliver and Pharr approach [24] as shown in Figure 2.15b.

Higher values of Taylor factors for electric current processed specimen compared to the as-built

specimen (as shown in Figure 2.13c) imply that the slip systems in the electric current processed

specimen provide more resistance to yielding than the slip systems in the as-built specimen,

which further reflects in their hardness value. Thus, we can postulate that the high EWF causes

local reorientation of the microstructure (-laths) and prior β grain that are harder or more

resistance to yield. This study corroborates the scaling up feasibility of EWF processing at low

56

temperature even without any loss in mechanical properties. In our present study, we have

annealed samples with a thickness of up to 100um. However, it is possible to scale up this low

temperature processing technique even for bulk-scale sample with a high-power supply capable

to deliver high output current, and we left this work for future study.

Figure 2.15. (a) Force-displacement plot obtained from nanoindentation experiment, (b)

calculated hardness and Young’s modulus.

2.3.4 Conclusion

In this study, we have presented a novel and energy efficient method of defects

elimination by electric current processing at low temperature without compromising mechanical

properties. Our study shows that electric current induced EWF can enhance mechanical

properties by reducing or eliminating defects from the sample. Investigation shows that under the

EWF defective atoms might have sufficient mobility to cause local reorientation, hence can

exhibit more resistance to deform. Due to this effect, higher external force is required to initiate

plastic deformation in the electric current processed specimen compared to the as-built specimen.

Additionally, study also shows that electric current processing reduces residual strain in the as-

57

built sample. All of these observations corroborate the scaling up feasibility of EWF processing

at low temperature without sacrificing mechanical properties.

58

Chapter 3

External Stimuli Effects on Two-dimensional (2D) Materials and Devices


Zahabul Islam, Kehao Zhang, Joshua Robinson and Aman Haque, Quality enhancement of low

temperature metal organic chemical vapor deposited MoS2: an experimental and computational

investigation, Nanotechnology, Volume 30, Number 39, Pages 395402 (9pp), 2019.

Author of this dissertation designed the experiment, performed the sample transfer,

device fabrication, experimentation, data analysis and manuscript writing. Kehao Zhang,

Joshua Robinson synthesized materials and involved in manuscript writing. Aman Haque

guided on experiment design, and involved in data analysis as well as manuscript

preparation.

Zahabul Islam, Azim Kozhakhmetov, Joshua Robinson, Aman Haque, Enhancement of WSe2

FET Performance Using Low-Temperature Annealing. Journal of Electronic Materials 49, Pages

3770–3779 (2020).

Author of this dissertation designed the experiment, performed the sample transfer,

device fabrication, experimentation, data analysis and manuscript writing. Azim

Kozhakhmetov, Joshua Robinson synthesized materials and involved in manuscript

writing. Aman Haque guided on experiment design, and involved in data analysis as well

as manuscript preparation.

59

3.1 Low Temperature-Electron Wind Force Synergy in Molybdenum Disulfide

In this section, we will discuss low temperature processing of two-dimensional material

such as molybdenum disulfide (MoS2). Crystallinity plays important role in electronic quality of

chemical vapor deposited MoS2, which tends to deteriorate with decrease in deposition

temperature. Thermal annealing can improve the quality but requires very high temperatures. In

this study, we investigate an alternative low temperature (room temperature to 400 C) annealing

process that exploits the electron wind force (EWF) during the passage of current. Electrical

current density gives rise to atomic scale mechanical force whenever the electrons encounter

defects in the lattice or grain boundaries. After hypothesizing that this EWF can significantly

enhance defect mobility even without any temperature field, we demonstrate the process using

in-situ transmission electron microscope (in-situ TEM) and molecular dynamics (MD)

simulation. To validate the hypothesis, we choose monolayer metal organic chemical vapor

deposited MoS2 deposited at 400 C, and process at low temperature. Experimental results show

5 times enhancement in electrical conductivity, which is supported by the selected area electron

diffraction (SAED) patterns indicating significant grain growth. Discrete spots in SAED pattern

also indicate evolution of high crystallinity at low temperature. Computational investigation

shows that atomic mobility, defects healing as well as reorientation mechanism at the grain

boundaries.


Two-dimensional (2D) transition metal dichalcogenides (TMDCs) such as MoS2 show

intrinsic band gap ranging from 1.2eV to 1.8eV depending on the number of layers [123]. MoS2

has attracted the materials community because of their exceptional electrical, optical, and

60

mechanical properties [124-126], thus making it desirable for field-effect transistors (FETs)

[127-129], flexible electronics [130-132] and sensors [133, 134] applications. Due to the limited

flake size and exfoliation-induced defects, exfoliated MoS2 layers remain concerns for scalability

[135-137]. Thus, a need for scalable production of large-area and highly crystalline MoS2 has

spurred advancement in chemical vapor deposition (CVD) processes [138-140]. It is well

established that quality of CVD grown 2D materials is a strong function of deposition

temperature. However, lower temperature deposition tends to decrease the materials quality. For

example, at low temperature multiple nucleation sites on supporting substrate could form

nanocrystalline MoS2 during the CVD growth, with randomly oriented grains connected by

grain boundaries (GBs) [141, 142]. These GBs in 2D materials have detrimental effects on

physical properties such as higher resistivity, lower carrier mobility, poor mechanical properties

and lower thermal conductivity of 2D materials [141, 143-147]. However, quality of low

temperature deposited 2D materials can still be improved with post-processing. Unfortunately,

for MoS2 the required thermal annealing temperatures may be in excess of deposition

temperature i.e., 800 C [148, 149]. From device fabrication process compatibility and residual

stress perspective, lower processing temperatures are desired.

The primary objective of this study is to minimize the defects and GBs in low

temperature CVD grown MoS2. Thermal annealing is a widely used technique to enhance defect

mobility to minimize their volume fraction in the solid, or to minimize the GB density (thus

increasing the grain size). Thermal annealing technique has been also employed for annealing of

2D materials. For example, post thermal annealing of MoS2 to enhance the crystallinity has been

studied by both ex-situ [150] and in-situ [151] experiments. However, reported thermal

annealing temperature was high, i.e., in the range of 700˚C-1000˚C, and time was in the order of

61

hours. A major drawback of, high temperature annealing is detrimental thermal stress in layered

materials. In addition, the required annealing time is also very long because highly uniform

temperature means random diffusion, where both lattice and defective regions are given equal

vibrational energy [152]. Thus in our present study, we propose low temperature electrical

annealing of MoS2. In our experiment, temperature was around 400˚C to avoid any undesirable

thermal stress, and time required for this annealing is in the order of minutes (i.e 5 minutes). We

proposed that defects in electrically conducting materials can be targeted at the GBs using a non-

thermal process, i.e., momentum transfer of electrons at defects. When we pass electrical current

through the sample electron scattering is massive at the defects, where the electrons lose their

momentum [58], developing atomic scale force known as the ‘electron wind force’ (EWF). Such

momentum transfer imparts high defect mobility and creates a ‘directional diffusion’ (compared

to the random diffusion in high temperature heating). Our hypothesis is that such directed, non-

thermal diffusion can be effectively exploited to control defect mobility and migration at low

temperature and high speed. Additionally, temperature from resistive Joule heating can be

helpful in enhancing defect mobility. In this study, a balance of Joule heating (i.e, below thermal

degradation failure) and EWF is achieved through careful consideration of specimen geometry

and thermal boundary conditions. Our proposed electrical annealing technique has been

successfully deployed for conductive thin films, additive manufactured materials and alloys

[116, 153, 154].

In this study, we investigate the proposed EWF induced low temperature annealing

process on 2D materials. Nominally, monolayer MoS2 specimens were synthesized at 400 C

using low temperature metal-organic chemical vapor deposition (MOCVD) technique. The

material was then transferred to a custom designed and fabricated micro-electro-mechanical

62

(MEMS) chip that can be mounted on a TEM specimen holder with electrical biasing capability.

The specimen resistance and selected area electron diffraction (SAED) patterns are recorded for

the input current density. Significant enhancement in electrical conductivity indicates the

effectiveness of the EWF induced low temperature processing. To obtain atomic-scale defect

mobility and migration mechanisms that lead to low temperature EWF induced annealing, we

develop MD simulation models using LAMMPS [75] simulation package.


3.1.2.1 MOCVD Synthesis

Low temperature MOCVD is used to synthesize MoS2, the reactor design and

experimental procedures are described elsewhere [155, 156]. To summarize, 4×10-4 sccm

molybdenum hexacarbonyl (Mo(CO)6) and 0.55 sccm diethyl sulfide (DES) were carried to the

hot wall MOCVD reactor by 50 sccm H2 and 565 sccm Ar. The growth is conducted at 400°C

for 1 hr on 300 nm SiO2 substrate that yielded nominally 1 layer of MoS2.

3.1.2.2 Device Fabrication and Sample Transfer

Wet transfer technique using PMMA and HF solution was employed to transfer MoS2

sample on the MEMS chip for electrical biasing. Details device design and fabrication details are

given elsewhere [91]. Figure 3.1 shows schematic of the transfer technique and experimental

setup. At first, MoS2/sapphire specimen was spin-coated with 950 K PMMA A2 at 1500 rpm.

After spin coating, MoS2/PMMA was released in 16% HF solution from sapphire utilizing the

sacrificial etching and surface tension to peel off the MoS2/PMMA film [157]. Then released

specimen was manipulated on the MEMS device for biasing purpose. Finally, acetone and

63

isopropyl alcohol (IPA) was used to dissolve PMMA film to obtain freestanding MoS2 sample.

The device is then wire bonded and mounted on an in-situ TEM holder for in-situ TEM study.

Figure 3.1. Schematic shows the transfer process of monolayer MoS2 on to a MEMS device and

subsequent experimental setup for in-situ TEM investigation [67].

3.1.2.3 Computational Details

We developed classical MD simulation model using Reactive Empirical Bond-Order

(REBO) potential [80] and implemented it in LAMMPS [75] package. Single layer

polycrystalline MoS2 were modeled with 10 randomly oriented grains with approximate grain

size of approximately 2.5 nm. We incorporate both Joule heating and EWF effects in our

simulation. We chose 0.5 fs simulation time step and maintained periodic boundary condition

along the length of the sample (i.e., x-direction) while keeping free surface boundary conditions

on transverse and film normal directions (i.e., y and z directions respectively). Energy

minimization was performed on the simulation cell using conjugate-gradient (CG) method

64

followed by NVT dynamics for several thousand steps in LAMMPS. Time integration was

carried out using the Verlet algorithm during the NVT dynamics. The EWF was applied on

individual atom. It is difficult for a large system to directly employ electron effects in classical

MD simulation as classical MD simulation solely depends on Newton’s equation of motion.

Thus in this present study, we indirectly quantify equivalent EWF as mentioned in section 2.1.2

on each atom using the Huntington-Grone [93] model. In our simulation, individual atom

experiences a total force that includes force from REBO potential and EWF, and can be written

as follows:

𝐹𝑇𝑜𝑡𝑎𝑙 = 𝐹𝑅𝐸𝐵𝑂 + 𝐹𝑤𝑖𝑛𝑑 (2)

To mimic the experimental condition i.e., direct current (dc) in the specimen, we impart the

electron wind force on individual atoms for 1.0 ns.

Figure 3.2. (a) Observation of almost one order of magnitude reduction in electrical resistance of

MoS2 specimens during EWF annealing, and (b) spatial temperature distribution for the highest

current density [67].

65


To investigate the effectiveness of the electrical current annealing induced quality

enhancement in low temperature MOCVD MoS2, we conducted in-situ TEM (electrical biasing

and resistance measuring) experiments. Experiments were performed inside a FEITM Talos

F200X scanning/transmission electron microscope (S/TEM) with a resolution of 1.2Å using field

emission gun and 200 kV acceleration voltage. In our experiment, we pass DC current through

the specimen in steps of 10µA until 0.1 mA limit is reached. We allow about 5 minutes between

each current density steps, which is remarkably shorter than conventional thermal annealing.

Figure 3.2a shows a representative specimen resistance vs. current density data, which shows the

resistance as a function of current density. We noticed a dramatic reduction in initial resistance

from 1 MΩ to about 180 kΩ after the electrical annealing. Electro-thermal simulation results

predict the temperature profile in the specimen as shown in Figure 3.2b. In our present

simulation, we consider low thermal conductivity of MoS2 [158] to ensure maximum

temperature rise during the electrical annealing of MoS2. Literature on Raman thermometry

[159] measurement indicates that MoS2 temperature could reach as high as 380°C prior to their

breakdown at a drain current of 210uA/um. This reported current is more than twice in

magnitude higher compared to our biasing condition. Thus actual temperature rise in our sample

might be even lower than the predicted simulation results. Due to the nature of the freestanding

geometry, the specimen temperature is highest (around 400 C) at the vicinity of the center

region of the sample. Whereas massive silicon electrodes effectively constrain the specimen ends

to be at room temperature as shown in Figure 3.2. The arrows in Figure 3.2b indicates the

locations for the SAED patterns.

