carbon material growth, characterization, and device

Carbon Material Growth, Characterization, and

Device Fabrication

by

Solomon Mikael

A dissertation submitted in partial fulfillment

of the requirement for the degree of

Doctor of Philosophy

(Electrical Engineering)

at the

University of Wisconsin Madison

2015

Date of final oral examination: 12/14/2015

The dissertation is approved by the following members of the Final Oral Committee:

Zhenqiang “Jack” Ma (Advisor), Professor, Electrical Engineering

Shaoqin “Sarah” Gong, Professor, Biomedical Engineering

Michael Corradini, Professor, Engineering Physics

Mikhail Kats, Professor, Electrical Engineering

Zongfu Yu, Professor, Electrical Engineering

i

To Anyone Taking The Less Traveled Path

ii

Acknowledgements

I would never have been able to finish this dissertation without the guidance of my

committee, professors, and help from my colleagues and friends.

I would like to express my sincere gratitude to my advisor Professor Zhenqiang (Jack)

Ma for offering me a position in his group at the University of Wisconsin Madison. His constant

support and input helped guide many of the projects I worked on through my Ph.D. study. I

would also like to thank Professor Shaoqin (Sarah) Gong, Professor Douglass Henderson,

Professor Michael Corradini, Professor Mikhail Kats, and Professor Zongfu Yu, for their

insightful comments and valuable suggestions on the research I worked on.

I would also like to thank Mr. Winslow Sargent for his generous support during my Ph.D

studies and The Graduate Engineering Research Students program (GERS) at the University of

Wisconsin Madison. Both of these programs provided support and encouragement throughout

my Ph.D studies. Both helped me reach goals I could have not reached on my own. Thank you.

I am thankful for the opportunity to work in Prof Ma’s research group and participate in

many research topics. I had the chance to meet many intelligent and thoughtful people during my

stay and wish them the best of luck in their lives and careers. Last but not least I’d like to thank

my family for their support over the years. Thank you.

iii

Abstract

The research of using different carbon allotropes has steadily developed over the years.

One of the allotropes of interest is graphene because of its unique optical and electronic

properties. Bilayer graphene unlike monolayer graphene has the potential to have the bandgap

modified. To date the largest bandgap opening for bilayer graphene is 250meV, but was done

locally (~10um) with very large bias voltages. A critical step to see materials like bilayer

graphene leave the lab is introducing wafer scale methods for electronic band modification. This

work will present the use of straining films to apply wafer scale stress to sheets of bilayer

graphene to modify the electrical properties of bilayer graphene. Using FTIR and raman

spectroscopy a bandgap of ~40meV was observed in large areas (~100umx100um).

The use of transparent neural electrode arrays with ultra-flexibility and biocompatibility

would provide an optimal platform for various applications, including optogenetics and neural

imaging. Neural electrode arrays with broad-wavelength transparency from the ultraviolet (UV)

to infrared (IR) spectrums are especially desirable, and provide unique opportunities to advance

these techniques that would otherwise be impossible with conventional opaque metal electrodes.

Also, the transparent neural electrode allows for simultaneous observation of tissue response

during optical or electrical stimulation. Graphene, a novel two dimensional carbon-based

material, is one of the most promising candidates because the material has a UV to IR

transparency of over 90 % in addition to its high electrical and thermal conductivity, flexibility,

and biocompatibility. Here we present a protocol for the fabrication of the transparent graphene

iv

neural electrode array and its operation for electrophysiology, fluorescent microscopy, optical

coherence tomography (OCT), and optogenetics

Finding appropriate measurement techniques in high temperature high radiation

environments present several challenges. This work will also introduce the development of a

temperature sensor for high radiation environments. Currently the most sensitive high

temperature thermocouples have sensitivities in the uV/C range. Next generation nuclear reactors

will have temperature ramps and power densities that will exceed the capabilities of current

thermocouples. I propose using single crystal boron doped diamond diode as a replacement for

next generation reactor temperature sensing. Due to its large bangap the sensitivity of the device

can be as high as mV/C allowing for detailed recording of quick temperature changes. In

addition to the high sensitivity the carbon in diamond and boron are two materials that are highly

radiation resistant allowing reliable operation over large fluxes and durations. I will show the

current progress on this project and the future plans for in-pile testing.

Replicating the human eye using conventional semiconductor materials and devices has

been a goal of photodetector arrays for many years. Artificial Eyes using silicon nanomembranes

on flexible polyimide substrates have been demonstrated. The device in conjuncture with the

collection setup and software allow for many of the unique capabilities of the human eye to be

realized in a process that is compatible with current semiconductor tools and methods.

v

TABLE OF CONTENTS

Acknowledgements……………….…………………….…………………………………….ii

Abstract………………………………………...………………………………………………iii

List of Figures…………………………………………………………………………….…viii

List of Tables……………………………………………………………………………..…xxii

List of Equations…………………………………………………………………….…….xxiii

CHAPTER 1 Introduction………………………………………………………………………1

1.1 Introduction and Motivation…………………………………………………….….…………1

1.2 Graphene’s Properties…………………………………………………………………………2

1.2.1 Graphene lattice and Band Structure……………………………...………………...2

1.2.2 Electronic Properties…………………………………………………………...……5

1.2.3 Optical Properties……………………………………………………………………9

1.3 Motivation and Objectives……………………………………………………...……………14

1.4 Synthesis of Monolayer and Bilayer Graphene………………………………………...……15

1.5 Dissertation Outline……………………………………………………………….…………22

Chapter 2 Bandgap Modification of AB Stacked Bilayer Graphene……………………..…23

2.1 Introduction and Motivation…………………………………………………………………24

vi

2.2 Strain Engineering of Graphene………………………………………………...……………26

2.3 Experimental Techniques & Results…………………………………………………………32

2.4 Discussion & Future Work………………………………………………………..…………68

2.5 Summary………………………………………………………………………………..……70

Chapter 3 Transparent Electrodes for Brain Implants………………………………………71


3.2 Current Methods for Brain Signal Recording…………………………………..……………72

3.3 Carbon Layered Electrode Array (CLEAR) Brain Electrode……………………..…………73

3.4 Future Work & Summary……………………………………………………………………82

Chapter 4 High Sensitivity Diamond Temperature Sensor………………….………………83


4.2 Diamond Properties………………………………………………………………….………84

4.3 Growth of Single Crystal Diamond………………………………………….………………92

4.4 Fabrication of PI Diodes for high sensitivity Temperature Sensors……………………..…111

4.5 Future Work & Summary……………………………………………………………..……118

Chapter 5 Electrical Artificial Human Eye Photo-detector Array……………...…………119

5.1 Introduction…………………………………………………………………………………119

vii

5.2 Background of Artificial Human Eye………………………………………………………120

5.3 Electrical Characterization and Image Acquisition…………………………………...……122

5.4 Summary……………………………………………………………………………………123

Chapter 6 Conclusion and Future Work……………………………………….……………125

6.1 Conclusions…………………………………………………………………………………125

6.2 Future Work……………………………………………………………………...…………125

6.3 References…………………………………………………………………………..………126

viii

List of Figures

Figure 1.1 Graphene in 0 dimensions (buckyball), 1 dimension (carbon nanotube), 2 dimensions

(graphene), and 3 dimension (graphite)…………………………………………………………...3

Figure 1.2 Graphene’s structure [1] and unit cell for both mono layer and bilayer graphene. The

unit cell for monolayer graphene has 2 atoms in it while bilayer unit cell has 4 atoms. The

rhombus is the conventional unit cell, The γ terms represent the energy of the bonding between

the respective atoms in graphene………………………………………………………………….4

Figure 1.3 Reciprocal lattice of monolayer and bilayer graphene with lattice points shown as

black dots, b1 and b2 are primitive reciprocal lattice vectors. The shaded hexagon is the first

Brillouin zone with Γ indicating the centre, and K + and K − showing two non-equivalent

corners……………………………………………………………………………………………..6

Figure 1.4 H = transfer integral matrix = describes the hopping of the π electrons between the

different carbon atoms, S = overlap integral matrix = gives us the strength of the overlap of the π

orbital's on different atoms, f (k) describes nearest-neighbor hopping. The respective gamma (γ)

terms describe interatomic hopping parameter between different combinations of atoms in the

unit cell………………………………………………………………………………………….…7

Figure 1.5 The energy dispersion for bilayer and monolayer graphene. The expression for the

energy is calculated by solving the determinate mentioned earlier in Equation 1.3. (a) shows the

monolayer dispersion where E(k) = ± kvF and for (b) bilayer graphene *2

22

m

kE

.(Michael S.

Fuhrer, University of Maryland) = planks constant, Fv = Fermi velocity, m* = effective mass

of electron in bilayer graphene…………………………………………………………………....8

ix

Figure 1.6 The types of raman bands in monolayer and bilayer graphene can be divided into (i)

defect inducted modes where the additional momentum to get total momentum transfer is almost

equal to zero is provided by elastic scattering from defects (ii) excitation of tow phonons with

wave vectors q and -q which doesn’t require any defect induced scattering for wave vector

compensation.[2]………………………………………………………………………………..11

Figure 1.7 The γ coupling terms have been observed in the Slonczewski-Weiss-McClure (SWM)

Tight binding model and been experimentally observed using Raman, FTIR, and photoemission.

For bulk bilayer graphene γ0=2.9eV, γ1=0.3eV, γ3=0.1eV, and γ4=0.12eV [3]…………..……13

Figure 1.8 The low pressure chemical deposition (LPCVD) growth system. Major components

include the furnace, label 1, flow meters, label 2, pressure sensors, label 3, and mechanical

pump, label 4. (b) The computer controls using for the LPCVD system, consisting of the

computer control system, label 5, flow controllers, label 6, and power supplies, label 7……….17

Figure 1.9 Show the growth recipies with both temperature and pressure plotted on the y-axis (a)

shows the entire growth process from start to fininsh (b) is a magnified version of the process

distinguishing the different steps in the growth process, anneal, monolayer growth, and bilayer

growth…………………………………………………………………………………………....20

Figure 2.1 (a) Top and bottom gated structure with exfoliated bilayer graphene [4] (b) Using

uniaxial strain the sample substrate is stretched while one of the layers is pinned down [5] (c)

Theoretical paper proposing using strain in bilayer graphene to open the bandgap and the

calculated change in the E(k) of the sample [6]…………………………………………...……25

x

Figure 2.2 Shows the process for growing the sample and the wet chemistry needed to get the

final graphene sample on a silicon substrate for further device processing…………………..…28

Figure 2.3 The graphene sample in different stages of processing (a) right after removal from the

LPCVD system (b) optical image of graphene on the copper foil (c) the graphene sample on Si

SiO2 (300nm) substrate (d) optical image of monolayer graphene on Si SiO2 (300nm) substrate

(e) optical image of bilayer graphene on Si SiO2 (300nm) substrate (f) scanning electron

microscope (SEM) of bilayer region on the monolayer graphene……………………………….29

Figure 2.4 The raman data for three types of graphene that are present in the LPCVD grown

material. The purple raman signal is for monolayer graphene. The green raman signal is for

bilayer graphene, and the red raman signal is for trilayer graphene…………………………..…30

Figure 2.5 The figure shows an atomic force microscopy (AFM) scan one of the bilayer regions.

The step between the stacked layers is visible in the AFM. The 0.5nm step between the layers is

close to the monolayer graphene thickness. Below a profile of the scan over the center of the

stack is shown………………………………………………………………………………...….31

Figure 2.6 (a) Schematic illustration of the layered structure of strained bilayer graphene with a

Si3N4 stressor layer. (b) Measured tensile (top)/compressive (bottom) stress values from the

layered structure which is described in Figure 1(a) with respect to the Si3N4 layer with various

thicknesses. Blue and red plots denote the stress value of a Si3N4 layer generated using a high

and medium stress Si3N4 recipe. (c) and (d) Microscopic images of the bilayer graphene layer

transferred on 4” SiO2/Si substrate before deposition of Si3N4 stressor layers. (e) Illustrations

showing the formation of wrinkles by tensile or compressive Si3N4 stressor layers (f) A

xi

microscopic image of the bilayer graphene layer after deposition of the tensile Si3N4 stressor

layer. (g) A microscopic image of the bilayer graphene layer after deposition of the compressive

Si3N4 stressor layer. The insets in Figure 1 (f) and (g) are the angled SEM images of the strained

bilayer graphene. Scale ………………………………………………..33

Figure 2.7 Microscopic images taken from (a) low compressively stressed and (b) highly

compressively stressed graphene. The images show different dimensions of wrinkles formed by a

low and high compressive Si3N4 stressor layer, respectively. The sample with a low compressive

Si3N4 stressor layer shows an average width of 2.96 μm, while the sample with a highly

compressive Si3N4 stressor layer shows an average width of 4.43 μm. Overall the wrinkles

formed by a high stressor layer have wider wrinkles. However, it is also noted that wrinkles

mostly formed around the bilayer graphene regions as indicated by white arrows……………...38

Figure 2.8 (a) A schematic cross section of the layered structure of the samples with different

degree of strains (Green: Low stress, Red: Medium stress, Blue: High stress). “Layer 1” and

“Layer 2” indicate the bottom and top graphene layer, respectively. (b) Raman shifts of the G

band (left) and 2D band (right) induced by the Si3N4 tensile stressor layer taken on the wrinkles

graphene region. (c) Raman shifts of the G band (left) and 2D band (right) induced by the Si3N4

tensile stressor layer taken on the bilayer graphene regions. (d) Raman shifts of the G band (left)

and 2D band (right) shifts induced by the Si3N4 compressive stressor layer taken on the wrinkles

graphene region. (e) Raman shifts of the G band (left) and 2D band (right) induced by the Si3N4

compressive stressor layer taken on the bilayer graphene regions……………………………....39

xii

Figure 2.9 (a) Illustration of the structures of the samples (left) without and (right) with the

Si3N4 stressor layer. Raman spectra compare the (b) G peak and (c) 2D peak without and with

the low stress Si3N4 layer. This shows that a low tensile stress (~15 MPa) does not change the G

peak position notably and only caused a minor blue-shift in the 2D peak position……………..41

Figure 2.10 (a) The measured sheet resistance under three different conditions, i.e. a bilayer

graphene (1) without any Si3N4 layer on top, (2) with a low stress Si3N4 layer, and (3) with a

highly compressive stress Si3N4 layer. (b) Microscopic images of the device used to measure the

sheet resistance. It should be noted that the result did not show any noticeable graphene sheet

resistances for case (1) and (2), whereas case (3) showed 70 % lower sheet resistance. It is

believed that the low sheet resistance is mostly caused by the high Si3N4 stressor layer, since the

effect by the deposition of Si3N4 layer or unwanted doping from the SiO2/Si substrate can be

ruled out. The average sheet resistance values for each case are 41 Ω, 48.6 Ω, and 27.3 Ω with

…………………………………..43

Figure 2.11 (a) A microscopic image of the wrinkles graphene. The arrows indicate the

measured spots and the colors of the arrows match each plot in Figure 2.11(b)-(c). (b) Raman

shifts of the G band and (c) 2D band measured by line scanning from the red region (body of the

wrinkle) to blue region (tail of the wrinkle), showing G band was spitted into two peaks (G+ and

G-) at the tail of the wrinkle. (d) A microscopic image of the wrinkles graphene indicating two

different spots with different degrees of compressive strains. The white arrows indicate the

bilayer graphene regions. (e) Raman shifts of the G band and (f) 2D band of the wrinkles

xiii

graphene taken at spot "a" and spot "b" shown in Figure 3(d). Red plots indicate the Raman

spectra taken from the “Low stress” Si3N4 layer as a reference…………………………………46

Figure 2.12 (a) Raman mapping of the highly compressive strained bilayer graphene (scan area:

100 μm2). Light blue and yellow indicate G band red-shifting. Yellow also indicates G band

splitting. Yellow indicates the location of G band splitting. (b) The overlay image of Raman

mapping and the microscopic image of the locations of splitting can be seen nearly all over the

sample……………………………………………………………………………………………48

Figure 2.13 The FTIR spectra and microscopic images of the strained bilayer graphene with a

red-arrow showing the line scanning direction. (a) high and (b) medium tensile stressed bilayer

graphene samples, respectively, (c) high and (d) medium compressive stressed bilayer graphene

samples, respectively. The band transitions that gave rise to the absorption spectra are shown for

the bilayer graphene (e) with Eg = 0 and (f) with Eg ≠ 0………………………………………...51

Figure 2.14 FTIR spectra taken over graphene with a low stress Si3N4 layer. For this particular

scan, the majority of the signal came from the Si/SiO2 substrate in the 1000 cm-1

to 200 cm-1

region. This shows that the Fabry Perot effect is difficult to completely remove from the

collected sample. The characteristic absorption peaks in the higher wavenumber values were not

observed. The absorption of single and bilayer graphene was very low resulting in the gold

standard distorting the final collected absorption………………………………………………..53

Figure 2.15 The FTIR spectra((a) and (c) and the respective Tauc’s method calculation of the

interband transitions((b) and (d) for the tensilely and compressively stressed measurements.(a)

and(b)for the high compressive-stressed sample,(c) and(d) for the high tensile-stressed sample.55

xiv

Figure 2.17 FTIR spectra of two monolayer samples stacked ontop of one another. The two

monolayers of graphene are placed one ontop of the other and a stressor layer of Si3N4 is applied

to the stack. The absorption spectrum shows minimal absorption (<1%) in the area relevant to

bilayer graphene……………………………………………………………………………….…57

Figure 2.18 Graphical illustration of the method of creating multiaxial strain by patterning

various number of spokes to generate (a) biaxial strain, (b) triaxial strain, (c) quadriaxial strain

and (d) quadriaxial axial strain, respectively, as examples……………………………………....59

Figure 2.19 Schematic illustrations and images of the fabrication process for creating triaxial

tensile strain in bilayer graphene. (i) Preparation of the CVD grown bilayer graphene. (ii) A

hexagonal shape patterning on a bilayer grpahene. (iii) Deposition of Cr claps to fix the patterned

graphene layer. (iv) Deposition of Si3N4 stressor layer on entire surface to apply a strain. (b) An

illustration to show the mechanism of the formation of tristar shape wrinkle. (c)-(e) Microscopic

images, corresponding to step (ii) – (iv). (f)-(g) Microscopic images after the deposition of low

and high Si3N4 stressor layers. Wrinkles are formed clearly. (h) A tilted SEM image taken at the

tristar shaped wrinkle…………………………………………………………………………….60

Figure 2.20 A measurement of the dimension of tristar wrinkle (Left) by SEM image and (Right)

calculation of its’ height……………………………………………………………………….....63

Figure 2.21 Raman shifts of the G band (a) and 2D band (b) on the tristar shape wrinkled

graphene induced by the Si3N4 tensile stressor with different degree of strains (Red: un-strained,

Green: low strained, Blue: high strained)………………………………………………………..65

xv



Green: low strained, Blue: high strained)………………………………………………………..66

Figure 2.23 (a) An AFM image to show the three dimensional surface profile of tristar shape

wrinkled bilayer graphene after the deposition of a Si3N4 tensile stressor layer. Inset show the top

view of the scanned region. (b) Simulated triaxial tensile strained graphene with high tensile

stressed Si3N4 layers by COMSOL Multiphysics………………………………………………..67

Figure 2.24 (A) shows a schematic of a bilayer RF graphene transistors without a straining gate

dielectric. (B) an optical image of the structure (C & D) Scanning electron microscope (SEM)

images of the bilayer RF transistor. The gate length of the transistor is 140 nm and source-to-

drain gap is 500 nm. The total gate width of two fingers is 12 μm……………………………...69

Figure 3.1 CLEAR device. a. Basic fabrication process: i. Metal patterning of traces and

connection pads on Parylene C/silicon wafer. The silicon wafer is the handling substrate. ii.

Transfer and stack four mono layers of graphene sequentially. iii. Graphene patterning to form

electrode sites. iv. Second Parylene C deposition and patterning to form device outline. v.

Removal of device from silicon wafer. b. Diagram of CLEAR device construction showing the

layered structures. c. Demonstration of CLEAR device flexibility. The device is wrapped around

of glass bar with a radius of 2.9 mm. d. Rat-sized CLEAR device: outlined by white dashed line.

e. Close-up of rat-sized device showing transparent graphene electrode sites and traces on a

Parylene C substrate. This side touches brain surface. Scale bar represents 500 µm. f. Mouse-

sized CLEAR device with ZIF PCB connector……………………………………………….…73

xvi

Figure 3.2 (a) shows the reduction of the sheet resistance as the number of layer stacked is

increased (b) compares the percentage of transmitted light since the laser light is 472nm it’s

critical that the transparent electrode has high transmission in that region of the spectrum (c)

compares the sheet resistance vs the transmission of a variety of metals and transparent

electrodes CLEAR aka graphene device is shown as a start and is comparable to many of the

materials but with much higher transmission capabilities……………………………………….76

Figure 3.3 In vivo recorded signal characterizations. a. Average longitudinal 1 kHz impedance

values for CLEAR and platinum micro-ECoG devices implanted in the same animal………….78

Figure 3.4 Optogenetic experiment (a) Schematic drawing of opto-experiment setup showing the

graphene/CLEAR device implanted on the cerebral cortex of a mouse with the light being

delivered by an optical fiber to stimulate the neural cells (b) Image of blue laser light stimulation

being delivered through the CLEAR/graphene device implanted on the cortex of a Thy1::ChR2

mouse. c. Optical evoked potentials recorded by the CLEAR device. d. Post-mortem control

data, with light impingent on electrode site 11, as is apparent by the stimulus artifact visible in

the signal for that channel. X-scale bars represent 50 ms, y-scale bars represent 100 µV……....79

Figure 3.5 In vivo imaging experiment. a. Bright-field image of CLEAR device implanted on the

cerebral cortex of a mouse beneath a cranial window. b. Fluorescence image of same device

shown in a. Mouse was given an intravenous injection of FITC-Dextran to fluorescently label the

vasculature. c. and d. Higher magnification bright-field and fluorescence images of same device

shown in a and b, respectively e. and f. Bright-field and fluorescence images of standard rat-

xvii

sized micro-ECoG array with platinum electrode sites, respectively. Scale bars in a-d represent

250 µm, while scale bars in e and f represent 750 µm…………………………………………..81

Figure 4.1 A comparison of the properties of Type IIa diamond and silicon………………...…84

Figure 4.2 Diamond unit cell with the cubic lattice structure, the lattice dimensions is about 0.36

nm and the interatomic distances are about 0.154 nm [7]………………………………….…85

Figure 4.3 The classification of different types of diamond, the different impurity levels, colors,

etc……………………………………………………………………………………………...…87

Figure 4.4 Band structure of diamond as calculated from the linear muffin tin orbital (LMTO)

method in the local-density approximation.[8]…………………………………………….….88

Figure 4.5 Comparison of diamond bandgap and dopant locations to other popular

semiconductor like Silicon, Germanium, and Gallium Nitride and the locations of the Fermi

levels for P type and N type doped materials. The bottom left plot [9] shows the resistivity and

type of conduction versus the concentration of boron acceptors at room temperature. The bottom

left image shows the conductivity in p-type diamond as a function of energy levels of boron

acceptors and temperature [10]………………………………………………………………....91

Figure 4.6 The compiled phase diagram for carbon [11]. There are two regions of interest CVD

and HPHT these two methods have allowed the creation of synthetic diamond at a much faster

rate that can be naturally mined………………………………………………………………….93

Figure 4.7 Rayleigh–Bénard convection occurs in a plane horizontal layer of fluid heated from

below, in which the fluid develops a regular pattern of convection cells known as Bénard cells.94

xviii

Figure 4.8 (a) shows the gas temperature T in Kelvin for a comparable PECVD reactor. (b)

shows the hydrogen atomic mole fraction as a percentage for substrate holder with a diameter of

9mm and power density ~120Wcm-3

(c) shows the C2 and (d) the CH3 mole fraction expressed as

a percentage [12]…………………………………………………………………………….….95

Figure 4.9 A simplified version of the Bachmann triangle showing the diamond growth region in

addition to regions where no growth and non-diamond growth occurs………………………….97

Figure 4.10 Illustration of the SiNM preparation and diffusion process for diffusion doping of

single crystal <100> Ib diamond. i. Heavy boron implantation on an SOI wafer and thermal

annealing to realize heavily doped top Si on SOI. ii. Heavily boron doped top Si layer released as

SiNM by selective etching of SiO2 . iii. Top Si picked up by an elastomeric stamp. iv. SiNM

transferred to a diamond plate. v. Bond forming between SiNM and diamond and thermal

diffusion with RTA. vi. SiNM removed by potassium hydroxide (KOH) etching………………98

Figure 4.11 Raman spectroscopy of three types of diamond, Green plot is natural Ib diamond,

Blue plot is synthetic high pressure high temperature (HPHT) diamond, and the red is synthetic

PECVD diamond. The blue dots are carbon while the white dots are hydrogen. If there is C-H

streaching the optical phonons will show up at ~3300cm-1

while if it’s only C-C stretching there

will be a strong peak at 1330cm-1

and another peak at 1550cm-1

……………………………..…99

Figure 4.12 Comparison of three types of synthetic diamond. The first left image is synthetic

PECVD diamond, the middle left is synthetic PECVD diamond with nitrogen incorporation, the

middle right is boron doped PECVD grown on a synthetic PECVD substrate, and the right image

is a heavily boron doped synthetic diamond sample……………………………………..…….101

xix

Figure 4.13 SIMS profile for boron for (a) PECVD grown samples (b) and diffusion doped

sample (c)shows the profile for additional materials that get incorporated into the film during

growth which include Si , N , O, and H………………………………………………………...103

Figure 4.14 Shows the effect on the XPS data as a reulst of the high conductively layer and a

fabricated device using the HCL as a diode at room temperature (green), 100oC (blue), and

200oC (red)………………………………………………………………………………...……104

Figure 4.15 Optical profilometry of the diamond samples showing very smooth (100) surfaces

with roughness RMS values <5nm……………………………………………………….…….107

Figure 4.16 FTIR spectra of several diamond samples. This compares the natural diamond to the

synthetic diamond spectra. For the natural diamonds the one phonon absorption peak as well as

the two phonon absorption peak is present. For the synthetic diamonds only the two phonon

absorption peak is present……………………………………………………………………....109

Figure 4.17 Diffusion doped diamond diode with XPS data. Shows the the SiNM also diffuses

nitrogen and silicon in addition to the boron. Great deal of leakage current as a result of this. The

smaller peaks to the right ~105eV and ~160eV correspond to Si incorporation into the to layers

of the lattice from diffusion ~4% in the lattice. Also nitrogen is also incorporated at ~408eV..117

Figure 4.18 Boron has two naturally occurring and stable isotopes, 11

B (80.1%) and 10

B (19.9%)

- 10

B is used in boron neutron capture therapy. The carbon in diamond is nearly all 12

C

xx

Lithium has two stable isotopes, 6Li (7.59%) and

7Li (92.4%) – the nuclear cross section of

6Li

940 barns while 7Li is 45mbarns [ 1 barn = 10

-28 m

2 ] making

7Li less affected by neutron

irradiation [KSU (P. Ugorowski) ]……………………………………………………………...112

Figure 4.19 The samples were irradiated with an average fast flux of ~ 2.63E+12 [n/cm2

s] and a

flux greater than 2.9eV of ~ 6.511E+11[n/cm2

s] for 15 minutes. This time attempts to replicate

the conditions the samples will experience during real operation……………………………...113

Figure 4.20 The left shows the ideal diode equation after some algebra extracting the sensitivity

which has the materials band gap in the exponent. The right plot shows how diamonds band gap

changes over a wide temperature range (<1%) meaning the sensitivity will stay the same even as

the environment changes [13]…………………………………………………………………115

Figure 4.21 The structure of the PECVD grown diamond and the respective IV curve from the

devices. The IV shows little leakage current while having ideality factors close to one……....116

Figure 4.22 The proposed design for the capillary for insertion into the reactor. The diode will

be inside the capillary and placed next to the fuel and the two connectors will be treaded through

an insulating material……………………………………………………………………..…….117

Figure 5.1 (a) Microscope picture of the doped membrane with etching holes. Different colors

indicate two types of doping. Shapes of each photodiode are marked out. (b), Microscope picture

for the finished silicon photodiode. Two metal layers clearly form interdigitated connection...121

xxi

Figure 5.2 (a) shows the optical setup for image creation (b) shows the convave photo detector

array (c) and (d) show the collected image using Labview and Matlab to extract and process the

collected IV data from the pixels……………………………………………………………….123

xxii

List of Tables

Table 1. Si3N4 film PECVD parameters and the measured stress on the layered samples

consisting of 25 nm Si3N4/20 nm SiO2/100 nm SiO2/ Si substrate…………………………...….35

xxiii

List of Equations

Equation 1.1 The lattice unit vectors/primitive lattice vectors for graphene……………………..4

Equation 1.2 The primitive lattice vectors which are related to the primitive lattice vectors by:

a1∙b1=a2∙b2=2π and a1∙b2=a2∙b1=0……………………………………………………………….....5

Equation 1.3 Solving the determinate will allow for the calculation of the energy levels (Ej) in

monolayer and bilayer graphene………………………………………………………………..…7

Equation 1.4 Expression for the boundary layer as a function of position on the sample……...18

Equation 1.5 The expression for the diffusion flux rate as a function of position……………..18

Equation 2.1 Expression for calculating the amount of strain in the graphene films using the

Raman data……………………………………………………………………………………….34

Equation 2.2 Expression for calculating the amount of strain in the graphene films using the

Raman data………………………………………………………………………………….……42

Equation 2.3 Equation for Tauc’s method for calculating the bandgap of a material based on the

optical absorption………………………………………………………………………………...54

Equation 2.4 Expression for the calculation of the Gruneisen parameter…………………..…..64

Equation 4.1 Position of second atom in the primitive unit cell in diamond lattice…………….86

Equation 4.2 Position of second atom in the primitive unit cell in diamond lattice…………….86

Equation 4.3 The reciprocal lattice basis vectors where V is the volume of the unit cell where i,

k , j = 1,2,3………………………………………………………………………………….…86

xxiv

Equation 4.4. Expression for the energy of electrons in the valence band maximum…………..89

Equation 4.5 Expression for the energy of an electron in the conduction band minima………..89

Equation 4.6 Expressions for the concentration of holes and the intrinsic charge carrier…...…89

Equation 4.7 The concentration of holes in the valence band due to boron…………………….90

1

CHAPTER 1 Introduction

The exponential growth of Si-based CMOS technology is rapidly approaching an end as

scaling down beyond the 10 nm node reaches fundamental physical limits [14]. The aggressive

scaling of CMOS devices has induced many drawbacks which include a dramatic increase in

fabrication costs, short-channel effects, high-field effect, quantum effect, gate leakage, process

variation, and heat dissipation issues. In the near future the cost of scaling of Si-based CMOS

will outweigh the benefits. One of the most promising candidates is graphene, which has

attracted a great deal of attention for electronic devices ever since its discovery in 2004 [15]. Its

material properties, such as intrinsic carrier mobility, saturation velocity, thermal conductivity,

and current carrying capacity, are far superior to those of silicon; moreover, its atomically thin 2-

D structure is naturally compatible with standard CMOS-based technologies. However graphene

is a gapless semimetal which is a large obstacle to overcome to be the next successor to silicon.

Thus, opening and tuning the band gap is the critical key for wider adoption of graphene for

electronic applications.

I will also introduce the development of a temperature sensor for high radiation

environments. Currently the most sensitive high temperature thermocouples have sensitivities in

the uV/C range. Next generation nuclear reactors will have temperature ramps and power

densities that will exceed the capabilities of current thermocouples. I propose using single crystal

boron doped diamond diode as a replacement for next generation reactor temperature sensing.

Due to its large bangap the sensitivity of the device can be as high as mV/C allowing for detailed

recording of quick temperature changes. In addition to the high sensitivity diamond and boron

are two materials that are highly radiation resistant allowing reliable operation over large fluxes

2

and durations. I will show the current progress on this project and the future plans for in-pile

testing.

In this dissertation the research work will focus on the growth, characterization and

straining of bilayer graphene. It will also include the growth, characterization and development

of a radiation hard temperature sensor. The thesis is organized as follows. Chapter 1 reviews the

electronic properties of graphene and reviews how others have modified the energy band

structure. Chapter 2 Will discuss the straining technique developed. Chapter 3 will look at

another application for graphene as a replacement for transparent electrodes for use in reading

brain signals in vivo. Chapter 4 will discuss the growth of single crystal diamond and the

development of the diamond diode. Chapter 5 will describe the artificial eye photo detector and

the image processing used to collect the data

1.2 Graphene’s Properties

1.2.1 Graphene lattice and Band Structure

Graphene has attracted a great deal of attention because of many unique optical and electrical

properties it has. Graphene has attracted widespread attention because of its superior properties

and enormous potential for various applications [16]. Graphene is the basis of all graphitic forms.

Graphene can be wrapped up into 0D buckyball, rolled into 1D nanotube, and stacked into 3D

graphite as shown in Figure 1.1 [17].

3

Figure 1.1 Graphene in 0 dimensions (buckyball), 1 dimension (carbon nanotube), 2

dimensions (graphene), and 3 dimension (graphite).