66

It is well known that thin film resistance is strongly influenced by surface and internal

(such as grain boundary) defects. For example, Figures 3.3a and 3.3b show the bright-field (BF)

and SAED images for the as-deposited specimen. After annealing at near room temperature with

EWF at 9x105 A/cm2 current density, the corresponding image is shown in Figure 3.3d. The

transformation of pre-annealed sample as shown in Figure 3.3c to apparent featureless

appearance in Figure 3.3d suggests a significant decrease in surface roughness, which may

contribute to the decreased surface scattering. This may have significant impact on the electrical

resistance as we have noticed in Figure 3.2a since the surface scattering of electrons plays a

dominant role [160] on electrical resistivity.

Figure 3.3. In-situ TEM electron wind force annealing results: (a) Bright-field image, (b) SAED

pattern of as-deposited nominally monolayer MoS2 specimen, and (c, d) The same location after

67

annealing at 9.5x105 A/cm2 current density [67].

Due to the high volume fraction of GBs pronounced scattering event takes place at the

GBs. This is in agreement with the SAED available in the literature [149]. Since the grain size

(~5nm) is close to the characteristic length scale of electron mean free path, a small increase in

grain size can result in very large decrease in resistance. Figure 3.3e shows a representative TEM

SAED pattern at low temperature locations (arrows in Figure 3.2b) in the specimens after EWF

annealing. Here, the initially diffused rings appear sharper after annealing. More remarkably, we

observe distinct hexagonal spots at a current density of 9.5x105 A/cm2, suggesting the existence

of large crystalline domains. This observation follows the scaling of EWF, which is the highest

at the grain boundaries. While the discrete spots in the SAED pattern are clear indication of

large-scale crystalline domains, the actual specimen morphology is difficult to achieve. Thus we

propose few possibilities such as (a) prominent secondary grain growth, where energetically

favorable grains may grow abnormally large at the expense of neighboring smaller grains [161]

and (b) electrical discontinuity in domains that lead to little annealing effect. Figure 3.3f

schematically exhibits the later scenario where some regions (labeled as region 2) may see lower

electron wind force as the current will flow mostly from region 1 to 3.

We performed a series of MD simulations to comprehend atomic scale mechanism of

defects annihilation. To investigate the quality enhancement mechanism such as GBs

reconstruction, we analyze simulation trajectories. Figures 3.4a-3.4c show such analysis, which

indicate dislocation migration between adjacent grains. Here, r1 and r2 indicates two rings at the

lower center of the sample to track the initial dislocation position as marked by () at the GB

during different simulation time. Figure 3.4a represents deep yellow colored grain (boundary

marked by black dotted line) and its three nearest neighbor grains. Initial dislocation position

68

adjacent to the ring r2 (as shown in Figure 3.4a) migrates to the next ring as shown in Figure

3.4b after 200ps. With the progression of simulation time dislocation marked in Figure 3.4c

further moves to the GB. This dislocation migration due to effect of EWF and Joule heating

during electrical current induced annealing could lead to the quality enhancement by rearranging

the atoms at the GBs. Figure 3.4d shows the transformation of a 6|8Mo ring defect to 6|6|4

defects due to the synergy of EWF and Joule heating. Thus one 6|8Mo ring defect could generate

two perfect 6 rings and one 4 ring during the electrical annealing as shown in Figures 3.4d and

3.4e.

Figure 3.4. (a)-(c) 6|8Mo type dislocation migration during the electrical annealing, (d)-(e)

Transformation of a 6|8 ring to 6|6|4 ring at the Grain boundary (GB) (Individual grains are

69

shown by different colors, smaller radius sphere indicates Sulphur atoms and larger sphere

indicates Molybdenum atoms) [67].

Figure 3.5 represents the annihilation mechanism of vacancy defects at the GBs of the

sample. Here, we mark r1 and r2 rings to track the defects during the annealing simulation.

Figure 3.5. Transformation of vacancy defects at the GB: (a) initial sample, (b) vacancy

transforms to 6|8 S defects, (c) 6|8 S transforms to 6|6|4 ring defects, (d) formation of 6|8 S due

to the dislocation motion, and (e) formation of 6|6|4 ring (Individual grains are shown by

different colors, smaller radius sphere indicates sulfur atoms and larger sphere indicates

molybdenum atoms) [67].

70

Figure 3.5a shows the vacancy defects at the GBs of the sample prior to the annealing simulation

marked by 0 ps. With the progression of simulation time, we notice vacancy defect transforms to

6|8S ring defects after 265ps as shown in Figure 3.5b. Figure 3.5c shows the intermediate

transition of 6|6|4 defects from 6|8S ring defects after a simulation time of 480 ps. These 6|6|4

defects further transform to 6|8 Mo ring defect after 600 ps as shown in Figure 3.5d. Thus,

dislocation migrates to one ring lower left (Figure 3.5b and 3.5d). After 615 ps of simulation

time, 6|8S ring defects transformed to 6|6|4 defects as shown in Figure 3.5e. This vacancy defect

annihilation at the GBs further facilitates the quality enhancement in the specimen.

3.1.4 Conclusion

In this section, we examined the effectiveness of the EWF induced quality enhancement

and defects mobility in 2D MoS2. Our study indicates that EWF annealing of MoS2 specimens

with defects (such as GBs in polycrystalline MoS2) can be achieved at low temperatures. This is

in stark contrast with the state of the art thermal annealing of 2D MoS2, where very high

temperatures (~800 C) are required. We synthesized monolayer MoS2 at low temperature (~400

C) to test our hypothesis. Then we performed in-situ TEM to observe and study the

enhancement in the quality of the specimen. The electrical resistance decreased by 5 times at a

dc current density of approximately 9.5x105 A/cm2. Such enhanced transport property is

supported by the EWF induced changes in SAED patterns in regions of specimens where the

temperature was constrained to the ambient. To gain insights on defects annihilation, we also

conducted MD simulation. Our simulation study captures the atomic scale defect migration and

grain reconstruction mechanisms that led to specimen quality enhancement. Simulation results

71

shows that during the proposed annealing, the grain boundary atoms could rearrange and re-

reorient due to impetus of EWF.

3.2 Low Temperature Processing of 2D Material based Thin Film Transistors

In this section, we will extend our proposed novel non-thermal annealing process for two-

dimensional (2D) materials based devices. Instead of high temperature, we exploit the electron

wind force (EWF) at near-room temperature conditions to process back-gated WSe2 field effect

transistor (FET). As mentioned in earlier section, EWF is an atomic scale mechanical force that

acts only in the defective regions, which is proposed to provide very high defect mobility. This

EWF processing is demonstrated on back-gated WSe2 transistors. EWF annealing was performed

by passing current through the device drain and source channel while actively removing the

Joule heating. We observed approximately one order of magnitude increase in the output current,

validating our hypothesis on the mobility imparted by the EWF to migrate and eliminate defects.

Our molecular dynamics (MD) simulation confirms the defects annihilation and local metallic

phase transformation, which further enhance device performance. Proposed technique will

potentially lead to time and cost-effective post-processing of two-dimensional materials based

devices.


Two-dimensional transition metal dichalcogenides (2D TMDCs) [162-166] possess

extraordinary electrical, optical and mechanical properties [71, 124, 165, 167, 168] which make

them as a potential candidate materials for next generation nano-electronic applications. Unlike

graphene, single layer TMDCs exhibit indirect to direct band gap transition [169, 170]. This

tunable band gap properties of TMDCs offer new types of electronics such as field effect

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transistor (FET), photo-detectors, and memories with selectable properties [127, 167, 171, 172].

Recently, WSe2 based electronic devices such as FET has attracted significant interest [173-175]

due to their p-type, n-type and ambipolar carrier transport properties via electrode engineering

and thickness modulation [176, 177].

Performances of the 2D electronics are a function of crystallinity of the TMDCs. Depending

upon the synthesis conditions different types of defects such as grain boundaries (GBs), ring

defects, dangling bonds, vacancy defects may appear in 2D materials [178-184]. These defects

potentially degrade the physical properties of 2D TMDCs such as carrier mobility, resistivity,

thermal conductivity and mechanical properties [141, 143-147, 185], which commute through

device performance. A conventional practice to control these defects and enhance device

performance is thermal annealing. However, thermal annealing requires very high temperature

(>800 ºC), which itself can deteriorate the material by creating more defects or stagnating the

existing ones due to the residual thermal stress between materials interfaces. Thermal annealing

is not suitable for low temperature materials as it is a slow process (at low temperature) based on

random diffusion. Instead of conventional thermal annealing [186], an electrothermal-annealing

(ETA) could have potential to enhance device performance after fabrication. ETA method uses

Joule heating as a driving force to eliminates defects and enhance device performance [187].

Recently, ETA techniques have been implemented on metal-oxide semiconductor field-effect

transistor (MOSFET) [187] and amorphous-oxide-semiconductor (AOS) thin-film transistors

[188] . ETA technique uses short-pulses (0.1-1 ms) at high voltages to repair devices. However,

their optimization efforts indicate that annealing time and voltage could be adjusted to achieve

desired outcome.

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In this section, we propose an alternative, non-thermal route compared to conventional

thermal annealing. Instead of using heat energy to increase the atomic mobility, we specifically

target defective atoms with mechanical force know as electron wind force (EWF) by passing

current through the sample. While current passes through the sample, electrons are minimally

scattered by the lattice, whereas they transfer their momentum whenever they meet defects and

GBs [58] thus enhance atomic mobility. EWF is also coupled with resistive (Joule) heating,

which is well known in the electromigration literature [100] Thus, EWF is minimum inside the

lattice and maximum at the defects. Previous studies on blanket films of thin metal [116, 154,

189] [189] and 2D materials [190, 191] confirms the effectiveness of EWF, even in presence of

the Joule heating. To ensure the effectiveness of the low temperature processing we eliminates

the Joule heating using active cooling and yet demonstrate enhancement of crystallinity. Active

cooling decouples and eliminates the Joule heating to maintain the transistor surface at room

temperature. This is in contrast to the literature, where defect mobility is viewed as a

predominantly Joule heating effect. Major contribution of this study is the low temperature

processing of WSe2 FET to enhance the performance of as-fabricated device. Our present study

thus demonstrates EWF annealing on fully functional electronic devices and provides insights on

the processing parameters and their optimization for time and cost-effective performance

enhancement.


3.2.2.1 Metal-organic Chemical Vapor Deposition (MOCVD) Synthesis of Epitaxial WSe2

The MOCVD growth of WSe2 thin films is performed in a custom-built vertical cold-

wall CVD reactor as previously reported elsewhere[192]. The tungsten hexacarbonyl (W(CO)6)

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(99.99 %, Sigma-Aldrich) and hydrogen selenide (H2Se (99.99 %, Matheson)) are used

precursors with high-purity H2 as the carrier gas. The epitaxial WSe2 was synthesized as

previously reported 3 step growth method (nucleation, ripening and lateral growth) at 800 °C and

700 Torr on c-plane sapphire substrates[192, 193]. Prior to the synthesis, substrates were cleaned

using ultrasonication in acetone (10 min) and isopropyl alcohol (10 min) followed by 20 min

immersion in commercial Piranha solution (Nanostrip, KMG Electronic Chemicals) at 90 °C and

deionized (DI) water rinse (10-14 times).

3.2.2.2 Device Fabrication and Characterization

The WSe2-based field effect transistors (FETs) were fabricated and characterized to study

the impact of the EWF annealing process on synthetic two-dimensional materials. The as-grown

monolayer WSe2 film was transferred to a 285 nm thermally grown SiO2 dielectric layer

supported by a rigid Si substrate. Deposited SiO2 serves as the global back-gate of the device.

Standard e-beam lithography (Vistec EBPG 5200) process was used to define the source and

drain electrodes. To isolate the WSe2 channels from blanket film Plasma-Therm Versalock 700

high-density inductively couple (ICP) plasma etch tool was used with a SF6/O2 30/10 sccm gas

mixture for 20 s. We used physical vapor deposited 40 nm Ni/ 30 nm Au metal stack as the

source and drain electrodes. All electrical measurements were carried out on a room temperature-

controlled stage in a high-vacuum (~10-5 Torr) CRX-VF Lake Shore probe station to minimize

threshold shifts due to moisture along with a Keysight B1500A Semiconductor Device Analyzer.

3.2.2.3 Experimental Setup, Modeling and Simulation

To study the effect of EWF we decoupled the Joule heating from the EWF and exploit the

latter to anneal the 2D WSe2 at near room temperature. In our experiments, the transistor was

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placed on a temperature-controlled stage maintained at 296K. Due to the resolution limitation of

commercially available microscope, direct measurement of the temperature is difficult since the

dimension of the channel length and width are 1 μm and 5 μm respectively. Thus to assess

temperature rise in the sample during the electrical annealing, we performed multi-physics

modeling of Joule heating using COMSOL® software as shown in Figures 3.6c and 3.6d. To

mimic the exact experimental condition we consider interfacial thermal boundary conductance of

12 MW/m2.K [194] and 125 MW/m2.K [159] between WSe2-SiO2 and Si-SiO2 respectively. Our

electro-thermal simulation did not show any significant temperature rise up to 30 V annealing

voltage, which could be attributed to the low annealing current density (~3.2×107 A/m2) and

controlled stage temperature (i.e., 296 K). However, at high biasing condition sample

temperature may rise as high as 380 K as shown in Figure 3.6c due to the Joule heating effects at

50 V annealing voltage accompanied by moderately high current density (~3.2×108 A/m2). Our

calculated temperature rise is in a good agreement with recent study [159].

We have adopted molecular dynamics (MD) method to investigate the microstructural

changes during EWF annealing. Though length and time scales are different in MD simulation, it

will provide qualitative understanding of the defects annihilation mechanism and quality

enhancement of the WSe2 film and FET devices. Stillinger-Weber [81] potential is used to

simulate monolayer WSe2 polycrystalline films using LAMMPS [75] code. Polycrystalline WSe2

samples with 16 numbers of grains (Figure 3.9a) and a grain size of 8nm were modeled using

voronoi tessellation technique. The grains were rotated at 0°, 30°, 45° and 60° with respect to

film normal i.e., c-axis. We performed energy minimization using conjugate-gradient (CG)

method followed by isothermal–isobaric (NPT) dynamics runs for several thousand steps to

obtain residual stress-free structure on the as-deposited sample. Once we obtain equilibrated

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structure, we apply wind force on each atom calculated from the Huntington-Grone [93] ballistic

model to mimic electrical annealing (EA) (as mentioned in section 3.1.2).

Figure 3.6. (a) Optical Microscope image of fabricated back-gated WSe2 FET transistor, (b)

schematic diagram of WSe2 transistor, (c) electro-thermal simulation model of WSe2 FET, and

(d) temperature profile across the cross-section of the sample obtained from model [68].


WSe2 FETs devices were annealed at different current density under vacuum condition.

During this annealing electrical annealing induces both Joule heating and EWF. However, due to

the controlled stage temperature Joule heating effects assumed to be insignificant. In addition to

this, transistor substrate itself acts as massive heat sink during the annealing process. The finite

element electro-thermal simulation also suggested negligible temperature (304 K) rise at 30 V

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annealing voltage. Thus, we can hypothesize that single stimuli such as EWF can eliminate

defects even at low temperatures.

Output current as a function of drain voltage (i.e., Id vs Vd) is shown in Figure 3.7a which

reveals the effectiveness of the low temperature EWF annealing. In Figure 3.7a, we noticed an

increment in drain current by more than one order of magnitude after annealing at 20V as shown

by cyan color line. During 20 V EWF annealing we anneal the sample for 5 minutes and the

channel current reaches a maximum value of 0.2 μA. The annealing time is sufficiently longer

compared to the other study [188]. This relatively long duration annealing time makes EWF

method effective without introducing any significant temperature rise in the device.

Enhancement in drain output current during due to the EWF could be attributed to the

improvement of film quality as well as contact resistance at the interface of electrode and the

device layer [186, 190, 195]. We noticed almost one order of magnitude increment in output

current, which suggests significant defect annihilation in the WSe2 channel layer due to the low

temperature annealing. We performed annealing at different voltages namely 20 V, 30 V, 40 V

and 50 V to figure out optimum voltage. We found 30 V as an optimum voltage where annealing

effects is maximum. Above 30V annealing voltage, the device starts to degrade and then gets

damaged at 50 V, probably because the EWF itself starts to migrate atom from the cathode to

anode accompanied by Joule heating. Essentially, the EWF minimizes different type of defects

that can form in 2D materials during synthesis or device fabrication [196]. Due to the high

surface to bulk ratio in 2D materials these atomic scale defects can strongly influence physical

properties. Defects such as tungsten (W) and Selenium (Se) atoms vacancy might introduce

semiconducting to metallic behavior in WSe2 [197] . P-type semiconducting behavior has been

reported in MoS2 with Mo vacancy [198]. Additionally, electronic properties such as electron

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mobility can be significantly limited by atomic scale defects such as vacancy, dislocations, and

grain boundaries [142].

Figure 3.7. (a) Output characteristics of WSe2 FET after annealing at different drain voltage, and

(b) Improvement in drain current after annealing while FET surface was maintained at 296K

[68].

Prior to the annealing, the device shows a maximum of 1.89×10-2 µA/µm drain current

achieved at 45 V gate voltage as shown by transfer characteristic curves in Figure 3.8, while the

drain voltage was kept constant at 5V.We notice significant improvement in drain current after

annealing at 20 V as indicated by pink color line in Figure 3.8a. Drain current reaches a

maximum plateau value of 5.47×10-2 µA/µm around 45 V gate biasing without any threshold

voltage shift (Figure 3.8a). Our optimization efforts shows a maximum drain current of 1.35×10-1

µA/µm (Figure 3.8b) at 30V annealing voltage, which is one order of magnitude higher

compared to the as-received device. This remarkable enhancement of device performance can be

attributed to the improvement of the channel quality. However, device starts degrading above 30

V annealing voltage and failed to operate at or above 50 V. Device degradation could be

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attributed to the high channel current and hot electron (electrons with high kinetic energy) induce

traps in the channel layer [199]. These traps hinder the electron mobility which potentially

reduces the output current of the device at high annealing voltage. In our present study, all

measurements were performed three times and three sets of experiments were conducted to

ensure consistency. As mentioned earlier degradation at higher annealing voltage (i.e., >30 V)

could be attributed to the high current density and traps which are outcomes of combined

electrical and thermal field. Due to this detrimental thermo-migration effect, atomic flux mobility

increases, and defects/voids appears in the channel layer. We can express the atomic flux

mobility of atoms due to the EWF (JEWF) and temperature (JT) as follows [57]:

𝐽𝐸𝑊𝐹 =𝑁𝑒𝑍∗𝐷𝜌

𝑘𝑇𝑗; 𝐽𝑇 = −

𝑁𝐷𝑄𝜌

𝑘𝑇2 ∇𝑇 (3.2a)

Where, D is diffusivity, k is Boltzmann’s constant, T is absolute temperature, N is vacancy

concentration, Z* is effective charge number, j is current density, e is the elementary charge, and

Q is the heat of transport. Eq. (1) shows that atomic flux mobility is proportional to both current

density (j) and temperature gradient (∇𝑇). Thus, at higher annealing voltage high EWF in

conjunction with high thermal field can induce directed diffusion of atoms from cathode to anode

side of the transistor (shown later in Figure 3.10). This directed diffusion can lead to the defects

such as voids and cracks formation near the cathode side, and hillock at the anode side.

Eventually the device fails due to this defects formation and migration effect.

Synthetic 2D TMDCs usually contains higher defect densities compare to the mechanically

exfoliated counterparts [31]. Different types of defects such as grain boundaries (GBs), metal as

well as chalcogen vacancies, and ring defects are commonly found in the synthetic materials,

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Figure 3.8. (a) Transfer characteristics showing WSe2 transistors performance after annealing at

different voltage, and (b) maximum drain current obtained after annealing [68].

which ultimately degrade transport performance of the devices [183, 200-202]. Since defect-free

synthesis of TMDCs remains as challenging, we propose an alternative approach using external

stimuli to improve the film quality and transport performance [116, 154, 189, 190]. Direct

visualization of the defect annihilation or GBs reconstruction is difficult, even with an in-situ

atomic resolution microscope. Therefore, we take resort in MD simulation. To investigate the

effectiveness of EWF annealing we model a nanocrystalline WSe2 film with 16 grains and

average size of 8 nm. The initial equilibrated structure primarily contains 2H phase a shown in

Figure 3.9a), and after annealing it contains 2H, 1T and intermediate phase as well as

reconstructed grain boundaries of WSe2 (Figure 3.9b). At the two opposite ends of the sample,

we define electrode by fixing atoms as shown by the green and red bands in Figure 3.9b. The GB

defects could be annihilated via phase transformation, intermediate structure and twinning as

shown in Figure 3.9b. Similar types of 2H to 1T-phase transformation in monolayer MoS2 has

been observed in recent experiment [203]. Interestingly, our study also shows semiconducting to

metallic behavior of WSe2 during 2H to 1T phase transformation [184]. This local 1T phase

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transformation might enhance the device performance due to the incorporation of metallic

behavior. Colored arrowhead in Figure 3.9b indicate possible electron flow paths without

significant scattering at the GBs. These defects annihilation at GBs might be the primary cause

for the higher output current as shown in Figure 3.7a.

Figure 3.9. Monolayer WSe2: (a) prior to the annealing, and (b) after annealing [68].

Our computational model also captured device degradation phenomenon beyond the

optimum value of the annealing voltage. This phenomenon could be attributed to the thermal and

electrical field induced failure, which is readily observed in metallic thin films [116, 204, 205].

At higher biasing voltage, both EWF and Joule heating are high enough to induce atomic

migration, which further degrades the WSe2 film. In order to investigate void formation and

migration in WSe2 FET, we apply one order of higher EWF in the WSe2 device layer and track

atomic motion as shown in Figure 3.10. Initial sample with anode and cathode marked by red

and blue color region is shown in Figure 3.10a. We observe void formation near the cathode area

of the sample after 5ps. As the simulation time progress, voids start to grow and we observe

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stress relaxation at the vicinity of the voids. With the progression of simulation time, we notice

gradual increment in void size near the cathode area and mass accumulation at the anode area.

Figure 3.10. (a)-(f) Failure at high biasing condition due to the electrical and thermal field, and

(g)-(j) side view of the sample showing void creation at the cathode side (left) and mass

accumulation at the anode side (right) (Color bar in Figure 3.10 shows the atomic stress

distribution in the sample) [68].

Thus, device fails due to this high atomic mobility and thermal field generated by the electrical

current. Figures 3.10g-3.10j represent side view of the sample under high EWF. At the

beginning of the simulation there is no void or mass accumulation (Figure 3.10g) however, with

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the progression of simulation time device degrades due to the voids enlargement (Figure 3.10h

and 3.10i) and mass accumulation in WSe2 sample as shown in Figure 3.10j.

In this section, we have employed electrical current annealing method to enhance the

WSe2 FET device performance at low temperature. We notice one order of magnitude output

current increment after annealing at an optimum annealing voltage (i.e 30 V). However, beyond

this optimum condition device performance starts to degrade, and device fails to operate after 50

V annealing voltage due to the high electro-thermal field. Present study provides insights on an

alternative, time and cost-efficient annealing technique using EWF, which can improve device

performance without introducing thermal stress, and can be applied to a wide range of devices

including flexible electronics.

3.2.4 Conclusion

We have adopted both experimental and computational approach to validate the

effectiveness of EWF annealing at low temperature. We chose WSe2 back-gated FETs, and our

investigation shows more than one order of magnitude drain current output. However,

enhancement of device performance is a non-linear function of annealing voltage, and device

performance degrades after an optimum annealing voltage. Experimental investigation indicates

device exhibits enhanced performance up to 30 V and fails to operate above 50 V annealing

voltage. Our computational study shows GB defects could be annihilated due to the atomic

mobility and phase transformation even at low temperature under the influence of EWF. Study

also shows that metallic 1T-phase transformation might enhance device performance. Simulation

results reveal the device degradation by capturing defects formation and migration, which

includes void creation near the cathode and hillock near the anode. In sum, low temperature

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electrical annealing take less than 5 minutes and could be applied on flexible electronics or other

devices during its operation without introducing any thermal stress. In addition, our study shows

that EWF process is order of magnitude faster and energy efficient compared to the conventional

thermal annealing.

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Chapter 4

Synergy of Stimuli on Operation and Degradation of Nanoscale Devices


Zahabul Islam, Nicholas Glavin, and Aman Haque, Potential and Challenges of in-situ

Microscopy on Electronic Devices and Materials, IOP eBooks, Edited By S. J Pearton, 2020

(submitted)


preparation, experimentation, data analysis and manuscript writing. Nicholas Glavin

guided on manuscript writing. Aman Haque guided on experiment design, and involved

in data analysis as well as manuscript preparation.

Zahabul Islam, Aman Haque, and Nicholas Glavin, Real-time visualization of GaN/AlGaN high

electron mobility transistor failure at off-state, Applied Physics Letters, Volume 113, 183102

(4pp), 2018.


preparation, experimentation, data analysis and manuscript writing. Nicholas Glavin

guided on manuscript writing. Aman Haque guided on experiment design, and involved

in data analysis as well as manuscript preparation.