Monolayer graphene is a single atomic layer of sp2 bonded carbon atoms. The carbon atoms are

organized in a two dimensional hexagonal lattice structure as shown in Figure 1.2. The unit cell

of graphene has two carbon atoms with interatomic spacing of 0.1421nm [18]. The lattice unit

vectors are expressed as:

4

2

3,

2

3

2

3,

2

3

2

1

aa

aa

Equation 1.1 The lattice unit vectors/primitive

lattice vectors for graphene.

(a)

(b)

Figure 1.2 Graphene’s structure [1] and unit cell for both mono layer and bilayer graphene. The

unit cell for monolayer graphene has 2 atoms in it while bilayer unit cell has 4 atoms. The

rhombus is the conventional unit cell, The γ terms represent the energy of the bonding between

the respective atoms in graphene.

The sp2 hybridization between the s and both px and py orbital’s forms the covalent C-C bonds

between the carbon atoms. This bond is called the σ bond and forms the honeycomb lattice

structure in graphene material. The sp2 hybrids have three electrons for σ bonding. One electron

remains in a π-orbital which is known as pz. This pz orbital forms the valence π and conduction

band π* as a result of the hybridization. Bilayer graphene is made of two coupled monolayers

5

of carbon atoms, each with a honeycomb crystal structure. For both the monolayer and bilayer

samples the primitive lattice vectors are the same [19]. Next I will briefly go over the electronic

properties of graphene that make it a promising candidate for future electronic devices.

1.2.2 Electronic Properties

The valence band (π) and conduction band (π*) meet at Dirac points; K and K’ given by in the

first Brillouin zone as shown in Fig 1.3 [20]. The primitive reciprocal lattice vectors b1 and b2 of

monolayer and bilayer graphene are show in Equation 1.2:

aab

aab

3

2,

2

3

2,

2

2

1

Equation 1.2 The primitive lattice vectors

which are related to the primitive lattice

vectors by: a1∙b1=a2∙b2=2π and a1∙b2=a2∙b1=0

6

Figure 1.3 Reciprocal lattice of monolayer and bilayer graphene with lattice points shown as

black dots, b1 and b2 are primitive reciprocal lattice vectors. The shaded hexagon is the first

Brillouin zone with Γ indicating the centre, and K + and K − showing two non-equivalent

corners.

The charge carriers in graphene, which behave like massless Dirac fermion for monolayer

graphene and fermions with mass for bilayer graphene and a high mobility of up to 2×105 cm

2 V

-1 s

-1 for monolayer , make it an excellent candidate material for future generation electronic

applications [21-27]. Many of the interesting properties of graphene are a result o the carbon

based sp2 hybridized lattice structure which results in a linear band dispersion for monolayer

graphene and hyperbolic band dispersion for bilayer graphene. The band structure of these two

types of graphene can be calculated by solving the determinate of the transfer integral matrix H

subtracted from the overlap integral matrix S. The expression would appear as:

7

0)det( SEH j

Equation 1.3 Solving the determinate will

allow for the calculation of the energy levels

(Ej) in monolayer and bilayer graphene.

For monolayer graphene and bilayer graphene the respective matrix values are shown in Fig 1.4

[19]:

(a)

(b)

Figure 1.4 H = transfer integral matrix = describes the hopping of the π electrons between the

different carbon atoms, S = overlap integral matrix = gives us the strength of the overlap of the π

orbital's on different atoms, f (k) describes nearest-neighbor hopping. The respective gamma (γ)

8

terms describe interatomic hopping parameter between different combinations of atoms in the

unit cell.

Once the expression for in Fig 1.4 are solved the respective energy dispersions energy versus

momentum plot are each of the dispersions occurs in the first Brillouin zone at K + and K −

points. Fig 1.5 shows a basic illustration of the band diagrams for monolayer and bilayer

graphene.

(a)

(b)

Figure 1.5 The energy dispersion for bilayer and monolayer graphene. The expression for the

energy is calculated by solving the determinate mentioned earlier in Equation 1.3. (a) shows the

monolayer dispersion where E(k) = ± kvF and for (b) bilayer graphene *2

22

m

kE

.(Michael

S. Fuhrer, University of Maryland) = planks constant, Fv = Fermi velocity, m* = effective

mass of electron in bilayer graphene.

The main distinguishing feature between monolayer and bilayer graphene is monolayer graphene

has electrons that are massless. This is characteristic of most metallic/semi-metallic materials.

9

This results in Fermi velocities near the speed of light in ideal conditions [23]. This is also one of

the major draw backs of graphene since it implies that a bandgap doesn’t inherently exist. While

in bilayer graphene the electrons do have a mass. This implies that there is the potential for use

of bilayer graphene to be used as a semiconducting device material and flexible electronic

applications.

Creating a band gap in graphene becomes one of the most important and significant

research topics to realize graphene’s true potential. Several approaches have been proposed and

implemented that open the bandgap they include: lateral confinement of electrons using

nanomeshes or nanoribbons [28-35] or by chemical functionalization [36-38]. The issues with

these methods is the additional defects they introduce to the graphene. Many of the intrinsic

properties are lost as a result of the modification necessary to create the band gaps. Additionally

there are processing challenges that must be addressed to be able to scale the process steps

necessary to realize the bandgap/electronic modification on a larger scale.

1.2.3 Optical Properties

Graphene also has a unique set of optical properties. Each layer of graphene is able to

absorb ~2.3% of visible light [39] making it nearly transparent. This makes graphene and ideal

candidate for transparent electrodes [40, 41]. Using this property graphene can be visible to the

naked eye once it’s deposited onto SiO2/Si that’s of ~300nm of silicon dioxide. Many of the

electronic properties of graphene can be studied by optical spectroscopy [42-44]. The main

method used is Raman spectroscopy and Fourier transform infrared spectroscopy (FTIR). A

great deal of information can be obtained using these no contact methods. Raman spectroscopy

provides information about the number of layers[43], doping [45, 46], and phonon properties

10

[47] near the K point of the Brillouin zone. The major Raman peaks in graphene are the G band

(E2g) which has a raman line at ~1580cm-1

in monolayer and multilayer graphene as well as in

graphite. The E2g raman line is associated the phonon near the Γ point in the Brillouin zone. The

G band is the only Raman mode in graphene originating from a conventional first order Raman

scattering process and corresponds to the in-plane, zone center, doubly degenerate phonon mode

(transverse (TO) and longitudinal (LO) optical) with E2g symmetry [48]. The 2D Band, at

~2700cm-1

, is the excitation of 2 phonons with wave vectors q and –q which doesn’t require

defects. The D line ~1350cm

-1 , is observed in all graphitic materials with disorder associated

with phonons near the K point in the TO branch along the K-Γ direction (intervalley phonon

induced scattering). Figure 1.6 shows the raman scattering phonons for single layer and bilayer

graphene.

11

Figure 1.6 The types of raman bands in monolayer and bilayer graphene can be divided into (i)

defect inducted modes where the additional momentum to get total momentum transfer is almost

equal to zero is provided by elastic scattering from defects (ii) excitation of tow phonons with

12

wave vectors q and -q which doesn’t require any defect induced scattering for wave vector

compensation.[2].

Figure 1.6 shows that both monolayer graphene and bilayer graphene have raman phonons that

are similar and the spectrum obtained for the one material can be used to learn about the other.

This technique has be used extensively in research and will be used in this thesis to explain what

is happening to the different types of graphene.

Fourier transfer infrared spectroscopy is another technique that can used to understand

the electronic properties of graphene. The spectrum obtained for monolayer and bilayer graphene

is different as a result of the massless fermions (electrons) in monolayer and electrons with mass

in bilayer graphene [49]. During FTIR spectroscopy All the incident power is either reflected,

absorbed, or transmitted 1 = R + A + T. The fractional change in reflectance associated with the

presence of a thin-film sample is proportional to the real part of its optical sheet conductivity

σ(ℏω), or equivalently, to its absorbance A = (4π∕c)σ(ℏω). The massless fermionic character of

monolayer graphene gives a constant FTIR spectra while electrons in bilayer graphene have

finite masses and are described by a pair of hyperbolic bands and strong FTIR absorption at

~0.37eV [50]. This peak is assigned to the interband transition in undoped bilayer between the

two conduction bands or two valence bands and near the interlayer coupling energy γ1. Figure

1.7 shows the bilayer lattice structure and the interlayer coupling terms and their associated

atoms.

13

Figure 1.7 The γ coupling terms have been observed in the Slonczewski-Weiss-McClure (SWM)

Tight binding model and been experimentally observed using Raman, FTIR, and photoemission.

For bulk bilayer graphene γ0=2.9eV, γ1=0.3eV, γ3=0.1eV, and γ4=0.12eV [3].

Using the fact that the FTIR spectra of monolayer graphene is nearly constant from 0-0.5eV,

bilayer graphene has ~2% absorption of IR-Visable light, and bilayer graphene has a strong

absorption peak at ~0.3eV that is a reflection of the γ1 bonding energy one can use the FTIR

spectra as a direct indicator of modification of the electronic properties of bilayer graphene. This

powerful technique allows for measurement of the bangap over large areas by calculating the

spreading of the 0.3eV peak. As the E(k) of the bilayer graphene has additional transitions

between the conduction and valence bands additional absorption will occur as a restul. This

additional absorption can be used to calculate the introduced bandgap.

14

1.3 Motivation and Objectives

A great deal of effort is being dedicated to the band gap creation and control in graphene.

Some approaches have been proposed, such as graphene nanoribbons [31], graphene mesh [34],

and chemical functionalization [35, 51], but all above methods introduce additional serious

problems, including edge roughness, disorder, and impurities which greatly reduce the carrier

mobility in graphene. Strain engineering, because of its simple implementation and easy

fabrication process is a particularly promising approach. Previous work on straining monolayer

graphene indicates that no actual band gap opens when the lateral strain is below 20% [52]. A

perfect alternative to circumvent this problem is to use bilayer graphene, since it not only

preserves some of the exceptional electronic properties of the monolayer graphene but a band

gap can be opened easily under the right conditions. A sizable band gap opening in strained

bilayer graphene has been predicted theoretically [5]. Yet another approach is to bias bilayer

graphene, but this requires large displacement field between two layers up to an order of 2 or 3 V

per nm [4]. Thereby, straining bilayer graphene is a far better and more effective way to create a

band gap for graphene-based electronic applications.

Instead, bilayer graphene holds even more potential for electronic and digit logic

applications since it not only preserves the exceptional electronic properties as of monolayer

graphene but also can have a band gap if the symmetry between the layers is broken [4, 5]. It is a

result of the configuration of bilayer graphene which is not simply two coupled carbon layers.

Bilayer graphene is mostly found in so-called A-B or Bernal stacking [53]. In such an

arrangement, one layer does not lie directly on top of the other layer, which means only half of

15

the carbon atoms have a counterpart in the other layer and the other half are projected right into

the middle of the hexagon.

Up to now, a few approaches have been developed to overcome the above issues and some

devices have been made using complementary like graphene FETs [54-58]. Current graphene

complementary devices have low on-off ratios, low voltage gain, and gain mismatch between p-

and n-type transistors all of which are symptoms of a zero bad gap device. These methods rely

on shifting the charge neutrality point of graphene to modulate p-type and n-type behavior or

using elaborate gating configurations to open the band gap. Additionally most of the measured

data was done at cryogenic temperatures 77 K. For graphene-based CMOS to compete with

current Si electronics it must be able to operate at room temperature with high on-off ratios, have

enough voltage gain, and avoid additional processing steps.

In contrast to the existing approaches, our proposed graphene-based material using

strained bilayer graphene will meet all the requirements of conventional electronics and some

CMOS requirements. Our process will have many advantages over Si CMOS and other state-of-

art graphene devices, such as higher maximum gain, lower power consumption and better on-off

controllability.

1.4 Synthesis of Monolayer and Bilayer Graphene

Since graphene’s first isolation from bulk graphite in 2004, there have been three major

approaches developed for obtaining high quality mono- and few layer graphene sheets:

16

1. Micromechanical exfoliation of highly oriented pyrolytic graphite (HOPG) by

peeling with adhesive tape and then rubbing onto, e.g., SiO2/Si wafers [15, 59].

Graphene is first produced by this approach, but it is clearly not scalable.

2. Epitaxial growth on SiC substrate in ultrahigh vacuum and high temperature by

desorption of Si [60, 61]. It needs very high temperatures up to 1,400 oC, which is

not compatible with the CMOS process. Furthermore, ultrahigh vacuum

conditions and large SiC

3. Chemical vapor deposition (CVD) by catalytic decomposition of a gaseous

precursor on transitional metal substrates such as nickel [62], ruthenium [63],

iridium [64] and copper [65]. In the approach using CVD, graphene is grown by

chemisorption or dissolution of C from hydrocarbons such as ethylene, methane,

acetylene and benzene on the transitional metal substrates such as Ni, Ru, Ir and

Cu. It is followed by transfer of the graphene layer to another substrate for further

processing. The number of graphene layers varies with the hydrocarbon and

reaction parameters. This direct CVD synthesis provides high quality layers of

graphene without intensive mechanical or chemical treatments.

All the graphene grown in this these was grown using the CVD technique. The system

used to grow the graphene is shown in Figure 1.8 with the respective components of the low

pressure chemical vapor deposition system (LPCVD) labeled.

17

Figure 1.8 The low pressure chemical deposition (LPCVD) growth system. Major components

include the furnace, label 1, flow meters, label 2, pressure sensors, label 3, and mechanical

pump, label 4. (b) The computer controls using for the LPCVD system, consisting of the

computer control system, label 5, flow controllers, label 6, and power supplies, label 7.

The system relies on computer software to time the growth process. The software controls the

furnace and the mass flow controllers. Using a combination of the two the gases can be turned

on/off and adjusted at precise times. In addition feedback from the pressure sensor provides

additional safety by preventing over pressurizing the system during growths. Nearly all CVD

process growing graphene are performed in the mass transport diffusion controlled growth

regime. The growth temperatures typically range from 800 oC to 1400

oC depending on the metal

the graphene is grown on. At such high temperatures the growth is controlled by the mass

transport of reagents through the boundary layer to the growth surface.

CVD graphene has also been grown at a variety of pressures ranging from atmospheric

pressure to ~5mTorr. This variation causes significant changes in the growth process [66]. For

pressure 760 Torr -10 Torr there’s a large boundary layer and kinetics and mass transport

18

influence the deposition. For pressures less than 1Torr the growth is predominately controlled by

surface reactions. The rate that the precursor reaches the surface is proportional to the system

pressure indicating pressure plays a significant role in graphene growth. Our goal is to ensure

growth of large domains, for which low pressure and high temperatures were used [67].

Beyond the growth parameters variations in the Knundsen number (Kn =λ/L λ=mean free

path L=length normal to flow direction) make it difficult to distinguish growth effects from the

chamber vs. effects of different growth recipes. For the system used to grow graphene, Figure

1.8, in this paper Kn ~0.68 and the Reynolds number (Re =Ux/v, U=bulk velocity, x=position

over sample, v=kinematic velocity) is ~0.0411x. Using these values the boundary layer is:

Equation 1.4 Expression for the boundary

layer as a function of position on the sample.

and the diffusion flux rate is:

Equation 1.5 The expression for the diffusion

flux rate as a function of position

The use of the Knundsen and Reynolds number will allow an easier way to compare different

types of growths, considering the large number of growth techniques reported.

The growth recipe has two steps as shown in Figure 1.9. First the monolayer is allowed to

form on copper (Cu) surface. This is a result of the decomposition of methane onto the Cu

surface and the formation of nucleiation locations of supersaturation on the Cu surface. Once the

19

monolayer has formed the pressure inside the chamber is increased. The time required to form

the monolayer depends on hydrocarbon concentration and the pressure of the system (~5 minutes

for our configuration). The pressure change increase the Reynolds number resulting in the

boundary layer above the Cu substrate to become thinner. This modification results in restarting

the graphene growth and forming the bilayer regions which are most likely attributed to the

increased flux to the surface. The growths only appear at the initial nucleation sites formed

during the monolayer growth. Once the system reaches its equilibrium the growth of the bilayer

stops. The large single crystal domains avoid the issue faced with polycrystalline bilayer

graphene. Transistors or other type devices, potentially of multi gate fingers, can be readily

created using the very large single grains of bilayer graphene. The domain sizes can be readily

controlled by changing the step size of the pressure change. An interesting observation is the

nucleation points never overlap one another. The nucleation points always start a certain distance

away from other nucleation regions reminiscent of NW growth and their dependence on spacing

[68] . In fact, it might be more useful to understand that the bilayer regions form from the large

change in pressure. The window for monolayer graphene growth is large but the window for

large bilayer regions is much smaller. Too large of a pressure change and the Cu foil will burn

and too small there will be no formation of bilayer regions.