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4.1 On-state Degradatation of High Electron Mobility Transistor

In this section, we have shifted our focus on thin film devices such as high electron

mobility transistor (HEMT), and studied their degradation behavior under external stimuli. Thin

film HEMT devices are in a plain strain condition due to the specimen geometry and boundary

condition. In addition to plane strain condition, external stimulus such as external biasing

develops thermal, electrical and mechanical stress fields in these HEMT devices. The unique

geometry and complex interaction of multiple domains make these thin film HEMT devices as

ideal candidates for stimuli-synergy study. Thus, in this section we have investigated thin film

HEMT device such as gallium nitride (GaN) based HEMT using in-situ transmission electron

microscopy (in-situ TEM) under external stimuli. Recently, AlGaN/GaN HEMTs have drawn

tremendous attention due to their superior breakdown voltage, energy bandgap, and electron

mobility. However, the realizations of potentials of these devices are impeded by several

reliability concerns. Thus, this chapter addresses reliability issues of these devices under external

stimuli using in-situ transmission electron microscopy (in-situ TEM) techniques.


High electron mobility transistors (HEMT) are a new class of heterostructure field-effect

transistors primarily composed of III-V semiconductor groups. Confined carriers at the

heterojunction of HEMT act as two-dimensional electron gas (2DEG). Both Gallium arsenide

(GaAs), gallium nitride (GaN), and their compounds are widely used as HEMT. Owing to the

higher electron mobility of GaAs HEMT, their narrow bandgap limits their applications in power

electronics. On the other hand, GaN has wide bandgap and high breakdown voltage, which

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makes them as a suitable candidate for high power and high-frequency applications. Figure 4.1

shows a comparison of properties among competing semiconductors [206].

Figure 4.1. Comparison of properties among competing semiconductors.

A comparison as shown in Figure 4.1 reveals that GaN has one order of magnitude higher

breakdown voltage, and both bandgap and electron saturation velocity of GaN is higher

compared to the competitor GaAs. Additionally, thermal conductivity of GaN is also higher

compared to other competing semiconductors. Thus superior material properties make GaN as an

excellent candidate for the next-generation high-frequency, high power applications. Literatures

reported that GaN exhibits more radiation-tolerance compared to SiC-based devices, which also

makes it an attractive choice for space applications [207-211]. Due to these extra-ordinary

properties, the scientific community has seen a surge in GaN-based HEMT applications in the

field of wireless, high voltage electronics, radar, automotive, space, to name a few.

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Owing to their outstanding performance, long-term reliability issues under external

stimuli of GaN HEMT need to be addressed. Conventional approaches of reliability studies

involve electrical characterization data to detect failure signatures under external stimuli.

However, most of the reliability studies are performed outside the microscopes which includes

atomic force microscope (AFM), micro-Raman, cathodoluminescence, infrared,

thermoreflectance, scanning electron microscope (SEM) and transmission electron microscopes

(TEM) study. Thus, most of these studies are post-mortem in nature where failed specimens are

investigated to figure out the failure mechanism. Likewise, ex-situ techniques and post-failure

analysis is strenuous and not accurate, whereas in-situ microscopy could provide both real-time

visualization and quantitative capability to identify failure modes in electronic devices. However,

true potential of in-situ electron microscopes (EMs) study for device degradation under external

stimuli effects remains unrecognized until recently. In recent years, in-situ TEM study has been

employed to investigate GaN HEMT at different biasing conditions [28, 212, 213] to capture the

degradation mechanism in real-time.

It is well known that the degradation of GaN HEMT devices is a function of electrical,

mechanical and thermal fields. GaN HEMT can be operated in four different modes namely, low

current low field, high current low field, low current high field and high current high field.

Among these four operating conditions, both high current low field and high current high field

suffer from high electrical and high thermal fields also known as “on-state” mode, whereas low

current low field and low current high field are known as “off-state” mode and experience only

high electrical field. This chapter will discuss the effects of both “on-” and “off-” state

performance of GaN HEMT using in-situ TEM study. In the following sub-section, we will

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briefly discuss the reliability issues of GaN HEMT under external stimuli at “on-state” modes

using in-situ TEM techniques.


To investigate the stimuli effects on degradation of HEMT deviecs, in our present study

we use a 6 W, 18 GHz, and 40 V rated depletion-mode GaN HEMTs on silicon-carbide

(Wolfspeed, CGHV1J006D) as shown in Figure 4.2. We use this HEMT die (Figure 4.2a) to

prepare electron transparent sample using FIB as shown in Figure 4.3. In-situ TEM experiments

are challenging due to the limited workspace of the TEM chamber, requirement of electron

transparent specimens and their transfer to the testing chip or specimen holder. For example,

preparation of an electron transparent (nominally 100nm thick) specimens involves Focused Ion

Beam (FIB) including series of steps of ion milling and cleaning without introducing significant

damage in the films.

Figure. 4.2. (a) Optical micrograph of GaN HEMT die, and (b) Low magnification SEM image

of the die.

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Specimen preparation for in-situ TEM using FIB requires three important steps: (i) 100

nm thin electron transparent specimen preparation (Figure 4.3a-d), (ii) specimen transfer on

micro-electromechanical system (MEMS) device (Figure 4.3e) [212, 213], (iii) wire bonding of

(MEMS) device [91], and careful placement of chip on a TEM holder with biasing capability

(Figure 4.3f). The specimen preparation using FIB involves deposition of a protective layer to

define a thicker section in the GaN HEMT device and subsequent trench cutting (Figure 4.3a and

4.3b). In our study, a coupon was lifted-off using a controlled manipulator and placed on a TEM

grid (Figure 4.3c). At first, high current such as 21nA was used to lift out the specimen from the

bulk HEMT as shown in Figure 4.3b. However, thinning down to 100nm requires much smaller

current in the range of 0.79nA-34pA (values of currents are based on 4mm working distance).

The thickness of the specimen needs to be monitored at regular intervals during the milling and

Figure 4.3. Details of the GaN HEMT specimen preparation and transfer technique using FIB for

in-situ TEM reliability study.

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cleaning process. In addition, both accelerating voltage and current need to be adjusted

depending on the specimen thickness during the cleaning. After electron transparent specimen

preparation, the final step requires a low accelerating voltage and current to transfer electron

transparent specimen from the TEM grid to the MEMS device (as shown in Figure 4.3e). Low

accelerating voltage was used to accomplish this goal without introducing any significant beam

damage and re-deposition on the specimen. Once the specimen is transferred on the MEMS

device, the specimen is ready for loading on the in-situ TEM holder (Figure 4.3f) and further

external stimuli study was conducted inside a transmission electron microscope (TEM) equipped

with EDS, EELS, SAED, Bright field (BF) and dark field (DF) imaging capability.


Electron microscopes (EMs) have been widely used to study the reliability of electronic

materials and devices due to their analytical as well as atomic resolution imaging capability.

Electron microscopes (EMs) provide high-resolution images that can be deployed to investigate

quality and degradation behavior of electronic materials and devices [214-216]. An EM equipped

with EDS, EELS, SAED, BF and DF makes it an excellent analytical tool for electronic materials

and device characterization. Thus, in this section we will employ in-situ TEM techniques to

investigate the effects of external stimuli i.e., voltage bias on GaN based HEMT under high

current high field state also known as “on-state” condition. During on-state mode operation

external stimuli such as electrical field and current is high enough to induce Joule heating, thus

the thermal field is unavoidable. Due to the high electrical and thermal fields, self-heating, as

well as hot electron effects, plays dominant role on device degradation [217, 218]. Figure 4.4

shows such failure mode obtained from in-situ TEM study. The threshold voltage of our device

is -3.1V, thus we apply 1V gate bias to make sure that the device is at on-state mode. A bright-

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Figure 4.4. In-situ TEM reliability study of GaN HEMT: (a) electron transparent specimen prior

to the application of voltage stimulus, (b) device after failure, and (c) output characteristic curve

obtained from “on-state mode operation”.

field (BF) TEM image of GaN HEMT specimen prior to the application of external stimuli is

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shown in Figure 4.4. Due to the specimen geometry and boundary conditions we noticed

significant amount of bend contour in the device layer, as well as in the substrate layer, (Figure

4.4a). For the purpose of reliability study under external stimuli, electrical connections were

made at the gate, drain and source pad during specimen preparation (Figure 4.3e). FIB deposited

conductive layer was used to accomplish this goal, which further introduced stress in the

specimen. Hence this specimen is in a plane strain condition and bend contour develops in the

specimen due to this residual stress. A failed device after on-state loading, and corresponding

output characteristic (i.e drain current vs drain voltage) is shown in Figure 4.4c. The device fails

at 10.8V drain voltage as shown in Figure 4.4c, and corresponding current density at this

breakdown point is in a good agreement with the reported value [219]. As mentioned earlier, on-

state biasing introduces sufficiently high current density and Joule heating, which further plays a

dominant role in the device degradation (Figure 4.4b). This high current density sublimates

buffer layers (GaN layer) as shown in Figure 4.4b. We have also noticed metal-semiconductor

diffusion near the drain electrode, which indicates thermal field is very high under on-state

biasing condition. Literatures have also reported similar types of failure mode [184, 220, 221]

during on-state biasing.

Figure 4.5a shows observed microstructural changes due to the thermal degradation in the

device layer. High thermal field during on-state operation is confirmed by the amorphization of

the passivation layer (Figure 4.5b) which requires very high temperature as the recrystallization

temperature of SiNx lies within the range of 1200-1400°C [222]. The diffused ring pattern

(Figure 4.5c) obtained from selected area electron diffraction (SAED) clearly indicates

amorphization and polycrystalline nature of the buffer layer. This structural change could be

partially attributed to the very high thermal field accompanied by high current density.

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Figure 4.5. (a) Degradation of the passivation layer, (b) Evaporation of the buffer layer due to the

high thermal field, and (c) SAED indicating the transformation of GaN from crystalline to

amorphous state.

In our study, we have also noticed hot electron induced structural degradation in the

buffer layer as shown in Figure 4.6a. Under high electrical loading conditions, electrons could

achieve high kinetic energy (known as hot electron) which can develop traps and lattice defects

in the buffer layer. As the time progress, these defects could percolate and induce structural

defects as shown in Figure 4.6a, where we notice interface breaching at the GaN/SiC interface.

Weak bonding at the interface of GaN and SiC might make it vulnerable to the hot electron-

induced defects, thus interface easily separates under high external stimuli (Figure 4.6b).

Additionally, high thermal field can evaporate GaN, and small crystallites of gallium

nanoparticle might appear as shown in Figure 4.6c. Due to the lower binding energy of N atom

compared to Ga atom [223, 224], N atom can diffuse out of GaN and facilitates Ga nanoparticle

formation (Figure 4.6c).

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Figure 4.6. GaN HEMT device degradation: (a) Hot electron induced failure at the source side,

(b) evaporation of the buffer layer and formation of small crystallites (nanoparticles) due to the

high thermal field, and (c) high-resolution image of a spherical crystallite.

In our present study, we have also captured failure incidence of GaN HEMT under

external stimuli as shown in Figure 4.7. This uniqueness of this study is to capture failure

initiation and propagation during real-time operation, thus it can pinpoint the exact failure

mechanism. Due to the high electro-thermal fields we have noticed metal pool formation near the

drain area during the on-state operation (Figure 4.7a). However, this metal pool rapidly diffuses

through the GaN buffer layer as shown in Figure 4.7b and 4.7c, and subsequent diffusion and

evaporation continue under high external field as shown in Figure 4.7d-4.7f. At the beginning,

diffusion was initiated in the buffer layer as shown by Figure 4.7g, and a rapid propagation of

diffusion through the substrate layer was observed after 12s as shown in Figure 4.7h and 4.7i. It

is obvious that the high thermal field originating from high drain current is responsible for the

device degradation during on-state operation. During on-state operation due to the minimum gate

leakage, no apparent degradation is observed at the gate (Figure 4.7) electrode. However, hot

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electron ejected from the source during high power operation can breach the GaN/SiC interface

as represented in Figure 4.7i near the source side.

Figure 4.7. In-situ TEM reliability testing showing real-time device degradation.

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To sum up, we have demonstrated a novel in-situ technique for the reliability study of

GaN-based HEMTs under external stimuli, namely drain current. We have observed high

electro-thermal fields induced degradation due to the high electric current. Study also indicates

hot-electron induced failure near the source side. However, no gate degradation was noticed due

to the minimum gate leakage under “on-state” operating condition.

4.1.4 Conclusion

In this present study we have studied complex material system such as HEMT devices

under external stimuli. In-situ TEM studies provide us quantitative as well as qualitative data at

the same time under external stimuli. Our “on-state” failure study shows that high electrical

current could induce high electro-thermal fields including Joule heating. Due to this high

electrical and thermal fields self-heating as well as hot electron effects are dominant and cause

the device failure during “on-state” operation. Though in-situ TEM provides useful information

on the degradation mechanism of HEMT, preparation of an integrated electron transparent GaN

HEMT specimen for in-situ TEM study is challenging and requires special attention. Lower

accelerating voltage and current need to be adopted to avoid any contamination and beam

damage during specimen preparation. During TEM experiment minimization of the e-beam

exposure is also required to enhance the accuracy of the measurements.

4.2 Off-state Failure of High Electron Mobility Transistor

In the previous section, we have studied degradation and failure phenomena in GaN HEMT

devices under external stimuli. It is well known that micro and nano-electronic devices are

complex functions of electrical, thermal and mechanical stresses as well as the quality of the

device materials and their interfaces. Unlike previous section, where we studied devices under

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high external field with high current density, in this section we will demonstrated device

performance under high field but lower current density also known as “off-state” mode inside the

transmission electron microscope (TEM). Thus, in this present study we have studied reliability

of HEMTs under external stimuli using bright-field (BF), dark-field (DF), selected area electron

diffraction (SAED) and energy dispersive spectroscopy (EDS) techniques to characterize the

lattice defects, diffusion of the various elements. The ‘seeing while measuring’ approach

presented in this study can be useful in pinpointing the dominant failure mechanisms and their

fundamental origin under external stimuli.


Recently, scientific community has seen a surge in Gallium nitride (GaN) based high

electron mobility transistors (HEMTs) for high-power and high-frequency applications [225].

Their reliability of HEMTs is influenced by external stimuli such as very high electrical, thermal

and mechanical stress fields during operation [219, 226]. As mentioned in earlier section,

HEMTs can degrade and fail at both ‘on’ (high temperature and electrical fields [227]) and ‘off’

(high electrical field [228, 229]) states. At high power conditions, thermally activated

mechanisms can lead to metal (ohmic contacts, gate and the interconnects) failure [230].

Additionally, high temperatures also gives rise to thermo-mechanical stress, which is added to

the pre-existing residual stress generated during deposition [231]. Due to the piezoelectric nature

of GaN, another source of mechanical stress could be the inverse piezoelectric effect due to the

large values of the electrical field at off-state condition. Again, both high kinetic energy or hot

electrons [230] and diffusion causes [232] electrical and structural degradation of the device.

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The motivation for this study comes from the opportunities for identifying external stimuli

effects on off-state degradation and failure [228]. In our present study, we apply large reverse

bias at the gate to operate the device at off-state mode. During off-state biasing thermally

activated mechanisms are known to be absent or insignificant due to the little or no current

output [233, 234]. Thus, moderate to large-scale defect forms [233, 235, 236] due to the inverse

piezo-electric stress. During off-state biasing condition degradation of AlGaN/GaN HEMTs is a

complex function of the gate bias (with respect to pinch-off), drain bias (with respect to

breakdown voltage) as well as the device geometry and more importantly, the crystallographic

quality of the device layers [237].

Typically, device degradation and failure mechanisms are inferred from the characterization

data obtained from time and stress-controlled tests followed by computational modeling and

microscopy. For example, electroluminescence [229, 232], Raman microscopy [226], and

transmission electron microscope (TEM) has been widely used to visualize the lattice defect and

diffusion induced damages [230, 235, 238, 239]. TEM offers the highest possible spatial

resolution in through-the-thickness imaging with additional features, such as selected area

electron diffraction (SAED), energy dispersive x-ray spectroscopy (EDS) and electron energy

loss spectroscopy (EELS) to identify crystallographic, chemical and electronic states and

damages. In the following sections we will demonstrate external stimuli effects on AlGaN/GaN

HEMTs at off-state using in-situ TEM study [240].


In our present study, we have used a 6W, 18GHz and 40V rated depletion mode GaN

HEMTs on silicon-carbide (Wolfspeed, CGHV1J006D). The specification of the device is as

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follows: gate threshold voltage, saturated drain current, on-resistance and drain-source

breakdown voltages are -3V, 1.1A, 2.3 and 100 V respectively. In our present study, we used

both semiconductor parametric analyzer (Keithley 2400) and in-situ TEM to investigate the

xetrnal stimuli effects on device degradation. In-situ TEM study requires an electron transparent

sample. Figure 4.8a shows a scanning electron microscope (SEM) micrograph of such TEM-

ready specimen. The specimen preparation involves 100nm thin electron transparent sample

Figure 4.8. (a) The experimental setup showing a MEMS chip on a TEM specimen holder, (b)

The specimen integrated with the MEMS chip, (c) SEM image of the electron transparent GaN

HEMT specimen, and (d) transfer characteristic of the HEMT die, and (e) Off-state loading of

the 100 nm thick HEMT sample [241].

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preparation using focused ion beam milling (FIB) and manipulation of this film to a micro-

electro-mechanical system (MEMS) chip equipped with mechanical actuators, heaters and

electrodes. Figure 4.8b shows a typical HEMT specimen integrated with the MEMS chip. The

custom-made chip was then mounted on to a commercially available TEM specimen holder with

electrical biasing capability, which is shown in Figure 4.8c. The in-situ TEM experiments were

performed in a 200 KV FEITM Talos F200X TEM with 1.2 Å resolutions.


Transfer characteristics as shown in Figure 4.8d suggests gate threshold voltage (Vth) is -

3.1V for depletion mode HEMTs. Thus to study off-state response, we kept the gate voltage

(VGS) below -3.1V i.e., at -5V. In this study, we performed the experiments in two phases, which

includes die-scale as well as 100 nm thin film of HEMTs characterization. These outcomes are

represented in Figures 4.9a and 4.9b respectively. A comparison between the die-level data and

Figure 4.9. Off-state characterization of (a) die-level transistor, and (b) 100 nm thick HEMT

during phase I loading [241].

electron transparent specimen data are shown in Figure 4.9. The higher drain current output of

100 nm thick specimens indicates higher leakage in the electron transparent sample compared to

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the die. This can be realized by the difference in the boundary conditions and the higher surface

to ratio of thin films compared to the die. Surface defects in GaN HEMTs can contribute to

higher leakage current compared to the bulk. Due to the nature of our thin film specimen

geometry surface defects are unavoidable, and is expected to further increase the leakage.

Phase II experiments involve the loading of electron transparent HEMT specimens inside

the TEM until it fails under high negative stimuli applied at the gate electrode. Outcome of this

condition is shown in Figure 4.8e. In our study, the electron transparent HEMT specimens

survived up to 30V drain bias prior to the catastrophic failure. This value is in well agreement

with the 100 V breakdown rating of the die with gate width of about 200 μm.

Both electrical field and the inverse piezo-electric stress are prominent at the drain side of

the gate edge. Due to the presence of gate fieldplate and small reverse bias applied at the gate,

we anticipated degradation and failure mechanisms that are not common in the literature, such as

the source injection and the surface hopping conduction. Figure 4.10 shows source-injected

degradation that leads to the failure in the source-gate region. As shown in Figure 4.10a the

virgin specimen with pre-existing defects. Figure 4.10a identifies two hubs for defect nucleation

as marked with an oval and a rectangle at the drain and source side respectively. Figures 4.10b-

4.10d show the specimen during the various phases of loading. With the increment of drain bias,

we notice an appreciable amount of changes in the dislocations network and structure

(comparing the same region such as oval or rectangle area). Figure 4.10e shows the output

characteristics, with the locations of each of the TEM snapshots. These changes indicate that

with TEM it is possible to capture defect nucleation and propagation events under external

stimuli loading.

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Figure 4.10. Bright field TEM images acquired at drain voltages: (a) 0V, (b) 7.2V (c) 11.6V and

(d) 23V, and (e) corresponding drain current vs. drain voltage data at gate voltage -5V for a 100

nm thick GaN HEMT specimen [241].

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Source injection dominated failure has been shown in Figure 4.10d. Interestingly, the drain

side showed lattice degradation as well, whereas region underneath the gate remains mostly

intact. This phenomenon is attributed to the competing leakage mechanisms such as source-gate,

gate-drain and source-drain, which are influenced by the test (biasing) protocol as well as device

dimensions and quality of the materials and interfaces [242] Additionally, incorporation of a

field plate reduces the electrical field strength under the gate. Test conditions under high external

stimuli (in this case gate voltage below threshold value), show breakdown primarily from gate-

drain leakage. This is not the case represented in Figure 4.10 since VGS and Vth are about -5V

and -3V respectively. Thus, source injection governs the failure mode as confirmed by Figure

4.10e. This is also known as premature breakdown in the literature [242].

The true potential of in-situ TEM technique can be realized by predicting the failure

behavior for a given state of pre-existing defects as demonstrated by a separate set of experiment

in Figure 4.11. We notice dislocations network right under the source electrode marked in Figure

4.11a with arrow 1. Additionally, two more sets of potential defect structures were present

between the gate and the drain (marked with arrows 2 and 3). Due to this pre-existing defects we

assumed a punch-through [228] failure mode during off-state biasing, which actually took place

at premature drain voltage of 3V (the gate voltage fixed at -5V). Figure 4.11b manifests the

HEMT specimen right after the punch-through, where the leakage is also shown to activate the

dislocations in the locations marked by arrows 2 and 3, which has further moved under the gate.

This activated dislocation motion increases the source-drain leakage, and a high current density

The electron transparent thin GaN HEMT specimen yields very high current density in the

sample, which can melt the metal contact at the drain and subsequently damages the interface

with the AlGaN. Figure 4.11d shows the thin HEMT specimen at current density of 2000

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mA/mm. The drain electrode forms a pool marked by arrow t1, which takes about 6.2 seconds to

diffuse completely through the GaN layer. This state is shown with arrow t2 in Figure 4.11b.

Figure 4.11. Real-time operation of GaN HEMT: (a) electron transparent HEMT specimen prior

to the loading, where arrows indicating pre-existing defects, (b) magnified view of the arrows

marked 2 and 3 at the on-set of source-drain leakage, (c) molten drain metal pool at this instant

the drain side at current density of 2000 mA/mm, (d) metal diffusion through the GaN layer, (e)

rapid breaching of the GaN-SiC interface, and (f) degradation of the SiC layer [241].

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Since an interface offers the weakest resistance to diffusive transport of atoms, thus molten metal

rapidly penetrates the GaN-SiC interface. Figure 4.11e shows the moment when the interface is

breached, and the near-amorphous metal starts diffusing through the substrate SiC layer. The

substrate is also completely degraded by diffusion, albeit at a slower rate than the rate observed

for the previous phenomena.

Chemical elemental mapping capability of the TEM can provide us useful information on

Figure 4.12. EDS mapping of a failed HEMT specimen at the (a-d) gate and (e-h) drain areas. (i,

j) normalized weight percentage of the various elements in the gate and drain area respectively

[241].

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multi-element diffusion processes. Energy dispersive (EDX) spectroscopy mapping is shown in

Figure 4.12 for both gate and drain regions. At the gate region (Figures 4.12a-4.12d), we noticed

diffusion of gallium and nitrogen through the gate. The gate region also shows oxygen, most

likely trapped during specimen handling in the ambient. Since the current density is lower during

off-state operation, temperature is also expected to be lower, which implies that the inverse

piezo-electric stress can enhance diffusion. Interestingly, we did not notice nitrogen diffusion at

the drain side (Figures 4.12e-4.12h). Figures 4.12j highlighted the compositional changes at the

gate and drain area respectively. This analysis will enable us to study the feasibility of this

technique to map the diffusion and predict the degradation in the HEMTs.

4.2.4 Conclusion

We summarize that degradation and failure mechanism in GaN-based HEMTs (and

microelectronic devices in general) devices are complicated phenomena. Our off-state failure

experiments demonstrate the feasibility of monitoring dimensional and microstructural aspects of

electronic devices and connect that to their performance and reliability. However, experimental

methods need to be improved for minimum contamination from the specimen preparation and

handling. Additionally, modified boundary conditions of an electron transparent thin specimen

require scaling to the standard device geometry through modeling and experimentation. Finally,

care must be taken to avoid the electron beam damage during the operation of the electronic

devices. These tasks are challenging, however the potential benefits of visualization of

operational and failure phenomena in real time can provide insights on device degradation

mechanism under external stimuli.

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Chapter 5

Ion Irradiation and External Stimuli Effect at Nanoscale


Zahabul Islam, Angela L. Paoletta, Anthony M. Monterrosa, Jennifer D. Schuler, Timothy J.

Rupert, Khalid Hattar, Nicholas Glavin, Aman Haque, Heavy ion irradiation effects on

GaN/AlGaN high electron mobility transistor failure at off-state, Microelectronics

Reliability,Volume 102, Pages 113493 (9pp), 2019.


device fabrication, experimentation, data analysis and manuscript writing. Angela L.