20

(a)

(b)

Figure 1.9 Show the growth recipies with both temperature and pressure plotted on the y-axis

(a) shows the entire growth process from start to fininsh (b) is a magnified version of the process

distinguishing the different steps in the growth process, anneal, monolayer growth, and bilayer

growth.

There are mainly three steps in graphene growth on Cu foil surfaces:

1. decomposition of hydrocarbon catalyzed by Cu,

2. nucleation of graphene from carbon atoms, and

3. lateral extension of graphene nucleus via carbon atom attachment.

The first step the hydrogen bonds are removed from the methane molecule. To break these bonds

roughly 4.8eV [69] is needed from the surrounding heat generated by the furnace. Typically the

temperatures used during graphene growth are close to 1000oC and at such high temperatures

many methane radicals will get created. During this process the Cu foil acts like a sink for theses

active methane species. [70]. The methane conversion rate at a typical growth temperature is

estimated to be on the order of magnitude of 1.0 s−1

[71]. Recalling the Renolds number for the

21

system this gives once can select a range of finance geometries, pump speeds that are ideal for

graphene growth based, and the placement of the Cu sample in the system. The effect of this is

there will be more methane and hydrogen radiacal at one end of the tube when compared to the

front end of the tube furnace show in Fig 1.8. On Cu(100) surface, there is also a large total

energy increase (2.75 eV) for methane dehydrogenation causing partially dehydrogenated

species, such as CHx , will combine with each other before going to the final hydrogen-free

product. Since the starting cooper foil is itself polycrystalline understanding how the methane

species interact with the different orientation of copper is relevant for growth control. The

different orientations (100), (110), and (111) have different adsorption energies with (100) at

6.54eV and (111) at 5.17eV [72]. Additionally since Cu (111)’s hexagonal surface lattice has

only a slight mismatch (∼3–4 %) with the graphene lattice only Cu(111) facets possess the

correct symmetry and low lattice mismatch for ideal growth. Despite this face graphene still

grows on all the facets of the polycrystalline but is most efficiently grown on samples that are

predominately (111).

Using the growth techniques mentioned above we are able to grow large quantities of monolayer

and bilayer graphene films that can be used in electronic applications. The regions of bilayer

have the potential to have their electronic properties modified and have semiconducting

properties while the monolayer graphene will continue to have its metallic properties. The

objectives of developing these growth techniques and characterization methods are to allow for

easy access to graphene materials for numerous applications such as transparent electrodes,

passive semiconductor material, and active semiconducting material.

22

1.5 Dissertation Outline

This dissertation consists of five chapters. The second chapter introduces bilayer

graphene band gap modification using straining films. First wafer scale straining is introduced

and later controlled tri axial strain application is demonstrated. The third chapter discusses using

graphene as a transparent electrode for reading brain signal in conjunction with ontogenetically

modified mice. The fourth chapter will introduce the growth of single crystal diamond for high

sensitivity temperature sensors in radiation hard environments. The fifth chapter will discuss the

development of an artificial eye and how the data from the device was collected to create images

from a hemispherical photo-detector array. The last chapter will discuss the conclusions of the

results from the different projects and future plans and ideas.

23

Chapter 2 Bandgap Modification of AB Stacked Bilayer Graphene

This chapter detail the implementation of a straining technique for bilayer graphene. The

modifications of the material and electronic structure are probed using a combination of

spectroscopic techniques to quantify the amount of strain and the effect on the band structure of

the AB stacked bilayer graphene. Strain has been applied on localized regions as well as over

large wafer scale areas. When substantial stress is applied modification of the bandgap is

observed and opening. Wafer-scale compressive and tensile strained bilayer graphene is

demonstrated by employing a silicon nitride (Si3N4) stressor layer. Different types (e.g.,

compressive vs. tensile) and magnitude of stress or strain can be engineered by adjusting the

Si3N4 deposition recipes. The strained graphene displayed significant G peak shifts and G peak

splitting when measured by Raman spectroscopy. Raman mapping of large regions of the

graphene films found that the largest shifts/splitting occurred near the bilayer regions of the

graphene films. Large area FTIR spectra showed asymmetric spectra indicating bandgap opening

in the bilayer graphene. Our unique method of graphene strain engineering can be performed

over large areas without sacrificing the desirable properties of monolayer and bilayer graphene.

Substantially large strains of up to 840 MPa were measured, corresponding to a bandgap opening

of about 40 meV on regions of bilayer graphene. Using this technique, bilayer graphene could

potentially be used to fabricate high performance graphene electronics including CMOS devices,

far infrared sensors, and terahertz sensors.

24

2.1 Introduction and Motivation

Graphene has attracted a great deal of attention for electronic devices ever since its

development in 2004 [15]. Some of its properties including intrinsic carrier mobility, saturation

velocity, thermal conductivity, and current carrying capacity, are far superior to those of silicon

[23]. Furthermore, graphene’s atomically thin 2-D structure is compatible with standard CMOS

(complementary metal-oxide semiconductor) processing technologies. However, intrinsic

graphene is a conductor, thus a band gap must be engineered into the graphene to make it

suitable for electronic devices [29]. A number of approaches have been investigated to creating

and controlling the size of the band gap in bilayer graphene including graphene nanoribbons

[31], graphene meshes [34] , and chemical functionalization [35, 51]; however, these methods

often introduce other undesirable characteristics to graphene, including additional edge

roughness, disorder, and impurities which greatly reduce the carrier mobility in graphene. Strain

engineering, due to its ease of implementation, is a particularly promising approach. Previous

reports on straining monolayer graphene indicates that no actual band gap opens when the lateral

strain is below 20 % [52]. A perfect alternative to circumvent this problem is to use bilayer

graphene, since it not only preserves the exceptional properties of the monolayer graphene

including excellent conductivity and mechanical strength, but also allows for the opening of a

band gap relatively easily under the right conditions. Additionally, previous theoretical work

predicted a sizable band gap opening in strained bilayer graphene [5] . An additional motivation

is trying to scale up the region where the band gap is forming. To date most of the work done on

bandgap engineering has been done using gated or flexible structures as shown in Figure 2.1.

25

Figure 2.1 (a) Top and bottom gated structure with exfoliated bilayer graphene [4] (b) Using

uniaxial strain the sample substrate is stretched while one of the layers is pinned down [5] (c)

Theoretical paper proposing using strain in bilayer graphene to open the bandgap and the

calculated change in the E(k) of the sample [6].

For the AB (Bernel) stacking of two graphene sheets the opening of the bandgap by pushing or

pulling the two graphene layers towards or away from each other is possible. Pulling and pushing

are inequivalent, the former is more effective in producing a band gap. For large strains a

bandgap of ~125eV has been calculated [6]. Potentially if large amount of strain is applied to

large regions of bilayer graphene a new platform for electronic devices can be made.

26

In this chapter, we report a strained bilayer graphene fabrication method capable of large

scale production without sacrificing its electrical/mechanical properties through direct deposition

of a silicon nitride (Si3N4) stressor layer on top of the bilayer graphene layer. Deposited Si3N4

layers have been shown to be able to induce different types (e.g., tensile or compressive) and

amounts (i.e., high or low) of strains by adjusting the various deposition parameters.

Additionally, it is widely used as a stressor layer in various microelectromechanical systems

(MEMS). The mechanical properties of the Si3N4 film are dependent on its chemical

composition. Both the tensile and compressive stresses are caused by the dissociation of the Si-H

and N-H bonds during the plasma deposition process. The rearrangement of the dangling bonds

on the target substrate form stable Si-N bonds during deposition. The compressive stress is

generated by a silicon rich Si3N4 layer deposited with a higher Si-N composition, while the

tensile stress is generated by a Si3N4 layer with a lower Si-N composition [73, 74]. The optical

and electrical characteristics of the strained bilayer graphene with different types and amounts of

strains were investigated using both Raman spectroscopy and Fourier transform infrared

spectroscopy (FTIR) spectroscopy. According to the Raman and FTIR analyses, a clear

relationship between the biaxial strains and the induced bandgap of the bilayer graphene is

successfully revealed.

2.2 Strain Engineering of Graphene

27

The first step in strain engineering graphene is growth and identification of large amounts

of bilayer graphene. Graphene was grown on copper foil, and then one size of the foil is spin

coated with PMMA. The sample with the PMMA protection layer floats in ferro chloride

solution to remove the copper and finally rinsed in water. The sample is then placed onto a SiO2

substrate and allowed to dry in nitrogen ambient for one day. The PMMA was then removed

through standard cleaning procedure by acetone, IPA, and DI water and annealed. Unsaturated

oxide was deposited as the straining material. With the sample completed Raman measurement

were performed over large regions. Figure 2.2 shows the process that was developed and

commonly used to transfer the graphene samples from the copper foil to the rigid silicon

substrates. The first step is growing the graphene samples in the earlier mentioned low pressure

chemical vapor deposition system with the mentioned conditions. After the growth the graphene

has formed on both the bottom and top of the copper sample. This is undesirable and the bottom

graphene is removed using a weak oxygen plasma etch. Following this step PMMA, either A2 or

A4, is spin coated onto the sample to serve as a carrier. Following a short anneal step to cure the

PMMA this stack is placed in iron chloride solution to allow the copper foil to be etched away.

What remains after the etch is a floating PMMA graphene stack. The sample is then scooped out

onto a silicon silicon dioxide (300nm thick) wafer. Once this is done the PMMA is removed

using acetone and diluted HF and the sample is annealed to ensure good adhesion to the substrate

wafer. Now the sample is read for measurement, further processing, and analysis.

28

Figure 2.2 Shows the process for growing the sample and the wet chemistry needed to get the

final graphene sample on a silicon substrate for further device processing.

Using this procedure both monolayer and bilayer graphene can be grown on large a scale. The

grown results are shown in Figure 2.3.

29

Figure 2.3 The graphene sample in different stages of processing (a) right after removal from

the LPCVD system (b) optical image of graphene on the copper foil (c) the graphene sample on

Si SiO2 (300nm) substrate (d) optical image of monolayer graphene on Si SiO2 (300nm)

substrate (e) optical image of bilayer graphene on Si SiO2 (300nm) substrate (f) scanning

electron microscope (SEM) of bilayer region on the monolayer graphene.

The samples after reaching Figure 2.3c are ready for analysis. The first property that is

characterized is the raman data for the graphene. In the image it’s clear that there are several

types of graphene present. There are monolayer regions, bilayer regions, and some trilayer

regions. Using the raman tool and the fact that the phonons in these materials is similar one can

collect and compare the raman data. The raman from these samples is shown in Figure 2.4.

30

Figure 2.4 The raman data for three types of graphene that are present in the LPCVD grown

material. The purple raman signal is for monolayer graphene. The green raman signal is for

bilayer graphene, and the red raman signal is for trilayer graphene.

A great deal of information about the quality of the film and the number of layers can be learned

from the raman specta of the graphene [45]. The identifying feature of the films is the intensity

ratio of the two peaks. It’s clear for monolayer that the 2D band is much larger. For the bilayer

the peaks are close to each other and there is a slight side lobe on the left of the 2D band. For the

trilayer the strong intensity of the G band relative to the 2D indicates strong multilayer evidence.

A Horiba micro-Raman spectroscopy (spectrometer resolution of 0.045 cm−1

) with a 50×

31

objective lens (a spot size of about 1 μm) and 18.5 mW of He-Ne (532 nm) laser light was used

to evaluate the biaxial in-plane tensile/compressive stresses in the bilayer graphene at room

temperature. The actual laser power directed to the sample is measured to be around 6.9 mW.

Using this information we can say we have quality graphene, but still don’t have any information

about the bilayer regions. To check the bilayer regions AFM scans were performed to verify the

structure and compare the known bond lengths to the measured data. Figure 2.5 shows a non

contact AFM scan over the bilayer/trilayer region.

32

Figure 2.5 The figure shows an atomic force microscopy (AFM) scan one of the bilayer regions.

The step between the stacked layers is visible in the AFM. The 0.5nm step between the layers is

close to the monolayer graphene thickness. Below a profile of the scan over the center of the

stack is shown.

In Figure 2.4 one of the stacked multilayer regions is scanned. This sample has monolayer which

is labeled 1st layer, bilayer which is labeled 2

nd layer, and trilayer which is labeled 3

rd. The

difference in high between the 2nd

and 3rd layer is shown in the plot in Figure 2.4. The 0.5nm

step between the layers is very close to the monolayer graphene thickness. Now we have optical,

Raman, and AFM data that are all saying the same thing about the LPCVD grown films.

2.3 Experimental Techniques & Results

To apply this strain we propose using compressive and tensile straining films to modify the

electronic structure of bilayer graphene. Then use Raman and FTIR to quantify the strain and

band modifications. Figure 2.6 shows the layered structure of the samples used .

33

Figure 2.6 (a) Schematic illustration of the layered structure of strained bilayer graphene with a

Si3N4 stressor layer. (b) Measured tensile (top)/compressive (bottom) stress values from the

layered structure which is described in Figure 1(a) with respect to the Si3N4 layer with various

thicknesses. Blue and red plots denote the stress value of a Si3N4 layer generated using a high

and medium stress Si3N4 recipe. (c) and (d) Microscopic images of the bilayer graphene layer

transferred on 4” SiO2/Si substrate before deposition of Si3N4 stressor layers. (e) Illustrations

showing the formation of wrinkles by tensile or compressive Si3N4 stressor layers (f) A

microscopic image of the bilayer graphene layer after deposition of the tensile Si3N4 stressor

layer. (g) A microscopic image of the bilayer graphene layer after deposition of the compressive

Si3N4 stressor layer. The insets in Figure 1 (f) and (g) are the angled SEM images of the strained

bilayer graphene. Scale bars in insets are 10 m.

The Si3N4 stressor layer was deposited on the entire surface of the SiO2/Si substrate (3 inch in

diameter) including the bilayer graphene (2” × 2”) using conventional PECVD (Plasma

34

Enhanced Chemical Vapor Deposition). As stated earlier, the type (e.g., tensile vs. compressive)

and the amount of stress generated in the bilayer graphene by the Si3N4 film can be easily

manipulated by changing the deposition conditions of the Si3N4 film. The film stress was

characterized using a stress measurement system (Tencor FleXus FLX-2320) after the

completion of the Si3N4 layer deposition. Table 1 shown below provides the detailed deposition

conditions as well as the stress types and values for the various Si3N4 films investigated during

this study.

A mixture of 2% silane (SiH4) in N2, 5% ammonia (NH3) in N2, and Nitrous oxide (N2O)

react to form silicon nitride with different amounts of stress. The detailed Si3N4 deposition

conditions are given in the supplementary section and each sample was post-annealed for 5

minutes in an N2 ambient followed by 3 minutes in high vacuum. In both cases, silicon nitride

films with a thickness of 25 nm were deposited and measured using an optical reflectometer

(Filmetrics F20). Changes in film stress were measured by using a stress measurement system

(Tencor FleXus FLX-2320). The stresses induced by the silicon nitride layer with and without

the bilayer graphene were determined by measuring the change in the curvature of the Si sample

substrate, and relating stress to curvature by the Stoney approximation Equation 2.1:

2

1 2

1 1

1 6

s

f

tE

v t R R

Equation 2.1 Expression for calculating the

amount of strain in the graphene films using

the Raman data.

where E is the Young’s modulus, ν is the Poisson ratio, R1 is the radius of the initial curvature of

the layered structure (i.e. the graphene film covered with a 20 nm SiO2 protection layer on the

35

SiO2/Si wafer), R2 is the curvature after the deposition of the Si3N4 stressor film, ts is the

thickness of the silicon substrate, and tf is the thickness of the thin film producing the stress (i.e.

the thickness of the Si3N4 layer in this experiment). Equation 2.1 is only valid when ts is larger

than tf. The Si3N4 films generate a maximum compressive stress of 745 MPa and a maximum

tensile stress of 840 MPa on the graphene for different deposition conditions.

2% SiH4

flow

(sccm)

NH3 flow

(sccm)

N2O flow

(sccm)

Pressu

re

(mTorr

)

RF

Power

(W)

Temp.

(C°)

Measured

stress (MPa)

Compress

ive stress

(High)

250 50 10 950 26 350 840

Compress

ive stress

(Medium)

175 125 10 950 46 350 530

Reference

low

stress

50 20 810 950 46 350 15

Tensile

stress

(Medium)

70 100 410 950 46 350 505

Tensile

stress

(High)

70 150 410 950 26 350 745

Table 1. Si3N4 film PECVD parameters and the measured stress on the layered samples

consisting of 25 nm Si3N4/20 nm SiO2/100 nm SiO2/ Si substrate.