Paoletta performed the die-level transistor characterization accompanied by Zahabul

Islam. Anthony M. Monterrosa, Jennifer D. Schuler and Khalid Hattar conducted

irradiation experiment at Sandia national Laboratory and involved in manuscript writing.

Timothy J. Rupert guided on manuscript writing. Aman Haque guided on experiment

design, and involved in data analysis as well as manuscript preparation.

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5.1 Irradiation Damage and Degradation in Nanoscale Transistor

In this chapter, we will discuss the ion irradiation effects on AlGaN/GaN high electron

mobility electron transistors (HEMTs) using in-situ transmission electron microscopy (in-situ

TEM). The experiments were performed inside an electron microscope (EM) to visualize the

defects, microstructure and interfaces of ion irradiated HEMTs during operation. Experimental

results exhibits heavy Au4+ ion induced different types of defects such as vacancies, interstitials

and dislocations in the device layer. It is hypothesized that these defects act as carrier traps in the

device layer and the resulting charge accumulation lowers the breakdown voltage under external

stimuli. Sequential energy dispersive X-ray spectroscopy (EDS) mapping allows us to track

individual chemical elements during the experiment, which further suggests that electrical

degradation in the device layer may originate from oxygen and nitrogen vacancies.


GaN based high-electron mobility transistors (HEMTs) are potential candidates for next

generation electronics such as power amplifiers, broadband communication and high-voltage

switches due to their high breakdown voltage and wide-bandgap [243, 244]. The high carrier

density and mobility of the two-dimensional (2D) electron gas (2DEG) channel yield low on-

resistance which further reduces switching losses when operated in switching mode power

converter [245]. In addition, the wide bandgap of GaN allows operation under high electrical and

temperature fields [246] which makes them an attractive choice for harsh environment

applications [247, 248].

Due to their size, weight and power effectiveness, GaN HEMTs are attractive for space

applications. However, high energy particles featuring the cosmic rays or solar flares in space,

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with energies up to 100 MeV for the protons, and up to 10 GeV for heavy ions [249] can create

electron-hole pairs and displace atoms from their original lattice position, leaving vacancies,

interstitials and dislocations in the crystal [250]. Over time, these defects further accumulate and

interact with each other to generate stacking faults, dislocation loops, and vacancy/interstitial

clusters [251, 252]. Such microstructural degradation degrades device performance [253].

Literatures also suggest lower carrier density and mobility in the 2DEG sheet [254] and a

decrease of the Schottky barrier height at the gate [255].

Ionization effects in GaN are not severe due to the absence of gate dielectric in the

HEMT structure and the higher surface state density in GaN [256], which makes them superior

candidate for radiation tolerant electronics compared to silicon. Measurement of transport

properties of proton-irradiated GaN thin films and devices corroborates the radiation tolerant

properties of GaN compared to GaAs [247, 248, 257, 258].

However, literatures also report the evidence of a decrease in DC saturation current and

transconductance at a fluence of 1014 𝑐𝑚−2 protons or even at lower dose such as 1012 𝑐𝑚−2

[256, 259]. Thus exact mechanism(s) for irradiation-induced degradation and their signatures in

the transistor characteristics demand further investigation. The current trend in the literature is to

measure electrical properties using ex-situ techniques [260]. However, pinpointing the exact

mechanism using ex-situ techniques is strenuous, where in-situ microscopy can be of tremendous

help. Ex-situ studies are post-mortem in nature and can indicate the extent of radiation damage,

but it cannot capture their interaction with the defects generated due to the high electrical,

mechanical and thermal fields during HEMTs operation. This situation can be expected to be

exacerbated with the presence of radiation induced defects. Thus, we suggest that in-situ

microscopy could be helpful in resolving these concerns.

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In the folowing section, we will demonstrate in-situ electron microscopy such as in-situ

transmission electron microscopy (in-situ TEM) technique for simultaneous quantitative and

qualitative Study of GaN HEMTs under external stimuli. In this study, we havecharaterized both

pristine and Au4+ irradiated GaN HEMTs die and electron transparent coupons using a

semiconductor parametric analyzer and TEM respectively. Our in-situ study not only provides

invaluable information on heavy ion induced defects geneartion in GaN HEMTs, but it also

predicts the performance and degradation mechanism of the device under heavy ion irradiation in

a harsh environment. Due to the nature of in-situ TEM study it allows us simultaneous

quantitative characterization and qualitative visualization during real time operation of HEMTs.


To study the effect external stimuli response on ion irradiated HEMTs’ electrical

performance and failure, we designed a series of experiments described as follows:

(a)At first, we characterized commercially available depletion mode GaN HEMTs

(Wolfspeed, CGHV1J006D rated at 6W, 18GHz and 40V) at die level using a Keithley 2400

semiconductor parametric analyzer at room temperature.

(b) In the next step, pristine HEMT dies were irradiated normal to the surface with 1.5

MeV Au+ ions using the 6 MV HVE Tandem accelerator at Sandia National Laboratories. The

ex-situ ion irradiation was performed on three dies with fluences approximately 6.5 x 1013, 6.5 x

1014, and 6.5 x 1015 ions/cm2. The ion energy was chosen based on a Stopping and Range of Ions

in Matter (SRIM) simulation [261] to have a relatively uniform damage profile in the device

layers. Irradiation damage level in term of dpa were 0.45, 4.5, and 45.0 dpa for each respective

112

fluence level. The irradiated dies are then electrically characterized using a Keithley 2400

semiconductor parametric analyzer to compare the die-level performance degradation.

(c) After characterization at die-level we prepared electron transparent (~100 nm thick)

HEMT specimens using focused ion beam (FIB) as described in earlier section [212, 240]. These

thin electron tarnsparent HEMT specimens were then in-situ ion-irradiated inside a TEM [262]

as mentioned in step (b) to obtain visual description of radiation induced defect generation.

However, in our present study we consider sample with highest dose for in-situ TEM study. Both

pre-irradiated and post-irradiated electron transparent HEMT specimens were characterized for

DC transfer and output characteristics inside a TEM. The specimens are electrically biased untill

they degrades.

During electron transparent specimens preparation, special care was taken to minimize

the FIB damage and to keep the three electrodes (i.e drain, gate and source) intact. The transfer

process of thin specimen from a custom TEM grid to a micro-electro-mechanical (MEMS)

device also requires low accelerating voltage i.e., 5kV and ion beam current exposure to avoid

any damage in the device layer. Figure 5.1b shows a mounted chip on a TEM specimen holder

with electrical biasing capability. Details of the sample preparation and transfer technique has

been described in the earlier section i.e., 4.1.2. Figure 5.1c shows the electron transparent thin

sample on the TEM grid prior to the mounting on the MEMS chip. We performed in-situ TEM

experiments in a 200 kV FEI Talos F200X S/TEM with 0.12nm resolution.

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Figure 5.1. Experimental setup for in-situ TEM experiment of electron transparent HEMTs: (a)

GaN HEMT die, (b) a MEMS chip with the HEMT specimen mounted on in-situ TEM electrical

biasing holder, and (c) FIB lamella of the HEMT before mounting on to the MEMS chip [28].

High resolution transmission electron microscopy (HRTEM) allows us to estimate the

atomic-scale strain in the sample using geometric phase analysis technique (GPA) [263]. In the

GPA technique, phase image, 𝑃𝑔 can be expressed by the component of the displacement field,

𝒖(𝒓), in the direction of the reciprocal lattice vector g as described below:

𝑃𝑔(𝒓) = −2𝜋𝒈𝒖(𝒓) (5.1a)

Thus, a two-dimensional displacement field can be obtained using Eq. (5.1a) by choosing

two independent phase images (Pg1 and Pg2):

𝒖(𝑟) = −1

2𝜋[𝑃𝒈𝟏(𝒓)𝒂𝟏 + 𝑃𝒈𝟐(𝒓)𝒂𝟐] (5.1b)

(𝑢𝑥𝑢𝑦

) = −1

2𝜋(

𝑔1𝑥 𝑔1𝑦

𝑔2𝑥 𝑔2𝑥)

−1

(𝑃𝑔1

𝑃𝑔2) (5.1c)

Where 𝑃𝒈𝟏 and 𝑃𝒈𝟐 are two phase images, 𝒂𝟏and 𝒂𝟐 are lattice vectors in real space, 𝑢𝑥

and 𝑢𝑦 are displacement fields, 𝑔𝑥 and 𝑔𝑦 are two components of vector g in reciprocal space.

114

Once displacement field is known, atomic strain maps were obtained from elasticity theory using

the following equations:

𝜀𝑥𝑥 =𝜕𝑢𝑥

𝜕𝑥 (5.1d)

𝜀𝑥𝑦 =1

2(

𝜕𝑢𝑥

𝜕𝑦+

𝜕𝑢𝑦

𝜕𝑥) (5.1e)

𝜀𝑦𝑦 =𝜕𝑢𝑦

𝜕𝑦 (5.1f)

Where 𝜀𝑥𝑥 , 𝜀𝑥𝑦 and 𝜀𝑦𝑦 are normal strain in x direction, shear strain in x-y plane and normal

strain in y direction respectively.


5.1.3.1 Die-level Irradiation Effects: Figure 5.2a shows the schematic cross-sectional view of

the die and irradiation direction (downward arrowhead), whereas Figure 5.2b shows the SRIM

simulation results on the damage level in displacement per atom (i.e., dpa) units as function of

depth into the device. Bright field (BF) TEM images of a pristine and one ion irradiated HEMT

at 45 dpa as shown by Figure 5.2c and 5.2d respectively, indicate ion irradiation effect in the

device layer. Figures. 5.2c and 5.2d exhibit dislocation density comparison between 45dpa

irradiated sample and pristine counterparts to estimate the relative difference in dislocation

density. We noticed significantly higher (more than 100x) dislocation density in the irradiated

HEMTs compared to the pristine one, which is related with the degraded performance measured

by the parametric analyzer.

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In our experiment, we first measured the performance of pristine and irradiated HEMTs

in their bulk form. Figure 5.3 shows both transfer and output characteristic curves of pre- and

post-irradiated GaN HEMT die. Transfer characteristics curve suggested that threshold voltage

(Vth) for a pristine HEMT is approximately -3.1V. However, for all irradiated HEMTs we do not

see any significant increase in output current compared to the pristine counterpart during the gate

voltage increment. This suggested that ion irradiation significantly damage 2DEG channel. In

our present study, we kept the gate voltage (Vg) at -5 V to ensure off-state condition. Figure 5.3b

shows output characteristics curve for all irradiated dies at off-state operating condition. Here,

Figure 5.2. (a) Downward arrowhead in the schematic diagram of the GaN HEMT showing

irradiation direction, (b) displacement per atom (dpa) profile as a function of depth for different

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doses of irradiation, (c) TEM image of a pristine HEMT showing mostly bend contours, and (d)

TEM image of an irradiated HEMT at 45 dpa showing very high dislocation density [28].

four orders of magnitude reduction in output current (for the highest damage level i.e., 45 dpa)

indicates ion irradiation can create significant amount of defects/structural damage and thus

degrade device performance. This drastic reduction in output current for all irradiated devices is

a clear indication of defect introduction in the AlGaN layer as well as the GaN layer during the

Au4+ ion irradiation, which lead to lower carrier density and mobility in the 2DEG channel.

Figure 5.3. Die-level HEMTs specimens characterization curve as function of ion irradiation

damage in dpa: (a) transfer characteristics, and (b) output characteristics [28].

5.1.3.2 In-situ Ion Irradiation Effects (No Biasing): The objective of this set of experiments

was primarily to monitor and identify the irradiation-induced defects in the lattice and interfaces

of the HEMT system. Formation of defects due to the ion irradiation influences the device

performance, since the HEMTs devices experience a very large amount of thermal, electrical and

mechanical stresses during operation. Figure 5.4 shows TEM micrographs of a specimen before

and after exposure to 2.8 MeV Au4+ ion species for 60 minutes to a fluence of 4x1014 ions/cm2.

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Micrograph shows shifts in sample contrast, particularly at the locations of pre-existing

dislocations marked by dotted rectangular region. The effect of ion bombardment is pronounced

right under the gate, where the mechanical stress is also the highest. During an irradiation

process, interstitials atoms might leave their regular lattice sites and occupy interstitials sites.

Figure 5.4. Electron transparent HEMT device: (a) before, and (b) after 2.8 MeV Au4+ ion

irradiation for 60 minutes to a fluence of at 4x1014 ions/cm2. The rectangular dashed box shows

contrast change due to point defect accumulation, while the arrows indicate dislocation activities

at the GaN-SiC interface [28].