It is important to note that these measurements were carried out using samples consisting of

25 nm Si3N4/20 nm SiO2/100 nm SiO2/Si substrate because wrinkles formed on the films in the

36

presence of the graphene layers, as elaborated in the next paragraph, which adversely affected

the accuracy of these measurements. Figure 1(b) shows the measured stress values with respect

to the thickness of the Si3N4 films. The tensile and compressive stress values generated by the

deposited Si3N4 layer increased as the thickness of the Si3N4 film increased. The Si3N4 film stress

reached a plateau as the thickness approached 25 nm. Thus, 25 nm thick Si3N4 films were chosen

to induce the maximum amount of tensile or compressive stress, as well as minimize the optical

interference for Raman and FTIR analyses. As shown in Table 1 and Figure 1(b), the maximum

tensile and compressive stress obtained was 745 and 840 MPa, respectively, which were

generated using the so-called high stress recipes and denoted as “high stress” in this chapter. The

medium tensile and compressive stress values were 505 and 530 MPa, respectively, which were

created medium stress recipes in this chapter. Lastly, the very low tensile stress of 15 MPa was

denoted as “low strain”. We believe that the types and amount of stress generated in the films

consisting of the 25 nm Si3N4/20 nm SiO2/bilayer graphene/100 nm SiO2/Si substrate would be

very similar to these data presented in Figure 1(b), because the atomically thin graphene layer is

not only extremely thin in comparison with the Si3N4 layer, but the Van der Waals forces also

make the graphene layer strongly adhere to the silicon substrate due to their intimate contact.

Intermolecular interactions vary in strength depending on the type of interaction. When

compared to other intermolecular forces the Van der Waals is the weakest with a strength from

0.4-4.0kJ/mol followed by hydrogen bonds at 12-30kJ/mol (1 kJ/mol = 1.04x10-2

eV per

particle). In hexagonal graphite each carbon aton in the basal plane is covalently bonded to three

nearest neighbors through 3 sp2 hybridized σ orbitals and an unhybridized 2pz orbital. Because of

these two different kinds of interactions i.e. covalent and van de Waals along different crystal

directions, the lattice structure of graphite is extremely anisotropic and is unusual in showing

37

both highest and lowest bond strengths in different directions in the same crystal. The weak

interplanar interaction between the π electrons in the adjacent planes has a spacing of 0.335nm

and the interaction is dominated by long range Van der Waals interactions. As a result of this

anisotropy the exfoliation energy has been reported to be 40-60 meV [75]. Other groups have

experimentally shown the interlayer shear strength of SiO2 graphite stack to be ~0.14 GPa[76].

This shows of the three interfaces the weakest bonding force is the interlayer force, and stress

applied to a graphite/graphene stack will modify this interaction. Annealing the samples at 350

oC during the Si3N4 deposition process further enhances the adhesion between the substrate and

graphene due to additional hydrogen bonding [77] . Figure 1(c)-(d) show the microscopic images

taken from the graphene layer transferred on a 3 inch SiO2/Si substrate before the deposition of

the Si3N4 stressor layer. Figure 1(f) and (g) were taken after the deposition of the highly tensile

(745 MPa) or compressive (840 MPa) Si3N4 layers, respectively. Irregular wrinkles with a width

of roughly 2 to 4 m were formed only at the areas where the monolayer graphene was present

for both tensile- and compressive-strained Si3N4 stressor layers. Notably, these irregular wrinkles

were typically formed along the boundary of the bilayer of the graphene at the single layer

graphene regions. Namely, the bilayer graphene regions remained flat and wrinkle free. We will

discuss a possible explanation for this phenomenon in a later section. Additionally, as shown in

Figure 2.7, larger stress led to the formation of larger wrinkles indicating that the stress was

transferred from the Si3N4 stressor layer to the graphene layer.

38

Figure 2.7 Microscopic images taken from (a) low compressively stressed and (b) highly

compressively stressed graphene. The images show different dimensions of wrinkles formed by a

low and high compressive Si3N4 stressor layer, respectively. The sample with a low compressive

Si3N4 stressor layer shows an average width of 2.96 μm, while the sample with a highly

compressive Si3N4 stressor layer shows an average width of 4.43 μm. Overall the wrinkles

formed by a high stressor layer have wider wrinkles. However, it is also noted that wrinkles

mostly formed around the bilayer graphene regions as indicated by white arrows.

39

Raman spectroscopy has been used to identify the number of layers of graphene in a sample

[43], study the edge characteristics [78] , determine the amount of doping and the disorder

related to it [46, 79], quantify thermal conductivity of the sheet [80], and most importantly,

quantify strain [47, 81-83]. The applied strain deforms the lattice which can be determined

through Raman analysis due to the Dirac point shifting from the K point. To monitor such

physical changes in the strained bilayer graphene, we carried out Raman spectroscopy analysis

on the samples with different strain conditions as illustrated in Figure 2.8(a).

Figure 2.8 (a) A schematic cross section of the layered structure of the samples with different

degree of strains (Green: Low stress, Red: Medium stress, Blue: High stress). “Layer 1” and

“Layer 2” indicate the bottom and top graphene layer, respectively. (b) Raman shifts of the G

40

band (left) and 2D band (right) induced by the Si3N4 tensile stressor layer taken on the wrinkles

graphene region. (c) Raman shifts of the G band (left) and 2D band (right) induced by the Si3N4

tensile stressor layer taken on the bilayer graphene regions. (d) Raman shifts of the G band (left)

and 2D band (right) shifts induced by the Si3N4 compressive stressor layer taken on the wrinkles

graphene region. (e) Raman shifts of the G band (left) and 2D band (right) induced by the Si3N4

compressive stressor layer taken on the bilayer graphene regions.

We first prepared the sample with a low tensile stress Si3N4 layer (~15 MPa) as a reference

sample (noted as a “low stress sample” in Figure 2.8(a)) and compared the positions of G band

peaks and 2D band peaks from samples with either a medium or high stress Si3N4 stressor layer.

It should be noted that the low stress Si3N4 layer did not affect the G band peaks and only caused

a slight blue-shift in the 2D band peaks as indicated in Figure 2.9 below.

41

Figure 2.9 (a) Illustration of the structures of the samples (left) without and (right) with the

Si3N4 stressor layer. Raman spectra compare the (b) G peak and (c) 2D peak without and with

the low stress Si3N4 layer. This shows that a low tensile stress (~15 MPa) does not change the G

peak position notably and only caused a minor blue-shift in the 2D peak position.

In general, tensile stress induces phonon softening (red shift) while compression stress causes

phonon hardening (blue shift). Figure 2.8(b) shows the G band peaks and the 2D band peaks red-

shifted by the tensile Si3N4 stressor layer on the wrinkled region. Both G band peaks and 2D

band peaks exhibited a red-shift as a higher tensile stress was applied. The amount of strain that

was directly applied to the graphene was extracted by the Equation 2.2 [84]:

42

0 0/ 2Bi graphene

Equation 2.2 Expression for calculating the

amount of strain in the graphene films using

the Raman data.

where is the Gruneisen parameter, and 0 is the unstrained peak position. Based on previous

biaxial strained graphene studies [85, 86], 1.8 was used as the value for this study. The biaxial

strain values for the medium and high tensile stressor layers were calculated to be 0.13 % and

0.18 %, respectively, with the peak shifting ratio of ~1.5 cm-1

/100 MPa. Figure 2.8(c) shows the

G band peaks and the 2D band peaks exhibited red-shift caused by the tensile Si3N4 stressor layer

on the bilayer graphene. The peak shifting ratio and biaxial strain value were calculated using the

same method described above and were 2.1 cm-1

/100 MPa and 0.26 % [86], respectively. The G

band and 2D band peak shifts (blue-shift) caused by the compressive Si3N4 stressor layer on the

wrinkled regions and bilayer graphene regions are shown in Figure 2.8(d) and (e), respectively.

The peak shifting ratios and biaxial strain values associated with the medium and high Si3N4

stressor layers were measured and calculated to be 1.52 cm-1

/100 MPa and 0.14 %, and 2.1 cm-

1/100 MPa and 0.19 %, respectively. G band peak splitting was also observed for both the

wrinkled areas and the bilayer graphene area with the high Si3N4 stressor layer. Roughly 1-1.5%

stain was reported for G band splitting for uniaxial strain [47, 85]. Since the width of the

wrinkles (i.e., ~4 m) in our film was bigger than the beam spot (1 m2), the calculated strain

values only reflect the local strain values. A higher strain at the bilayer graphene regions may

allow for significant modification of its intrinsic phonon properties thereby causing a larger

43

effect on the crystal lattice. Similarly, the sheet resistance of the bilayer graphene slightly

increased with the addition of the Si3N4 stressor layer as shown in Figure 2.10 below.

Figure 2.10 (a) The measured sheet resistance under three different conditions, i.e. a bilayer

graphene (1) without any Si3N4 layer on top, (2) with a low stress Si3N4 layer, and (3) with a

highly compressive stress Si3N4 layer. (b) Microscopic images of the device used to measure the

sheet resistance. It should be noted that the result did not show any noticeable graphene sheet

resistances for case (1) and (2), whereas case (3) showed 70 % lower sheet resistance. It is

believed that the low sheet resistance is mostly caused by the high Si3N4 stressor layer, since the

effect by the deposition of Si3N4 layer or unwanted doping from the SiO2/Si substrate can be

ruled out. The average sheet resistance values for each case are 41 Ω, 48.6 Ω, and 27.3 Ω with

44

Nevertheless, while G band peak splitting is known in mono layer graphene, G band peak

splitting in strained bilayer graphene has not been reported [6, 86, 87]. It will be shown that the

G band peak splitting is related to the position of the wrinkles. As mentioned earlier, wrinkles

tended to form around the boundary of the bilayer of graphene. As can be seen in Figure 2.8(c)

and (e), the G band peaks and 2D band peaks at the bilayer graphene regions were shifted in the

presence of tensile or compressive stresses, which agree with the previous findings of Raman

studies on biaxial graphene. This confirms the idea that the bilayer graphene regions were indeed

also strained by the Si3N4 stressor layer. Figure 2.6(e) is used to illustrate how tensile or

compressive strain is created and its effect on the graphene layers. The bottom graphene is

strongly bonded to the substrate. The graphene samples were transferred on to the SiO2/Si

substrate, and Van der Waals forces [88-91] and hydrogen bonding [77] hold this layer firmly in

place, which is typically much stronger than the interlayer bonding strength between the bottom

layer and the top layer of graphene. Due to the loosely bonded interlayer, only the top graphene

layer can be expanded or shrunk by tensile and compressive strains. For the tensile strained

graphene, the tensile Si3N4 stressor layer attempts to expand the top and the bottom graphene

layers. But, due to the tight bonding force between the bottom graphene layer and the SiO2/Si

substrate, only the top graphene layer was expanded. Because the Si3N4 stressor layer on the top

graphene layer keeps exerting to expand. However, the extents of the top graphene layer

expansion and the Si3N4 stressor layer on top of it were eventually limited by the graphene

interlayer bonding force which was estimated to be 1-7 GPa [92-94]. The force difference

mostly concentrated near the boundary of the top graphene layer. Therefore, wrinkles were

formed due to delamination of the Si3N4 layer. On the other hand, for the compressively strained

45

bilayer graphene, the compressive Si3N4 stressor layer applied compressive stress to both the top

and bottom graphene layers at the same time. However, similar to the tensile strain scenario, due

to the weak interlayer bonding force between the top and the bottom graphene layers, the top

graphene layer ended up experiencing less compression. The difference in force mostly

concentrated near the boundary between the top and the bottom graphene layers and caused

delamination of the Si3N4 stressor layer at the boundary of the top graphene layer to form the

wrinkles.

To understand the uniformity of the stress over the shape of the wrinkle, we conducted a

Raman line scan from the middle to the tip of the wrinkle on the highly compressive strained

graphene sample as shown in a Figure 2.11(a)-(c).

46

Figure 2.11 (a) A microscopic image of the wrinkles graphene. The arrows indicate the

measured spots and the colors of the arrows match each plot in Figure 2.11(b)-(c). (b) Raman

shifts of the G band and (c) 2D band measured by line scanning from the red region (body of the

wrinkle) to blue region (tail of the wrinkle), showing G band was spitted into two peaks (G+ and

G-) at the tail of the wrinkle. (d) A microscopic image of the wrinkles graphene indicating two

different spots with different degrees of compressive strains. The white arrows indicate the

bilayer graphene regions. (e) Raman shifts of the G band and (f) 2D band of the wrinkles

graphene taken at spot "a" and spot "b" shown in Figure 3(d). Red plots indicate the Raman

spectra taken from the “Low stress” Si3N4 layer as a reference.

47

The color in the arrow in Figure 2.11(a) corresponds to the plots in Figure 2.11(b) and (c).

Interestingly, as can be seen in Figure 2.11(b), the G band peak splitting was only observed near

the tail of the wrinkle (labeled with purple, blue, and green colors), which implies that the stress

becomes more concentrated at the tail of the wrinkle in comparison to its body area (labeled with

orange and red colors). It is likely due to the fact that the tail of the wrinkle was subjected to a

stress concentration effect. As shown in Figure 2.11(d)-(f), the location and the shapes of

wrinkles were mostly determined by the location and shapes of the bilayer graphene regions. The

stress values on the wrinkles were somewhat different (0.21% at a spot (a) and 0.26% at a spot

(b)) at different wrinkle locations. Therefore, narrow and curvy wrinkles were formed when two

neighboring bilayer graphene boundaries were very close. These wrinkles exhibited larger

Raman shift than the ones located at the less curved regions. Figures 2.11 (e)-(f) show the G

band peaks and 2D band peaks taken from the two different spots. By using a low stressed

graphene peaks as a reference point (i.e., red lines in Figure 2.11(e) and (f)), the G band shifts

and the 2D band peak shifts vary from 11.7cm-1

to 16.2 cm-1

and from 17.5cm-1

to 23.0 cm-1

,

respectively. Figure 2.12 shows the Raman map of the highly compressive strained bilayer

graphene on a 100 m2 area.

48

Figure 2.12 (a) Raman mapping of the highly compressive strained bilayer graphene (scan area:

100 μm2). Light blue and yellow indicate G band red-shifting. Yellow also indicates G band

splitting. Yellow indicates the location of G band splitting. (b) The overlay image of Raman

mapping and the microscopic image of the locations of splitting can be seen nearly all over the

sample

Light blue, orange and red indicate G band red-shifting. Particularly, red indicates G band

splitting (i.e., G peaks split into G- and G+ peaks as illustrated in Figure 2.11(b)). The light blue

in Figure 2.12 indicates the onset of G peak splitting and the orange and red color corresponded

to strong G band splitting and the start of G+ and G- peak formation.

Fourier transform infrared spectroscopy (FTIR) allows one to measure whether the grown

graphene’s band structure has been modified. A Thermo Scientific Nicolet iN 10 infrared

microscope with a Nicolet IZ10 FT-IR module was used to collect the FTIR spectra. The

49

microscope uses a DLaTGS detector for room temperature operation and is capable of measuring

samples as small as 25 μm with a special resolution of 16 cm-1

. The FTIR was able to measure

the absorption spectra in the frequency range of 400 cm-1

to 4000 cm-1

with the samples in a N2

ambient and at room temperature. A key characteristic to identifying modifications to the band

structure is to observe whether the symmetry of the band structure has been disrupted. Should

such a disruption occur, the FTIR spectra of the strained graphene will be shifted to higher

energies. The infrared absorption spectra were collected from 0.04 eV or 0.5 eV. The photon

energy in this region is within the energy range of the valence to conduction band transition in

bilayer graphene with and without a band gap. The opening of the band gap brings additional

features to the measured spectra. All four bilayer graphene bands in the considered energy range

are involved in the electronic transitions that affect the optical properties. The extra features are a

result of the flattening of the valance and conduction bands at the K point in reciprocal space.

According to previous reports [95, 96], around 0.4 eV is the region where the band gap will

influence the collected FTIR spectra. The flattening of the bands results in an increase in the

density of states of the bands in the measured energy range. A good indication of band gap

opening is the presence of asymmetry in the recorded spectra for samples subjected to different

amounts of strain. The electron-hole asymmetry results in a more pronounced broadening at

~0.35 eV when the bilayer graphene is hole-doped than when it is electron-doped. As with most

graphene, standard processing steps unintentionally dope the films; this results in the films being

slightly hole- or electron-doped. Natural doping as such leads to a higher absorption in the

graphene bands. The features most relevant to band gap opening are also stable in a large

temperature range. This allows for measurements conducted at room temperature to provide

relevant information of changes in the band structure.

50

Figure 2.13 shows the FTIR spectra of the compressively stressed and tensile stressed

graphene films.

52

Figure 2.13 The FTIR spectra and microscopic images of the strained bilayer graphene with a

red-arrow showing the line scanning direction. (a) high and (b) medium tensile stressed bilayer

graphene samples, respectively, (c) high and (d) medium compressive stressed bilayer graphene

samples, respectively. The band transitions that gave rise to the absorption spectra are shown

for the bilayer graphene (e) with Eg = 0 and (f) with Eg ≠ 0.

The strong peak near 0.15 eV (wavenumber of ~1209 cm-1

) is attributed to the Fabry Perot

effect in the SiO2 layer [95-97]. Previous studies have established that monolayer graphene

spectra are near constant and do not contribute to the collected spectra [49, 91, 97]. The absolute

absorption spectra of the subsrtrate+Si3N4, substrate+monolayer, and graphene+Si3N4 were taken

using a gold standard as a reference. The collected spectra in Figure 2.14 are a measurement of

the change in the absorption between these two positions.

53

Figure 2.14 FTIR spectra taken over graphene with a low stress Si3N4 layer. For this particular

scan, the majority of the signal came from the Si/SiO2 substrate in the 1000 cm-1

to 200 cm-1

region. This shows that the Fabry Perot effect is difficult to completely remove from the collected

sample. The characteristic absorption peaks in the higher wavenumber values were not

observed. The absorption of single and bilayer graphene was very low resulting in the gold

standard distorting the final collected absorption.

The FTIR absorption spectra of the samples with medium and high tensile and compressive

stress are also collected in this fashion. As shown in Figure 2.13 (a)-(d), the distinctive signals

near 0.3 eV and 0.4 eV (wavenumber of ~2400 – 3200 cm-1

) were present indicating the

54

transitions between the valance and conduction bands. The pronounced asymmetry in the

measured spectra for both tensilely and compressively stressed samples occurred as a result of

injection of electrons or holes in the bilayer graphene. Their absorption intensity also showed a

dependence on the amount of strain applied by the Si3N4 stressor layer. The most pronounced

FTIR spectrum from the highly compressive stressed bilayer graphene is shown in Figure

2.14(c). Despite taking measurements at room temperature, the characteristic asymmetries in the

bands are still visible [49]. The band transitions were calculated from the FTIR absorption

spectra using the following form of Tauc’s expression [98-100]:

2

2

gE

Equation 2.3 Equation for Tauc’s method for

calculating the bandgap of a material based on

the optical absorption.

,

where ε is the absorption intensity, ω is the angular frequency of incident radiation, and Eg is the

optical bandgap. Previous reports have plotted ε1/2

/λ against the energy h∙c/λ and extrapolated to

the linear regions of the curve to the x-axis to give the value of the optical transitions between

bands [98, 99]. Because of the Fabry Periot effect at the lower energy values, the bandgap was

not directly measured as a result of the two interfering signals. The Tauc’s method allowed for

the calculation of the band to band transitions that occurred in modified bilayer graphene under a

high stress.

55

Figure 2.15 The FTIR spectra ((a) and (c)) and the respective Tauc’s method calculation of the

interband transitions ((b) and (d)) for the tensilely and compressively stressed measurements. (a)

and (b) for the high compressive-stressed sample, (c) and (d) for the high tensile-stressed

sample.

Figure 2.15 shows the calculated expressions for the samples in Figure 2.13(a) and (c). Figure

2.15(b) and 2.15(d) show the Tauc plots for the sample shown in Figure 2.13(a) (high tensile-

56

stressed) and 2.13(c) (high compressive-stressed. The x-intercepts were in the range of 0.137 to

0.177 eV, and 0.495 to 0.619 eV for Figure 2.15(a) and (c), respectively. These values were too

large to be a measurement of the bandgap. Instead these values measure the transitions from the

lower to upper graphene bands and the intraband transitions as shown in Figure 2.13(e)-(f). For

bilayer graphene, the increased absorption in the ~0.45 eV range occurred as a result of the

transitions numbered “7” and “8” in Figure 2.13(f). These transitions are indications that the

bandgap has formed since they occurred with the absorption in the ~0.28 eV range. Bilayer

graphene that does not have a bandgap shows absorption as a result of the band to band

transitions of “2” and “3” as shown in Figure 2.13(e), but this is limited to the 0.35 to 0.41 eV

range. Using the γ1 peak the additional peak spreading is used to calculate the bandgap. A

satellite peak at roughly γ1 +∆g /2 shows up. In addition, a shoulder at about γ1 − ∆g /2 which

corresponds to interband transitions that cannot occur without the presence of a bandgap. Using

the data from the compressively stressed bilayer graphene a bandgap can be calculated using the

above expressions. The opening of the bandgap disrupts the symmetric nature of the absorption

spectra and causes additional optical transitions in the 0.5 to 0.28 eV range. In accordance with

previous published work [91], a bang gap opening of ~40 meV was estimated for the

compressively strained graphene. The absorption spectra of 2 monolayers stacked as was also

collected with a stressor layer and no additional absorption was observed in the FTIR spectra as

shown in Figure 2.17. The bandgap opening is supported by the FTIR data as well as the band

splitting observed in the Raman analysis. Using strained films allows for a transition to a wafer

scale band gap opening of bilayer graphene.

57

Figure 2.17 FTIR spectra of two monolayer samples stacked ontop of one another. The two

monolayers of graphene are placed one ontop of the other and a stressor layer of Si3N4 is

applied to the stack. The absorption spectrum shows minimal absorption (<1%) in the area

relevant to bilayer graphene.

At this point we have show how to apply strain wafer wide and quite randomly. This was

the first step in verifying that the straining method we developed is capable of modifying the

band structure of bilayer graphene. The next step was developing a technique that can viably

58

create the triaxial strained bilayer graphene on any desirable location by simple patterning was

developed. Unlike the conventional graphene strain engineering methods, the

photolithographically defined spoke patterns and tensile strained Si3N4 layer deposited by

plasma-enhanced chemical vapor deposition (PECVD) system enable to create the locally

confined triaxial strained bilayer graphene at the desire location by forming a unique tristar

shaped wrinkle. The tristar shaped wrinkle was investigated with high resolution micro-Raman

spectroscopy and atomic force microscopy (AFM) analyses, and confirmed that the 0.45% of

maximum triaxial tensile strain were created. The mechanical simulation was used to verify the

strain distribution and confirm the strain value which calculated from the Raman spectroscopy

and AFM profile. The technique presented here not only provides the practical route to create a

strained graphene at the desired location but also offers the potential of the creation of multiaxial

strain in the graphene for various types of graphene-based electronic and optoelectronic devices.

In this section we report a simple and viable method to generate triaxial strain in bilayer

graphene at the desire location by using a conventional patterning and deposition technique. The

location of the strain in the bilayer graphene was easily controlled by a spoke shaped pattern.

The use of Si3N4 stressor layer allows us to manipulate the desire intensity of the strain in the

bilayer graphene. Additionally, as proposed in a Figure 2.18, the various types of strains, namely

from biaxial to multiaxial strain, can be simply realized by changing the shape of the spoke

pattern.

59

Figure 2.18 Graphical illustration of the method of creating multiaxial strain by patterning

various number of spokes to generate (a) biaxial strain, (b) triaxial strain, (c) quadriaxial strain

and (d) quadriaxial axial strain, respectively, as examples.

As an example demonstration in this section, a pattern with three spokes was employed and

successfully created the tristar shaped wrinkle at the center of the spokes. The unique tristar

shape wrinkle formed at the center of spokes by the compressive strain enables to realize the

triaxial strain in bilayer graphene for the first time. The triaxial strain created using this

technique influences the optical phonon properties of bilayer graphene and such changes were

characterized by using high resolution micro-Raman spectroscopy and atomic force microscopy

(AFM). The mechanical simulation was used to reveal the strain distribution and the simulation

result agrees well with the strain values calculated from the Raman spectroscopy and AFM

surface profile analyses.

The schematic illustration of the device fabrication is shown in Figure 2.19

60

Figure 2.19 Schematic illustrations and images of the fabrication process for creating triaxial

tensile strain in bilayer graphene. (i) Preparation of the CVD grown bilayer graphene. (ii) A

hexagonal shape patterning on a bilayer grpahene. (iii) Deposition of Cr claps to fix the

61

patterned graphene layer. (iv) Deposition of Si3N4 stressor layer on entire surface to apply a

strain. (b) An illustration to show the mechanism of the formation of tristar shape wrinkle. (c)-(e)

Microscopic images, corresponding to step (ii) – (iv). (f)-(g) Microscopic images after the

deposition of low and high Si3N4 stressor layers. Wrinkles are formed clearly. (h) A tilted SEM

image taken at the tristar shaped wrinkle.

The process began with cleaning the 100 nm SiO2/Si substrate, followed by the layer

transfer of bilayer graphene. The growth and transfer processes of CVD bilayer graphene can be

found in section 1.4 of this thesis. The bilayer graphene was defined to hexagonal shaped with 20

nm thick SiO2 etching protection layer by photolithography. The hexagonal patterned bilayer

graphene was then firmly tied by using titanium clamps in order to prevent bilayer graphene

layer from sliding after the deposition of a stressor Si3N4 layer. The 25 nm thick Si3N4 stressor

layer was deposited on the entire surface including the patterned bilayer graphene by using

conventional PECVD (Plasma Enhanced Chemical Vapor Deposition). A mixture of 2% silane

(SiH4) in N2, 5% ammonia (NH3) in N2, and Nitrous oxide (N2O) react to form silicon nitride

with different amounts of stress. The detailed Si3N4 deposition conditions and the strain

mechanism of bilayer graphene via the Si3N4 deposition have been previously reported in Table

1. In both high and low tensile stresses, Si3N4 films with a thickness of 25 nm were deposited and

measured using an optical reflectometer (Filmetrics F20). Changes in film stress were measured

by using a stress measurement system (Tencor FleXus FLX-2320). The amount of stress

generated in the bilayer graphene by the Si3N4 film can be easily manipulated by changing the

deposition conditions of the Si3N4 film. In this experiment, two types of tensile stresses were

used; namely the 745 MPa and the 505 MPa for the high stress and low stress Si3N4 film,

62

respectively. It is important to note that these measurements were carried out using samples

consisting of 25 nm Si3N4/20 nm SiO2/100 nm SiO2/Si substrate because wrinkles formed on the

films in the presence of the graphene layers, as elaborated in the next paragraph, which adversely

affected the accuracy of these measurements. The illustration shown in Figure 2.19 (b) presents

the mechanism of the generation of strains showing the tristar shape wrinkle forms at the central

intersection of three strips. As a reference, the sample with the very low Si3N4 tensile stress of 15

MPa was prepared and denoted as “un-strain”. We believe that the types and amount of stress

generated in the films consisting of the 25 nm Si3N4/20 nm SiO2/bilayer graphene/100 nm

SiO2/Si substrate would be very similar to these data measured without the bilayer graphene

layer, because the graphene layer is not only extremely thin in comparison with the Si3N4 layer,

but the Van der Waals force also makes the graphene layer strongly adhere to the silicon substrate

due to their intimate contact. Annealing the samples at 350 oC during the Si3N4 deposition

process further enhances the adhesion between the substrate and graphene due to additional

hydrogen bonding [77]. Figure 2.19(c)-(e) show the microscopic images taken during the

fabrication process (Figure 1(c), (e) correspond to the step ii and iii in Figure 1(a), the more

process images can be found in Figure 2.18). Figure 1(f) and (g) were taken after the deposition

of the low (505 MPa) and high (745 MPa) tensile stressed Si3N4 layers, respectively. Figure

2.19(h) showed the angled scanning electron microscope (SEM) image which indicates the tristar

shape wrinkles with a height of ~70 nm. The detailed measurement and calculation are shown in

in Figure 2.20. Notably, the tristar shape wrinkle was formed at the central intersection region

where the tensile stresses from three strips were neutralized.

63

Figure 2.20 A measurement of the dimension of tristar wrinkle (Left) by SEM image and (Right)

calculation of its’ height.

The triaxial strain of the patterned bilayer graphene samples were investigated by using a

Horiba micro-Raman spectroscopy (spectrometer resolution of 0.045 cm−1

) with a 50× objective

lens (a spot size of about 1 μm) and 18.5 mW of He-Ne (633 nm). Raman spectroscopy has been

widely used to investigate the phonon vibration properties of graphene. The influence on phonon

vibration of graphene by mechanical strain can be correlated with the change in Raman

characteristic peaks of graphene. Therefore, the actual strain applied to the graphene layer can be

accurately calculated by the Gruneisen parameter and the measured Raman shifts [81, 101]. The

amount of tensile strain that was applied to the bilayer graphene at the center of tristar wrinkle

was extracted by the following equation [81, 101]:

64

0

1

h

Equation 2.4 Expression for the calculation of

the Gruneisen parameter.

where 0 is the Raman band frequency without strain, δω is the Raman band shift, γ is Grüneisen

parameter for corresponding band, and h is the hydrostatic strain in the graphene film. For

triaxial strain, h can be expressed as h =1 + 2 + 3 , i.e., the sum of the three directional strain

components of the strain, and these three components have the same value (1 =2 =3).

According to Equation 2.4, based on the previously reported Grüneisen parameter [81] and the

blue-shift in the G band peak, the calculated triaxial tensile strain was calculated to be 0.45 %.

This averaged triaxial strain value extracted from Raman shift agrees well with the value of 0.4

% derived from the angled SEM and AFM image as described in Figure 2.20.

Raman spectra taken from the sample with un-strained, low strained, and high strained

Si3N4 layer were shown in Figure 2.21. Both G band and 2D band peaks clearly show the blue-

shifting as the tensile strain value increases. G band peaks started to split when the low strained

Si3N4 stressor layer was applied, and the G band and G’ band splitting (namely, the G’ band is

the subband that is splitted from the original G band) became distinctive with the ~31 cm-1

of

shifting when the high strained Si3N4 stressor layer was applied. In the Raman spectrum of

graphene, the G band is related to the doubly degenerated E2g at the center of the Brillouin zone,

while the 2D band is related to the momentum conservation of the scattering of two phonons

with opposite waver vectors [81, 102]. Thus, strain can influence the phonon variation in the

crystal structure of graphene [103]. Specifically, the change in phonon vibration at the center-

zone

65



Green: low strained, Blue: high strained).

and the change in the double-resonance condition by the triaxial strain causes the G band and the

2D band shifts. As previously reported, the G band splits into G and G' bands and blue-shifts

under tensile strain. For uniaxial strain, the G band splits into two peaks and the G' band can

broaden [48, 104]. For biaxial or triaxial strain, in contrast, the widths of the G and G' bands are

unaffected by strain, but it is caused by the interaction of electrons, LO phonons, and interior

folded phonons at the intravalley [105]. Lu et al. revealed that the G′ peak is attributed to the

twisting between two graphene layers with a twist angle of 3–8o [106].

To further investigate the differences of triaxial strain on various spots near the center

region, the line scanning of Raman spectroscopy on the graphene was performed. The

66

microscopic image of the scanning direction near the center region is shown in an inset of Figure

2.22. Raman spectra shown in Figure 2.22(a) and (b) demonstrated that the graphene film

displayed significant blue-shifts in G band (2.4 cm-1

/%) and 2D band (4.4 cm-1

/%) with clear G

band splitting, which is undoubtedly attributed to the tri-directional compressions to the center.

The blue-shifts of G and 2D bands were measured from the other tristar wrinkled centers of the

graphene and showed consistent shifts.



Green: low strained, Blue: high strained).

As shown in Figure 2.23, the morphology of the wrinkled graphene was carefully measured

by a non-contact mode AFM over a 20 × 20 m2 area. The center part of the strained graphene

film clearly shows tristar shape triaxial strained graphene with about 70 nm in height. The small

wave-like wrinkles on each arm that is shown in the inserted 2D graphene surface profile in

67

Figure 2.23(a) suggested that the highest strain intensity was concentrated at the center of the

wrinkle, but the small strain was also existed on the arms where the Si3N4 stressor layer was

deposited. In order to analyze and evaluate the tristar shaped strain distribution at the center

region, a strain simulation was carried out by COMSOL Multiphysics with a solid stress strain

model of structure mechanics to. As shown in Figure 2.23(b), we incorporated the actual

dimensions, structure, and material parameters which were analyzed by AFM into the simulation

model. For graphene, we employed Young's modulus of 1 TPa, and Poisson's ratio of 0.3 [88,

107]. We simulated the high strain case as we experimented. As shown in Figure 2.23(b), the

simulated results correspond well with the calculated results based on the values from the Raman

spectroscopy and AFM surface profile analyses.

Figure 2.23 (a) An AFM image to show the three dimensional surface profile of tristar shape

wrinkled bilayer graphene after the deposition of a Si3N4 tensile stressor layer. Inset show the

top view of the scanned region. (b) Simulated triaxial tensile strained graphene with high tensile

stressed Si3N4 layers by COMSOL Multiphysics.