These interstitials atoms might appear as dark contrast in a BF TEM image [34]. Thus, the

contrast change (dashed rectangular box in Figure 5.4b) is an outcome of increased number of

interstitials defects. We also noticed remarkable irradiation effects at the GaN-SiC interface.

Figure 5.4b shows how the initially sharp interface is deteriorated due to the introduction of

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dislocations (indicated by arrows). However, these dislocations might be an outcome of the

radiation damage or resulted as a strain relief mechanism due to interface stress from the

radiation damage.

5.1.3.3 In-situ TEM Electrical Characterization: In addition to die-level characterization, we

also conduct experiments on the electron transparent specimens inside the TEM. We prepared

the specimens from the ex-situ irradiated dies because of the well-characterized nature of the

irradiation simulation and boundary conditions. The objective is to visualize the defect and

damage evolution processes under external stimulus such as electrical biasing. Major advantages

of in-situ analytical TEM study is simultaneous quantitative (device characterization) as well as

qualitative (microscopic visualization) characterization during real time operation. Electron

transparent specimens have different aspect ratio and boundary conditions compared to the die-

level specimens [241] thus their transfer and output characteristics are not numerically identical

to HEMT die. However, their characteristics functions follow similar trend, which suggests that

our in-situ TEM experiments can be useful for capturing the mechanics and physics of

degradation under external stimuli.

Figure 5.5a shows output characteristics of an irradiated sample during failure test. Data

shown here represents the transistor irradiated with highest damage (45 dpa). We observed that

drain current increases after each loading step up to a drain bias of 8V. However, after 8V drain

bias, the current shows large fluctuations prior to the failure at 10.2V bias. We also labeled 5

distinct data points (a, c, d, e, f) on Figure 5.5a, for which the corresponding bright-field (BF)

TEM images are presented in Figure 5.6. To identify the influence of irradiation on device

performance we have compared a pristine and an irradiated HEMT sample as shown in Figure

5.5b. We noticed more than two times higher operating drain voltage, and one order of

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magnitude higher output current for pristine sample compared to the irradiated sample. This

degradation in device performance after irradiation is attributed to the defects formation during

an irradiation process [259, 260].

Figure 5.5. (a) Drain current vs. drain voltage plot of electron transparent GaN HEMT specimen

during off-state operation inside the TEM. The data labels correspond to the in-situ TEM images

in Figure 5.6., and (b) Comparison between pristine and irradiated conditions [28].

Figure 5.6 represents BF TEM images at different loading condition during the off-state

failure tests. These low magnification images (except Figure 5.6b) show numerous bend contours

as well as a number of defect clusters. Examples of these heavy Au4+ ion irradiation induced

defects such as dislocation is captured at very high magnification and is shown in Figure 5.6b.

Atomic resolution imaging enables us to observe individual dislocation indicated by red colored

perpendicular sign () in the sample. However, we did not notice any significant amount of

defects in pristine dies. Figure 5.6c shows the HEMT specimen at a drain voltage of 5V, which is

large enough to trigger very minute changes in the drain current (as shown by point c in Figure

5.5a). This drain voltage i.e., 5V eventually accelerates microstructural damage nucleating from

preexisting defects. For irradiated specimens, irradiation damage can trigger such nucleation.

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Figure 5.6 also shows bend contours that arise from elastic bending during specimen preparation

(due to the mechanical constrains at the drain, gate and source electrode).

Figure 5.6. TEM BF images showing source, gate and drain at the same time. Drain voltage: (a)

Vd= 0V, (b) TEM image of screw dislocations in buffer layer, (c) Vd= 5V, (d) Vd= 7V, (e) Vd=

8.5V, and (f) Vd= 10.2V [28].

Low magnification micrographs as shown in Figures 5.6c-5.6f capture the microstructure

at the drain, gate and source at different drain bias as indicated on Figure 5.5a. Three arrows as

shown in Figure 5.6c (labeled 2, 3 and 4) identify the location of defect clusters that appear to

reproduce dislocations with current stressing. As we increased the drain bias, significant

microstructural changes are observed in the GaN layer. To identify damages those extended from

the arrows 2 and 4, we marked two dashed rectangular areas are shown in Figure 5.6e. Later on,

we noticed interfacial breaching in between these two locations. Figure 5.6f shows micrograph

of the sample just after the failure at 10.2V drain bias. Lower breakdown drain bias of the

irradiated sample could be attributed to defects generated at the interface of AlGaN layer as well

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as GaN buffer layer during the ion irradiation. This scenario is exacerbated by the pre-existing

irradiation induced defects, and inverse-piezoelectric stress induced dislocations during the off-

state loading [233, 236]. These studies also indicate that the resulting mechanical stresses create

defects induced traps in the GaN layer through local charge accumulations. The contribution of

our study is therefore provide visual evidence of the lattice defects generation (Figure 5.2d and

5.6b) and interfacial (Figure 5.4b) defects, which lower the breakdown strength [264] and

facilitates failure through their percolation [265].

High-resolution in-situ TEM imaging of defects allows us to track individual defect and

thermo-mechanical strain in the sample during the real operation (as shown in Figure 5.7). For

example, Figure 5.7a exhibits breakdown of the buffer layer and generated dislocations at the

drain-gate channel. Figure 5.7b and 5.7c represent BF low magnification and high-resolution

TEM image of dislocations in the buffer layer respectively. These defects could act as surface

traps in the buffer layer, and percolation of these defects possibly lowers the breakdown voltage

due to the stress field around individual dislocation. To validate stress field associated with a

dislocation, in our present study we quantified atomic strain in the sample by employing GPA

technique as shown in Figure 5.7c. The atomic scale strain mapping process involves several

steps as follows: (a) Fast Fourier Transformation (FFT) of HRTEM image, (b) selection of two

diffraction spots along lattice direction, (c) reconstruction of image using inverse FFT (IFFT),

and (d) phase image, displacement field and strain field calculation using eqns. (5.1a), (5.1c) and

(5.1d-5.1f). Calculated normal strain, shear strain and corresponding lattice fringes of the sample

with dislocations (as marked by white upward arrow) are shown in Figures 5.7d-f. Dislocations

in the sample are discernable even by visual inspection as shown in Figure 5.7c and 5.7f. On the

other hand, atomic strain mapping as shown in Figure 5.7d and 5.7e allow us to identify

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individual dislocation in the sample and associated strain field. In our present study, we notice

basal plane slip (i.e., [0001] plane) associated with [1120] prismatic dislocation. At the location

of dislocations, we notice the existence of both tensile and compressive strain field as shown in

Figure 5.7d and 5.7e, which further introduces stress field around a dislocation.

Figure 5.7. (a) TEM BF image at the drain side of drain-gate region, (b) Dislocations in the GaN

layer, (c) High-resolution TEM (HRTEM) image of dislocations in GaN layer, (d, e) atomic

strain mapping in the sample showing normal and shear strain field associated with individual

dislocations, and (f) simulated lattice fringes with the dislocations [28].

We also monitored real-time diffusion of chemical elements during the experiments using

EDS. Figure 5.8 shows such an experiment to capture chemical elements diffusion at different

drain bias. Figure 5.8 indicates nitrogen deficiency under the gate electrode. Literatures also

suggests that nitrogen vacancies are the primary defect forms in GaN [266] under ion irradiation

123

due to the lower displacement energy of nitrogen (N) atom compared to gallium (Ga) [267]

atom. This defect could act as a trap and reduce the conductivity in GaN [268], which in turn

reduces the carrier density and, increases carrier scattering in the 2DEG channel at the interface

of AlGaN/GaN. Thus, accelerated breakdown of an irradiated device could be attributed to these

nitrogen vacancy defects. The breakdown zone in the sample indicates nitrogen concentration

deficiency in GaN layer (as shown in Figure 5.8b). Figure 5.8c shows the quantitative data on

relative changes in the weight percentage of other elements in the region between the gate and

the drain (the dotted rectangular box in Figure 5.8a). We notice reduction of both Ga and N after

breakdown, which could be attributed to the diffusion of GaN buffer layer out of rectangular box

and possibly into the SiC substrate layer. The reduction of nitrogen is more prominent than

gallium. On contrary to this, weight percentage of gold increases after failure, which could be

attributed to the diffusion from the gate electrode to the buffer layer.

Figure 5.8. EDS mapping of GaN HEMT showing diffusion of chemical elements at different

drain bias: (a) Vd=0V, (b) onset of failure at Vd=10.2V, and (c) relative changes in diffusion of

chemical elements obtained from EDS at these two voltages [28].

124

To summarize, we demonstrated an in-situ TEM technique for visualizing nucleation of

ion irradiation induced defects as well as their proliferation during electrical stressing of GaN-

based HEMTs. Ion irradiation generates significant number of dislocations that contributed to

gate leakage, which further led to the failure.

5.1.4 Conclusion

In this study, we have investigated the effects of heavy ion irradiation on defects evolution

and performance of GaN HEMTs under external stimuli. Heavy ion (such as Au+) irradiation

generates lattice defects such as vacancies and dislocations, which degrade device performance

and accelerate permanent damage. We observed increased defect concentrations in the device

layer with the increment of irradiation dose. Study shows that defects could act as traps thus

degrade the carrier density and mobility in the 2DEG channel. Device degradation could be

attributed to the gate injection induced impact-ionization in the channel, which in turn dictates

the off-state failure. Oxide formation near the breakdown region of the channel layer suggests

that chemical oxidation also accelerate/partially contribute to the electrical degradation.

Continuation of this study in future will include (a) scaling the physics of degradation from

electron transparent device to die-level counterparts, and (b) investigation of the synergistic

relationship between defects and diffusion which leads to degradation and failure. On-state

device reliability is also required to comprehend fundamentals of the ion irradiation damage in

wide bandgap electronics.

5.2 Low Temperature Stimuli Synergy on Recovery of Irradiation Damage in Thin Films

In this section, we will explore the effectiveness of external stimuli such as electrical current

induced electron wind force (EWF) in annihilating defects originating from ion irradiation. In

125

this present study, self-ion irradiation to a dose of 5 x 1015 ions/cm2 (45 displacement per atom

(dpa)) was used to generate high density of displacement damage in a nanocrystalline gold

sample. Afterwards, an electron transparent specimen was prepared form bulk irradiated gold

sample to investigate external stimuli effects on ion irradiation defects using in-situ transmission

electron microscope (in-situ TEM). Unique geometry and boundary condition of electron

transparent specimen effectively decoupled Joule heating from the electron wind force (EWF),

thus allowing us to study the EWF effects on defects annihilation. Our study shows that EWF

can impart significant defect mobility even at low temperature, resulting in the migration and

elimination of defects in a few minutes. We propose that (EWF) interacts with defects to create

highly glissile Shockley partial dislocations, which makes the process fast and energy efficient.

This effect is even more pronounced for nanocrystalline materials with large fraction of grain

boundaries (GBs) and surface area, which subsequently act as active sinks for the migrating

defects.


Irradiation is a process where highly energetic particles (ions, neutrons, electrons) collide

with atoms in materials, energizing them to be displaced from their original lattice positions,

thereby generating various types of defects [269]. For example, point defects, such as vacancies

and interstitials are prominent and they further agglomerate, migrate and form defects clusters

amorphous zone, dislocations loops, three-dimensional defects, ripples, adatoms, craters, to name

a few [270, 271]. These defects can significantly degrade materials properties due to the

embrittlement, swelling and hardening [272, 273]. The literatures have primarily focused on

radiation effects on microstructure and properties of materials using both experimental and

computational techniques [272, 274, 275]. This has led to foundational understanding of the

126

evolution and mechanics of defects. However, processes for defect annihilation are not studied,

other than high temperature annealing [276, 277].

In this study, we propose electron wind force (EWF) as a non-thermal route to actively

control defects and microstructure in an irradiated nanocrystalline materials. During EWF

annealing a highly localized mechanical force imparted to defects due to the collision and

momentum transferred to the defective atoms by the electrons [278]. EWF is strongly coupled

with Joule heating when electrical current is passed through metals [279]. Electro-thermal effects

have been demonstrated to control defects and microstructures [48, 280, 281] in metallic

materials. In our present study, our sample geometry allows us to decouple thermal contribution

from EWF. We achieve this by making a thin film specimen freestanding, so that the regions

near the edges are constrained to lower temperatures due to the missive silicon heatsinks attached

to it. Our hypothesis is that the high specificity of defects can make EWF a fast and low

temperature driver for defect annihilation. Nanocrystalline materials are attractive from this

perspective since they contain large volume fraction of grain boundaries (GBs), phase

boundaries and free surfaces that act as defects sinks [269, 271, 282].