68

2.4 Discussion and Future Work

We have demonstrated the method which enables realization of both tensile and

compressive stress onto a CVD grown bilayer graphene. The method has been applied

successfully to large pieces of graphene, and has demonstrated that electrical and optical

properties of graphene can be modified with Si3N4 stressor layer. The significance of our

approach lies in the fact that it can be performed in a conventional microfabrication process, i.e.

the PECVD system, and thus easily implemented for large scale production. The large shift in

the G band peak over the bilayer regions as well as the asymmetry observed in the FTIR data

reveals that a band gap in bilayer graphene can be opened by applying the appropriate amount of

stress. We have taken this idea a step further and developed the technique that can viably create

the triaxial strained bilayer graphene by using tristar shaped patterning. Unlike the conventional

strain engineering method to graphene, our technique allows us to define the strain at the desire

location by the simple patterning, which, in other words, gives more flexibility and freedom to

apply strain at the local regions. Therefore, the technique presented here could be readily applied

a triaxial strain not only for various types of graphene-based electronic and optoelectronic

devices but also with other two dimensional materials.

Future work will include transitioning this concept to active devices like transistors. To

date the largest obstical for graphene adoption is creating a large enough bandgap. Figure 2.24

shows an example of a bilayer graphene transistor fabricated using the bilayer graphene grown in

our LPCVD system. Untailored graphene’s lack of a bandgap results in low current on/off ratios

69

in its transistors. Therefore, the majority of high speed applications using untailored graphene

have been concentrated on analog RF electronics.

Figure 2.24 (A) shows a schematic of a bilayer RF graphene transistors without a straining gate

dielectric. (B) an optical image of the structure (C & D) Scanning electron microscope (SEM)

images of the bilayer RF transistor. The gate length of the transistor is 140 nm and source-to-

drain gap is 500 nm. The total gate width of two fingers is 12 μm.

Future work would be to integrate the straining mechanism that was developed and detailed in

this chapter and integrating it with the transistor. Using a combination of the straining film and

70

patterned strain one can anticipate that a bilayer graphene transistor with switching capabilities

can be fabricated.

2.5 Summary

There are still many challenges awaiting graphene but this work has shown there is a path

to using graphene in digital applications and well and optoelectronic applications. Method to

reduce the extrinsic parasitics and properly increasing the graphene transistor size are still

needed in order to practically apply graphene transistors in RF applications. Nevertheless,

combination of electrical and physical properties [4] should ensure bilayer graphene to play an

important role in high speed RF devices and even beyond. Additional techniques in controlling

the placement of the bilayer regions have been developed by others using ion implantation [108-

111]. One could imagine using a combination of photolithography and ion implantation to

control where the bilayer regions formed then using a straining film on the graphene after it has

been transferred. In summary the method has been applied successfully to large pieces of

graphene, and has demonstrated that electrical and optical properties of graphene can be

modified with Si3N4 stressor layer. The significance of our approach lies in the fact that it can be

performed in a conventional microfabrication process, i.e. the PECVD system, and thus easily

implemented for large scale production. Ohters have using the material in combination with

block copolymers to allow for a host of strain engineering possibilities [112]. One can foresee

the feasibility of higher performance electronic applications by using a stressed bilayer graphene

such as CMOS devices, far infrared sensors, or terahertz sensors.

71

Chapter 3 Transparent Electrodes for Brain Implants


Neural interfaces allow an interface between the nervous tissue and the ex vivo external

environment. These tools provide an avenue for researchers to gain a better understanding of the

brain and provide therapy for patients with neuronal disorders. Optogenetics is a new technique

involving genetic modification of neural cells to make them respond to light stimulation. This

technique is stimulating a new wave of brain research allowing another degree control over the

brain [113]. It is now desired to optogenetically stimulate the cortex with light while

simultaneously recording the evoked response. Neural surface electrode arrays, such as micro-

electrocorticography (micro-ECoG) devices, strike a balance between invasiveness and recorded

signal quality [114-117]. One of the issues with these devices is they require the use of opaque

metallic conductive materials. If the conducting material is not opaque it is not possible to

stimulate the optogenetically modified brain cells since the metal contacts will block the

incoming stimulating light [118]. Also having transparent electrodes would allow simultaneous

imaging and brain stimulating and recording to occur. This could potentially allow correlations

to be made from the optical imaging and the brain responses [119]. The imaging under the

electrodes to date has been limited by the type of electrode materials used. We propose

development of a completely transparent micro-ECoG device to allow for simultaneous imaging

72

and brain signal recording from optical stimulation. This type of product would allow for further

insight into brain functions and advance developing therapeutic interfaces.

3.2 Current Methods for Brain Signal Recording

Transparent micro-ECoG arrays have been fabricated using indium-tin-oxide (ITO) a

transparent conductor typically used in solar cells [120]. ITO may work well in rigid devices but

is not ideal for micro-ECoG devices that are required to be flexible. This limitation prevents the

micro-ECoG from being used in many applications since intimate contact with the brain is

necessary for ideal signal recording. Also ITO deposition requires high-temperature processing

not suitable for use with the low-glass-transition-temperature Parylene substrate of the micro-

ECoG array [121, 122]. One of the biggest challenges is ITO has process dependent transparency

which limits the films ultra violet (UV) and infrared (IR) light transmission [123, 124]. Neural

imaging and optogenetics applications require the use of a wide range of wavelengths (from UV

to IR) for stimulating various opsin types and visualizing fluorescently tagged cells. Therefore,

for maximum versatility, neural interfaces that can allow light transmission with high

transparency over a broad spectrum are beneficial. Because of all these draw backs ITO based

transparent micro-ECoG devices have yet to see their full potential realized. Toward the creation

of a completely transparent, chronically stable device, useful over a broad light spectrum, we

propose a graphene-based transparent micro-ECoG array. Graphene has been widely researched

for a variety of applications due to its excellent conductivity, transferability, strength, and

tunable electronic properties [125]. In addition to its idea electrical and optical properties

73

graphene is also biocompatible and able to be produced in large scales making it an ideal

candidate for replacement of ITO for neural interfacing devices [39, 126, 127] .

3.3 Carbon Layered Electrode Array (CLEAR) Brain Electrode

Following graphene characterizations using Raman Spectroscopy, as described in Chapter 2 of

this thesis implantable graphene/CLEAR neural electrode arrays were fabricated on a 4-inch

silicon wafer. Figure 3.1 shows a simplified schematic of the fabrication process

74

Figure 3.1 CLEAR device. a. Basic fabrication process: i. Metal patterning of traces and

connection pads on Parylene C/silicon wafer. The silicon wafer is the handling substrate. ii.

Transfer and stack four mono layers of graphene sequentially. iii. Graphene patterning to form

electrode sites. iv. Second Parylene C deposition and patterning to form device outline. v.

Removal of device from silicon wafer. b. Diagram of CLEAR device construction showing the

layered structures. c. Demonstration of CLEAR device flexibility. The device is wrapped around

of glass bar with a radius of 2.9 mm. d. Rat-sized CLEAR device: outlined by white dashed line.

e. Close-up of rat-sized device showing transparent graphene electrode sites and traces on a

Parylene C substrate. This side touches brain surface. Scale bar represents 500 µm. f. Mouse-

sized CLEAR device with ZIF PCB connector.

First the wafers are coated with Parylene C films using a chemical vapor deposition

system. The connection pads and initial portions of the traces were patterened with gold via

electron beam evaporation and lift-off techniques. The use of gold for the traces and pads was to

ensure a good mechanical connection to the zero insertion force (ZIF) printed circuit board

(PCB) connectors used for reading the brain signals into the computer (Imagineering Inc, Elk

Grove Village, Illinois). The electrode sites and parts of the traces that are going to be in contact

with the brain were left for the following graphene transfer and patterning to allow the brain

contact are of the electrode to remain transparent. Four graphene monolayers were transferred

and stacked creating a four layer stack onto the wafer surface using the wet transfer technique. A

sacrificial layer of silicon dioxide was deposited to protect the graphene layers from being

damaged during layer reactive ion etching (RIE). The graphene was then patterned to form the

electrode sites and another layer of Parylene C was deposited. RIE was then used to expose the

75

electrode sites and pads and form the array outlines. The last step was peeling the device from

the wafer and the SiO2 protection layer was removed using a diluted HF etch and the array was

inserted into the PCB connectors.

To test the effectiveness of the device the impedances were evaluated at 30 different

frequencies ranging from 10 Hz to 30,937 Hz. If electrode sites had impedance values less than

600 kOhms at 1 kHz frequency, they were considered to be viable for implantation. 1 kHz

frequency was selected for evaluation because it is known to be a common benchmark for neural

impedance analysis [128, 129]. The phase angle is higher in the graphene device when compared

to a similar platinum device. This means that the value of the reactance is higher for the graphene

sites than for the platinum. The average magnitude of the impedance at 1 kHz frequency was

only slightly higher for the graphene device than for the platinum array (243.5 ± 5.9 kΩ for

graphene vs 188.8 ± 92.9 kΩ for platinum). This means that the graphene device is a viable

alternative and will allow comparable signals to be recorded. Cyclic voltammetry (CV) was also

performed on the devices to determine the amount of charge the devices can carry. When the

graphene device is compared to the platinum and gold devices it’s clear that the platinum device

is capable of moving the most charge. The graphene device was more similar to the gold device

and it’s well known that gold is also another viable material for brain recording electrodes

suggesting that the graphene device would perform as well as gold which is known to work [130,

131] . To test the artifact, the devices were placed face-down in saline solution and a 200 μm

optical fiber connected to a 100 mW, 473 nm diode LASER (Laserglow Technologies, Ontario,

Canada), was used to shine light onto the backs of the electrode sites. The light pulses were

delivered by applying 3 V to the LASER for 3 ms (up to 80 mW/mm2). Figure 3.2 compares the

graphene device to many other structures used as transparent electrodes. It’s clear that the

76

graphene device has a combination of properties that make it ideal as a transparent electrode

which moderate sheet resistance.

Figure 3.2 (a) shows the reduction of the sheet resistance as the number of layer stacked is

increased (b) compares the percentage of transmitted light since the laser light is 472nm it’s

critical that the transparent electrode has high transmission in that region of the spectrum (c)

compares the sheet resistance vs the transmission of a variety of metals and transparent

electrodes CLEAR aka graphene device is shown as a start and is comparable to many of the

materials but with much higher transmission capabilities.

77

For the graphene/CLEAR device, an average of about 90% of the light impinging on the

substrate is transmitted at the desired wavelengths (470 nm for excitation of channelrhodopsin

and 570 nm for halorhodopsin). This is similar to previously reported results, and is sufficient for

many optogenetic and imaging applications [41, 120]. To demonstrate in vivo performance of

the CLEAR/graphene devices, the arrays were implanted in four rats and five mice, one wild-

type for imaging and four Thy1::ChR2 (Jackson Labs, 012350) for imaging and optogenetic

testing. Surgical procedures and in vivo imaging sessions were performed under anesthesia, and

all efforts were made to minimize animal discomfort. Once the samples were implanted the

baseline signals for impedance were collected and the device was connected to a TDT PZ2

amplifier before being sent to the RZ2 system. In addition to baseline signal recording and

impedance measurements the animals were testing for electrical evoked potentials. This was

done by stimulating the hindlimbs of the animals with electrodes above and below the sciatic

nerve. The potentials were recorded with stimuli applied on the same and opposite sides to the

implanted device to verify the response was a somatosensory response to the electrical stimulus

signal. If this was true, evoked potentials would be seen only when the stimuli were applied

contralateral to the implanted electrode array, due to the crossing of the neural pathways in the

brainstem and spinal cord. The graphene electrode sites are capable of recording both

spontaneous baseline activity and evoked neural signals with the same level of clarity as the

platinum sites, and generally similar impedance behavior and stability over time as shown in

Figure 3.3.

78

Figure 3.3 In vivo recorded signal characterizations. a. Average longitudinal 1 kHz impedance

values for CLEAR and platinum micro-ECoG devices implanted in the same animal.

Three mice with the expressed Thy1::ChR2 gene were implanted with a

CLEAR/graphene device for the purpose of evaluating the viability of these devices. The mice

had neurons expressing the Channelrhodopsin-2 protein, making them susceptible to excitation

when in contact with blue (473 nm) light. The next step was implanting the graphene electrode

onto the surface of the cortex. In this experiment the brain was left open and an optical fiber

attached to the 473 nm laser was brought close to the opening as shown in Figure 3.4.

79

Figure 3.4 Optogenetic experiment (a) Schematic drawing of opto-experiment setup showing the

graphene/CLEAR device implanted on the cerebral cortex of a mouse with the light being

delivered by an optical fiber to stimulate the neural cells (b) Image of blue laser light stimulation

being delivered through the CLEAR/graphene device implanted on the cortex of a Thy1::ChR2

mouse. c. Optical evoked potentials recorded by the CLEAR device. d. Post-mortem control data,

with light impingent on electrode site 11, as is apparent by the stimulus artifact visible in the

signal for that channel. X-scale bars represent 50 ms, y-scale bars represent 100 µV.

Concentrated blue light, with a maximum power of 80 mW/mm2, was then directed onto various

80

regions of the brain, through the CLEAR device as shown in Figure 3.4b, while simultaneously

recording the neural response to the optical stimulation. The average evoked response is shown

for three different stimulation levels in Figure 3.4c. The initial peak is the stimulus artifact

resulting from the, and the second, longer peak is the evoked neural response. that the CLEAR

device is a suitable technology for optogenetic experiments.

Once experimentation was complete, the animal was euthanized with an intraperitoneal

injection of Fatal PLUS pentobarbitol solution, and a control experiment was conducted with the

electrode on the brain of the euthanized animal, to verify that the signals recorded were from

neurons affected by the light stimulation, and not solely due to the artifact. From Figure 4d we

can see that the signal magnitude is significantly lower for the recordings obtained from the post-

mortem control experiment than for the signals recorded from the living animal. Furthermore, for

the control, there was only an evoked signal on the channels which experienced direct light

stimulation, whereas Figure 4c shows a large spatial distribution of the signals from the live

animal. These results demonstrate that the signals in Figure 4c were evoked neural responses to

the light stimulation, while those in Figure 3.4d were a result of the artifact.

A subset of the implanted animals were imaged via the cranial window imaging method

previously described by Schendel et al [119]. Representative images of the cortical vasculature

through the CLEAR micro-ECoG device are shown in Figure 3.5a-d. Images in the left column

were taken in bright-field, while those on the right were taken under blue (470 nm) light with the

aid of a tail vein injection of FITC-Dextran to fluorescently label the vasculature. These images

demonstrate the clarity of the graphene electrode sites and the ability to view the underlying

cortex and cerebral vasculature through the CLEAR device

81

Figure 3.5 In vivo imaging experiment. a. Bright-field image of CLEAR device implanted on the

cerebral cortex of a mouse beneath a cranial window. b. Fluorescence image of same device

shown in a. Mouse was given an intravenous injection of FITC-Dextran to fluorescently label the

vasculature. c. and d. Higher magnification bright-field and fluorescence images of same device

shown in a and b, respectively e. and f. Bright-field and fluorescence images of standard rat-

sized micro-ECoG array with platinum electrode sites, respectively. Scale bars in a-d represent

250 µm, while scale bars in e and f represent 750 µm.

82

3.4 Future Work & Summary

The results of this study demonstrate that the CLEAR micro-ECoG device is capable of

recording neural signals with the same degree of clarity as the platinum array, and a comparable

longitudinal tissue response. Unlike the platinum array, the CLEAR device allows for

optogenetic stimulation and both fluorescence and OCT imaging directly through the electrode

sites, due to the broad spectrum transparency of graphene. Although future studies will be

necessary to determine the long-term stability of this device, both in terms of biocompatibility

and recorded signal quality, these findings, and previous studies reporting the biocompatibility of

CVD graphene, suggest that the CLEAR device is a viable micro-electrode array for neural

interfacing applications. This graphene device is superior to the present ITO-based transparent

electrode technology, for its dramatically increased mechanical flexibility and greatly enhanced

transparency in relevant spectral ranges. The tunable electrical properties of graphene could lead

to future integration of active electronic elements into these devices. Future directions for

transparent neural interfacing studies may include exploration and implementation of these

properties with CLEAR technology.

83

Chapter 4 High Sensitivity Diamond Temperature Sensor


We propose to improve upon conventional temperature sensors (improving

both lifetime and sensitivity) with thin film synthetic diamond semiconductors. The proposed

semiconductor films act as diodes (PN junctions) in which the intrinsic carrier concentration

changes with temperature. If a constant current is forced through the PN junction, the voltage

drop across the forward-biased junction will be a linear function of the junction temperature.