In our present study, polycrystalline gold thick films were irradiated normal to the surface

with 1.5 MeV Au+ ions using the 6MV High Voltage Engineering (HVE) Tandem accelerator at

the Sandia National Laboratories. The fluence was approximately 6.5 x 1015 ions/cm2 and

corresponding displacement per atom (dpa) was 45 dpa as shown in Figure 5.9a. After ion

irradiation we used Focused Ion Beam (FIB) based lift out technique using FEI Helios Nanolab

660 dualbeam with a working distance of 4.0mm to prepare nominally 100nm thick sample as

127

mentioned in section 4.1.2. Then we transfer electron transparent specimens on a custom

designed MEMS device, and integrated it with in-situ TEM holder (Figure 5.9b). In-situ TEM

experiments were performed inside a 200 kV FEI Talos F200X S/TEM with 1.2 Å resolution.

We also model defect annihilation phenomena using molecular dynamics (MD) simulation to

qualitatively investigate the atomic scale processes behind defect annihilation. Both specimen

preparation [212] and simulation details can be found elsewhere [189].


We passed electrical current through the specimens while observing the low temperature

region inside the Transmission electron microscope (TEM). This is shown with the dotted box in

Figure 5.9c, where the massive heatsinks (electrodes) constrain the temperature to approximately

400K as indicated by our electro-thermal simulation. This temperature rise is sufficiently low

compared to the conventional thermal annealing temperature (~0.5Tm where Tm is homologous

temperature). We choose an individual grain to investigate specific types of defects and their

annihilation mechanism under EWF as shown by TEM bright field (BF) image in Figure 5.9d.

Irradiation induced dislocation lines (green color arrow), dislocation loops (cyan color arrow),

and vacancy cluster (blue color circle) in a specimen are shown in Figure 5.9d. During EWF

annealing we increased the electrical current density in finite steps of 0.5× 105𝐴/𝑐𝑚2 and with a

hold time of 5 minutes. A comparison between Figure 5.10d and 5.9e clearly shows the same

area transformed into a highly ordered crystalline area, corroborated by the selected area electron

diffraction (SAED) pattern in Figure 5.9f. BF images as shown in Figure 5.9e indicates

dislocations annihilation during EWF annealing.

128

Figure 5.9. (a) Displacement per atom (dpa) profile for irradiation dose of 6.5×1015 ion/cm2, (b)

micro-electro-mechanical system (MEMS) device mounted on in-situ TEM holder, (c)

temperature profile obtained from electro-thermal simulation, (d) TEM BF image showing

irradiation damage, (e) BF image showing dislocation annihilation at 9.5×105 A/cm2, and (f)

SAED pattern after EWF annealing.

Figure 5.10a shows a TEM Bright field (BF) image of the 45 dpa irradiated sample

before any EWF processing. It is evident that ion irradiation generates significant damage and

defects in the specimen. The defect types and their annihilation mechansim have been discussed

in details in the following sections from both experimental and computational view points. In

this section, we choose an individual grain to investigate specific types of defects and their

annihilation mechanism under EWF. The area was chosen close to the electrode area as shown

by black dotted box in Figure 5.9c. The simulated temperature in this area is approximately 400

129

K. This is sufficiently low compared to the conventional thermal annealing temperature (~0.5Tm

where Tm is homologous temperature). TEM dark field (DF) image as shown in Figure 5.10f

indicates complete annihilation of defects and transformation towards a highly ordered

crystalline area. To investigate the defects such as dislocation lines annihilation mechanism we

studied dislocation and defects evolution at different current density near the low temperature

region of the sample (Figure 5.9c). Figure 5.10a shows TEM BF image of high defect density

with dislocation lines (green color arrow), vacancy cluster (blue color circle), dislocation loops

(cyan color arrow) and grain boundary (light red color arrow). To unambiguously trace the

defects during EWF processing, we use weak beam dark field (WBDF) imaging technique by

activating diffraction vector = 220. Figure 5.10b-5.10f show such TEM WBDF images at

different current density and corresponding dislocations evolution. Figure 5.10b shows DF image

and dislocation lines before EWF processing. At a current density of 3 × 105𝐴/𝑐𝑚2 we notice

that dislocation lines interact with each other and start to migrate towards the GB as indicated by

cyan color arrow in Figure 5.10c. In Figure 5.10b, the pink color dotted circle encloses three

different types of dislocation lines which further interacts among them under EWF. It is well

known that GBs act as defects sink [271, 283] thus could annihilate dislocation lines. Such an

annihilation is shown in Figure 5.10d near the GB as shown by pink color dotted circle. As the

current density increases dislocation lines further moves towards the GB as indicated in Figure

5.10e. Complete annihilation of dislocation lines is observed at a current density of 9.8 ×

105𝐴/𝑐𝑚2 as shown in Figure 5.10f. Our results indicate annihilation of defects is a non-linear

function of current density, which agrees with recent study on grain growth [189]. Figure 5.10b-

5.10f is an indication of defects mobility at low temperature (~400 K) under the impetus of

EWF.

130

Figure 5.10. (a) Irradiated sample before processing, (b) dark field image of pre-processed

sample showing dislocation lines, (c) dislocation lines interaction during EWF processing at

3.5×105 A/cm2, (d) partial annihilation (pink circle) of dislocation lines, and migration towards

grain boundary at 7×105 A/cm2, (e) migration of dislocation lines towards GB and partial

annihilation of dislocation lines at 9×105 A/cm2, and (f) complete annihilation of dislocations and

defects at 9.8×105 A/cm2.

To investigate how the EWF interacts with the defects and migrate them toward the sinks, we

performed MD simulation. Extensive computational experimentation shows that the fundamental

mechanism is the splitting of defects into Shockley partial dislocation. Figures. 5.11a-5.11c show

this phenomena for dislocations and vacancy clusters, while Figures 5.11d-5.11e are for stacking

fault tetrahedra (SFT). These are highly mobile (through gliding motion) defects [284] that can

131

be swept by the EWF towards the defect sinks. Our experiments took less than 5 minutes to

achieve a defect free single crystalline area where the temperature was constrained below

~400K.

Figure 5.11. Computational results on (a)-(c): Annihilation of dislocations and vacancy clusters,

and (d)-(f): SFT annihilation under EWF.

In low stacking fault energy materials, irradiation generates SFTs that are very difficult to

remove below 873K [285]. Figure 5.11d-5.11f (computational) and Figure 5.12 (experimental)

show such annihilation of SFT during EWF annealing. Figure 5.12a manifests high-resolution

TEM (HRTEM) BF image with high density of SFTs (cyan color arrow and yellow color dotted

triangle). Figure 5.12b represents an atomic resolution BF image of a single SFT defect before

EWF processing. Surface induced SFT annihilation under the effect of EWF is shown in Figure

5.12c. Such annihilations of SFT are marked by cyan color dotted arrow in Figure 5.12c. At

higher magnitude of EWF SFT could further interact with GBs as shown in Figure 5.12d, and

subsequently absorbed by the GBs. Complete defects annihilation is observed at a current density

of 9.5×105 A/cm2 and corresponding BF images is shown in Figure 5.12e and 5.12f. MD

132

simulation trajectory of SFT annihilation under EWF are shown in Figure 5.11d-5.11f, which

indicate SFTs are composed of stair-rod dislocation, and disintegrates to Shockley partial

dislocation during EWF annealing. This Shockley partial dislocation is further annihilated by

GBs as shown in Figure 5.12e. Dislocation density calculation shows that stair-rod dislocations

diminishes after 60ps simulation time as represented by Figure 5.12f.

Figure 5.12. HRTEM images of low temperature processing of irradiated materials: (a) defects in

the irradiated sample before processing, (b) HRTEM image of SFT, (c) surface induced

annihilation of SFT at a current density of 7×105 A/cm2, (d) SFT-GB interaction at 9×105 A/cm2,

(e) after EWF annealing at 9.5×105 A/cm2, and (f) HRTEM image after annealing at 9.5×105

A/cm2.

133

Our in-situ TEM study shows that EWF can introduce significant defect mobility such as

dislocations and SFT mobility even at room temperature. This non-thermal process is explained

from the interaction of defects with the EW, producing very large density of glissile and reactive

(with other defects) Shockley partial dislocations. An important feature of this process is targeted

defects momentum transfer at the non-defective regions, therefore the electrical energy input is

specifically targeted towards the defective regions, resulting in highly efficient ‘just in location’

annihilation of defects.

5.2.4 Conclusion

We studied the effects of external stimuli such as EWF on ion irradiation induced defects

annihilation in polycrystalline gold specimens at low temperature. Using heavy-ion irradiation,

we prepared specimens with significantly high density of dislocation lines, loops, SFTs and

vacancy clusters. We then passed electrical current through the specimens to investigate the

effectiveness of the EWF effect on defect annihilation by decoupling (exploiting the specimen

geometry and boundary conditions) it from the Joule heating. In-situ TEM evidence shows fast

(less than 5 minutes) and efficient transformation to defect-free crystalline regions. The

experimental and computational results support our hypothesis that EWF targets defects in a

highly localized manner to create highly mobile Shockley partial dislocations. The reactivity and

mobility of this species of defects generated by EWF lead to the fast and efficient defect

annihilation. In comparison, thermal annealing takes high temperature and longer time because

all the atoms need to be heated and their mobility is governed by random diffusion.

134

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Md Zahabul Islam

Phone: (814) 321 3140 E-Mail: [email protected], [email protected]

Education

• Ph.D.: Department of Mechanical Engineering, The Pennsylvania State University, University Park,

PA.(USA), August 2020

• M. S.: Department of Mechanical Engineering, Bangladesh University of Engineering and

Technology (BUET), Dhaka-1000 (Bangladesh), July 2013

• B. S.: Department of Mechanical Engineering, Bangladesh University of Engineering and

Technology (BUET), Dhaka-1000 (Bangladesh), February 2011

Selected Publications (Asterisk denotes undergraduate mentee)

Book Chapter 1. Potential and Challenges of in-situ Microscopy on Electronic Devices and Materials, Zahabul Islam,

Aman Haque, Wide Bandgap Semiconductor-Based Electronics, IOPscience (2020)

Journal Publications

1. Enhancement of WSe2 FET device performance using low temperature annealing, Zahabul Islam, Azimkhan Kozhakhmetov, Joshua Robinson, Aman Haque, Journal of Electronic Materials, volume 49, pages 3770–3779 (2020)

2. β-Ga2O3 Schottky diode failure under forward biasing condition: In-situ TEM study, Zahabul Islam, Minghan Xian, Aman Haque, Fan Ren, Marko Tadjer, Nicholas Glavin, S.J. Pearton, IEEE

Transactions on Electron Devices, doi: 10.1109/TED.2020.3000441 (2020)

3. Heavy ion irradiation effects on GaN/AlGaN high electron mobility transistor failure, Zahabul Islam,

Angela L. Paoletta⁕, Anthony M. Monterrosa, Jennifer D. Schuler, Timothy J. Rupert, Khalid Hattar,

Nicholas Glavin, Aman Haque, Microelectron. Reliab.,Vol. 102, pp. 113493 (2019)

4. Quality enhancement of nanocrystalline MoS2 via electrical annealing: an experimental and

computational investigation, Zahabul Islam, Kehao Zhang, Joshua Robinson, Aman Haque,

Nanotechnology, Vol. 30 (39) pp. 395402 (2019)

5. Synergy of elastic strain energy and electron wind force on thin film grain growth at room

temperature, Zahabul Islam, Huajian Gao, Aman Haque, Mater. Charact., vol. 152, pp. 85-93 (2019)

6. Real-time visualization of GaN/AlGaN high electron mobility transistor failure at off-state, Zahabul

Islam, Aman Haque, Nicholas Glavin, Appl. Phys. Lett., Vol. 113, pp.183102 (2018)

7. Departing from the mutual exclusiveness of strength and ductility in nanocrystalline metals with

vacancy induced plasticity, Zahabul Islam, Baoming Wang, Khalid Hattar, Huajian Gao, Aman

Haque, Scripta Materialia, Vol. 157, pp. 39-43 (2018)

8. Current density effects on the microstructure of zirconium thin films, Zahabul Islam, Baoming Wang,

Aman Haque, Scripta Materialia, vol. 144, pp.18-21 (2018) Awards/Honor

• Kulakowski Travel Award (Dept. of Mech. And Nuc. Engg.) for Attending MRS Fall Meeting, 2017

• First Annual Student Paper Competition Finalist (Applied Mechanics Division) IMECE-ASME, 2019

Google website: https://sites.google.com/site/mzahabulislam/

mailto:[email protected]

https://www.sciencedirect.com/science/journal/00262714/102/supp/C

https://iopscience.iop.org/volume/0957-4484/30

https://iopscience.iop.org/issue/0957-4484/30/39

https://www.sciencedirect.com/science/article/pii/S1359646218304561

https://www.sciencedirect.com/science/article/pii/S1359646218304561

https://sites.google.com/site/mzahabulislam/

Documents

MULTI-STIMULI EFFECTS ON THIN FILMS AND DEVICES