The most appropriate semiconductor material is diamond, due to its stable chemical properties,

large bandgap energy, and very low surface recombination velocity. Also, diamond has a high

thermal conductivity, which, given its small size will provide a superior response time (<< msec)

governed by the material with which it is in contact. Using a series of cleanroom processing

steps single crystal diamond diodes are fabricated. The primary tasks to be addressed by this

sensor development task will be evaluation of the concept (temperature range, sensitivity,

compatibility, etc.), assessment of thermal aging and radiation damage issues, and optimization

for test reactor applications. We propose to develop a new class of temperature sensors that will

outperform conventional sensors in terms of response time, sensitivity, linearity, stability, cost,

lifetime, operating range, and shock resistance. We propose to use PN junctions fabricated from

diamond films to realize these new sensors. Others have attempted to use other wide bandgap

material [132] but diamond still has the largest bandgap and is the most radiation resistant

making it ideal for use near or around nuclear fuel rods.

84

Figure 4.1 A comparison of the properties of Type IIa diamond and silicon.

4.2 Diamond Properties

Single crystal diamond has many unique electrical, thermal, and mechanical properties.

Diamond is simply one of the allotropes of carbon that can be formed the others include graphite,

fullerenes, and carbon nanotubes. Despite the same elemental composition the differences in

85

structure result in starkly different material properties. There are two stable isotopes of carbon in

natural diamond, 98.9% of the natural abundance is 12

C and the rest is 1.1% 13

C. Diamond

consists of tetrahedrally bonded (sp3 hybridized) carbon-12 atoms while materials like graphite

or graphene are made of layers of sp2 bonded carbon sheets. This difference results in a host of

different materials properties as shown in Figure 4.1 when compared to a well known material

like Silicon [133]. These unique properties are exploited in the development of heat sink

products, cutting tools, and electronic detection of high energy particles.

The cell strucutre of diamond is face centered cubic (FCC) lattice with a bases of two

atoms at (000) and (1/4,1/4,1/4) [134] as shown in Figure 4.2. Each carbon atom joins four other

carbon atoms in regular tetrahedrons:

Figure 4.2 Diamond unit cell with the cubic lattice structure, the lattice dimensions is about 0.36

nm and the interatomic distances are about 0.154 nm [7]

The position of one atom is at the origin 0, and the orthogonal coordinate system made up of the

unit vectors . The position of the second atom in the primitive cell is given by:

86

Equation 4.1 Position of second atom in the

primitive unit cell in diamond lattice.

Where a is the point base. The FCC lattice is described by a four point bases:

Equation 4.2 Position of second atom in the

primitive unit cell in diamond lattice.

These points help define the electrical properties of diamond through the reciprocal lattice bases

as defined as:

Equation 4.3 The reciprocal lattice basis

vectors where V is the volume of the unit cell

where i, k , j = 1,2,3

It’s this unique crystal structure that gives single crystal diamond high thermal

conductivity and low thermal expansion coefficient of making it ideal for uses where dissipation

of heat is critical for device performance, such as in the creation of x-ray lenses [135] or Raman

lasers [136]. Diamond finds use in optical windows [137] because of its lack of any significant

absorption of electro-magnetic radiation (ranging from the far infrared to deep UV ). Its high

resistance to radiation, high density and chemical stability to hydrogen plasmas also makes

diamond a promising material for use as shielding within fusion reactors [138]. It is also

possible to dope diamond and create a semiconductor. This combined with the chemical

inertness and bio-compatibility of diamond allows for a number of potential applications in the

production of bio-sensors [139] or bio-electronics, such as eye implants [140].

87

There are several different types of diamond available. Figure 4.3 give a brief over view

of the different types of diamond and their classification and names.

Figure 4.3 The classification of different types of diamond, the different impurity levels, colors,

etc.

There are four types of natural diamond (Ia, Ib, IIa, IIb), classified according to the presence of

nitrogen in the crystal and certain other properties. Type-IIb diamonds contain so little nitrogen

that the crystal is a p-type semiconductor due to trace amounts of boron. Unfortunately this

type is very rare and expensive. Therefore alternative sources must be developed but first the

electronic properties for diamond will be discussed.

The energy band structure of diamond comes from a superposition of wave function of

the electrons in the primitive cell. At the point near the band edge the maximum energy of the

valence bands touches the wave vector at k0. The ground state of the electron is in the carbon

structure corresponding to diamond is in the sp3 structure. The valence band is three fold

degenerated when there isn’t any spring orbit split. The conduction band has a spherical energy

88

surface orientated along a <100> crystal axis, with symmetry point at X1 , if the valence band

lies at center zone as shown in Figure 4.4 [8, 141, 142].

Figure 4.4 Band structure of diamond as calculated from the linear muffin tin orbital (LMTO)

method in the local-density approximation.[8]

The coordinates in Figure 4.4 describe the energy band structure of diamond along <100> and

<111> axis of the Brillouin zone at L, Γ, and X [143]. The band diagram indicates that diamond

is an indirect band gap material. The indication of ∆min shows the minimum location of the

valence band of diamond while the valence band maximum occurs at the Γ position giving

diamond and indirect bandgap of 5.4 ± 0.0005 eV at room temperature. If once looks at the

momentum point at Γ a direct gap can be measured of roughly 7.02 ± 0.02 eV at room

temperature. Since diamond is such a wide bandgap material it’s important a understand of the

electron energies is understood before implementing electronic devices. The energy of an

electron in the valence band maxima is given by Equation 4.4

89

Equation 4.4. Expression for the energy of

electrons in the valence band maximum.

In the above expression the terms are defined as with K = wave vector of holes in the valence

band, Ev is the energy of the valence band, h = planks constant, mhh and mlh are the mass of the

heavy holes and mass of the light holes respectively. So in type IIa diamond mhh = 1.1 me, where

me is the rest mass of a free electron, and mlh = 0.3 me. The energy of an electron in the

conduction band minimum is given by Equation 4.5 [134]

Equation 4.5 Expression for the energy of an

electron in the conduction band minima.

In the above expression the terms are defined as Ec is the energy in the conduction band, Eg is the

energy of the bandgap. The conduction band energy for diamond is spherical along the <100>

axis. Since the focus of this section will be pi diamond diodes expressions for the concentration

of holes and intrinsic carriers are given in Equation 4.6 [134].

Equation 4.6 Expressions for the

concentration of holes and the intrinsic

charge carrier

90

Figure 4.5 show diamond with negative and positive electron affinity compared to other

materials. Using the above expressions one can determine the types of carrier concentrations that

will be present in the diamond material. But the next issue that results is for the large band gap

materials is getting the activated dopants into the conduction or valence bands. For diamond the

donors and accpetors are extremely deep when compared to other semiconductors. For single

crystal diamond the most popular dopants are boron for p-type and phosphorus for n-type. For

the shallowest dopant in diamond the activation energies (Ea) at room temperature is 0.37 eV.

For comparison in silicon <100> the same dopants can be using in single crystal silicon but their

activation energies are much lower at room temperature for <0.03 eV [144]. The theoretical

result of about 0.41 eV [144] for the ground state is very close to the experimentally obtained

value of about 0.37 eV. The activation energies introduced by the impurity states in diamond

define the type of conduction regime of the material as a function of temperature and

concentration of the boron acceptor. The conduction types are classified as band, hopping [145]

[10] and metallic conductions. The band conduction describes the conductivity in the valence

band and the hopping conduction is made up of two types: nearest neighbor hopping (NNH) and

the variable range hopping (VRH). The conduction through NNH has transition of an electron to

a nearest unoccupied level, and for VRH the conduction happens between levels separated by a

hopping distance and has an associated probability [9]. The metallic conduction regime is

achieved at concentrations that are much larger than 1x1020

cm-3

which is above the Mott

transition limit. Equation 4.7 expresses the hole concentration:

Equation 4.7 The concentration of holes in the

valence band due to boron

91

Figure 4.5 Comparison of diamond bandgap and dopant locations to other popular

semiconductor like Silicon, Germanium, and Gallium Nitride and the locations of the Fermi

levels for P type and N type doped materials. The bottom left plot [9] shows the resistivity and

type of conduction versus the concentration of boron acceptors at room temperature. The bottom

left image shows the conductivity in p-type diamond as a function of energy levels of boron

92

acceptors and temperature [10].

Despite this diamonds still has many properties that make it ideal for extreme environment

application including high breakdown field strength of 10 MVcm–1

, inertness to many chemicals

and radiation hardness, thermal conductivities >2000 Wm–1

K–1

, etc. Now we have discussed

many of the relevant electronic properties of diamond the next step is the growth of usable

material for the diode fabrication. The next section will discuss the growth techniques used in

this study.

4.3 Growth of Single Crystal Diamond

Over the years many methods have been developed for growing diamond. In the past

researchers had to rely on diamonds found in naturally occurring deposits in mines. The mail

issue with this is the inconsistent supply of diamond that is useful for semiconductor

applications. This all changed when a technique to grown diamond using a high pressure high

temperature method was developed [146]. This method was an attempt to copy the method of

natural diamond formation in the earth. The HPHT processing requires heating graphitic carbon

to over 2000 K while also compressing it to pressures greater than 5 GPa within the present of a

metal catalyst. This process allows for the creation of single crystal diamond up to a few

millimeters in size.

93

Figure 4.6 The compiled phase diagram for carbon [11]. There are two regions of interest CVD

and HPHT these two methods have allowed the creation of synthetic diamond at a much faster

rate that can be naturally mined.

An alternative technique to grown diamond synthetically is chemical vapor deposition

(CVD). The process involves starting with some diamond material initially, either in powder or

single crystal form and growing additional diamond using those seeds by maintaining an

appropriate growth environment. The growth process is carried out of a hot substrate ranging

from 700-1200oC [12]. Diamond can also be grown on metal surface [147] but this section will

focus on homoepitaxial growth of diamond. Addition of gas phase species (methane radicals) to

the surface of these crystals results in their growth. Depending on the gas-phase conditions and

the nucleation density either a large single crystal of diamond or a polycrystalline film (which

can have a large variety of grain sizes ranging from µm to nm) is grown. Using the CVD process

the growth can be divided into three phases: (1) the activation of the gas mixture and the

reactions between the gas phase species within the mixture (2) gas-phase reactions, and (3) gas

surface and surface reactions which incorporates the gas species into the bulk diamond

94

sample/substrate. Of the two method of hot filament and plasma chemical vapor deposition the

most popular is PECVD since the plasmas balls allow for larger deposition areas an more

uniform depositions. This section will focus on the plasma enhanced chemical vapor deposition

of single crystal diamonds. The PECVD reactors function by coupling the microwave power via

a waveguide and antenna through a quartz window. A rectangular waveguide is connected to a

cylindrical vessel. TE01 mode of microwave with a frequency of 2.45 GHz is excited at a cross

section of the waveguide[148]. Quartz is transparent to the microwave frequencies used there is

minimal interaction between the microwave power and the quartz window. The system used is

imaged in Figure 4.7 with the plasma ignited. Pressure ranges from 20 to 200 Torr depending on

the desired deposition conditions and the mixture of gases is typically CH4 in H2 with TMB used

as the boron dopant source.

Figure 4.7 Rayleigh–Bénard convection occurs in a plane horizontal layer of fluid heated from

below, in which the fluid develops a regular pattern of convection cells known as Bénard cells.

The plasma is made by electrons in the chamber absorbing the energy from the electric

component of the microwave field. These electrons then collide with gas phase species and

95

transfer their kinetic energy. The collisions result in the heating of the gas mixture and

dissociation, excitation, and ionization of gas phase species. Additional reactions that occur

between the various gas phase atoms, molecules, and ions results in a complex set of chemical

reactions occurring the chamber. In addition to all this diffusion of the gas phase species toward

the substrate occurs and they react with the surface allowing for growth of the diamond. A

detailed simulation of the gas chemistry and the plasmas temperature are shown in Figure 4.8.

Hence there there is a steep temperature gradient between the substrate and the region where

temperature is highest. The rise of the temperature causes the so-called Rayleigh–Benard

convection Figure 4.7. Rayleigh–Bénard convection is a type of natural convection, occurring in

a plane horizontal layer of fluid heated from below, in which the fluid develops a regular pattern

of convection cells known as Bénard cells.

Figure 4.8 (a) shows the gas temperature T in Kelvin for a comparable PECVD reactor. (b)

shows the hydrogen atomic mole fraction as a percentage for substrate holder with a diameter of

9mm and power density ~120Wcm-3

(c) shows the C2 and (d) the CH3 mole fraction expressed as

a percentage [12]

So inside the chamber the power density around the substrate edge is more intense than those in

the central region. This means that production rates of radical/excited species in the edges would

96

be higher than that in the central region. Flow direction points from the edges to the center above

the substrate. Hence, species are transferred from edges into the center. The temperature along

the edges on the top surface of the substrates is higher than that in the central region similar to

the power density. Hence, reactions of the surface with radical species on the edges are expected

to be more intense than that in the central region. Using the growth recipe of 400 sccm hydrogen

and 0.001-5% CH4 at 800oC at 40 Torr and 750 W [149, 150] a region where smooth boron

doped diamond can be repeatedly grown. These conductions allow a large number of hydrogen

atoms to form, these hydrogen atoms play an important role in the CVD process. In the most

ideal growth conditions methane concentration is kept to a minimum to ensure smooth surfaces

after the growth process. The gas chemistry is mostly made of various hydrocarbons as shown in

Figure 4.8. It’s these hydrocarbons that attach to the diamond seeds and cause the growth to

occur. In general the higher the temperature in the plasma more variety in hydrocarbon species

exits allowing for diamond growth, to low a temperature means the methane remains methane

and no growth happens as a result. Since it’s not possible to detail every growth recipie a general

environment that is favorable for growth has been established based on the ratios of hydrogen,

carbon, and oxygen known as the Bachmann triangle [151] shown in Figure 4.9. The ideal

conditions for growth require that the C:O ratio is as close to 1 as possible. During the growth

diamond surface is presented with a great deal of hydrogen atoms making the surface hydrogen

terminated. For the CH3 radical and other radicals to reach the surface the hydrogen bond must

be severed. Besides the gas chemistry an important factor in diamond growth is using the

appropriate orientation diamond substrate. Of the four possible orientations of (100), (110),

(311), and (111) the (100) surface is the most ideal for homoepitaxial growth and boron doping.

97

Figure 4.9 A simplified version of the Bachmann triangle showing the diamond growth region in

addition to regions where no growth and non-diamond growth occurs.

The (100) surface is ideal for growth as a result of its low defect density and because this surface

only requires one carbon atom on the surface to form part of a new layer [152]. The basic

incorporation process requires that a surface radical site be created, typically by the hydrogen

plasma, and methane radical CH3 is abstracted to the surface. This process of etching and adding

continues randomly on the surface until CH2 is incorporated into the crystal lattice.

An alternative to PECVD growth of diamond samples is to start off with naturally

occurring diamond and using heavily boron doped silicon nanomembranes and diffusion dope

the diamond samples. This process will be referred to diffusion doped natural diamond (DDND).

This process begins by using SOI and etching periodic holes through the box layer. Once this is

98

done the sample is placed in HF to allow the box layer to be wet etched. The etching allows the

top silicon in the SOI to separate. Once separated a PDMS stamp can be used to transfer the

silicon nanomembrane (SiNM) to the diamond. The SiNM diamond stack is them placed in a

rapid thermal annealing (RTA) furnace and diffusion annealed Figure 4.10 illustrates the process.

Both techniques have been developed and explored in the following sections.

Figure 4.10 Illustration of the SiNM preparation and diffusion process for diffusion doping of

single crystal <100> Ib diamond. i. Heavy boron implantation on an SOI wafer and thermal

annealing to realize heavily doped top Si on SOI. ii. Heavily boron doped top Si layer released

as SiNM by selective etching of SiO2 . iii. Top Si picked up by an elastomeric stamp. iv. SiNM

transferred to a diamond plate. v. Bond forming between SiNM and diamond and thermal

diffusion with RTA. vi. SiNM removed by potassium hydroxide (KOH) etching

Once the diamond has been grown several techniques were used to characterize the

quality of the diamond growth. Theses techniques include Raman, FTIR (Fourier transform

99

infrared spectroscopy), XPS (x-ray photoelectron spectroscopy), XRD (x-ray diffraction), and

SIMS (secondary ion mass spectroscopy). The quickest way to characterize the diamond samples

is by using Raman. Figure 4.11 shows raman data for three types of diamond samples.

Figure 4.11 Raman spectroscopy of three types of diamond, Green plot is natural Ib diamond,

Blue plot is synthetic high pressure high temperature (HPHT) diamond, and the red is synthetic

PECVD diamond. The blue dots are carbon while the white dots are hydrogen. If there is C-H

streaching the optical phonons will show up at ~3300cm-1

while if it’s only C-C stretching there

will be a strong peak at 1330cm-1

and another peak at 1550cm-1

.

Raman is also a great tool for determining if there is any damage done to the lattice. A

low energy region (400-1500 cm-1

) due to stretching of the C-C cage, and a high energy region

(2700-3100 cm-1

) due to C-H stretches and bends The fingerprint of diamond is a single sharp

100

Raman line at 1332 cm−1

. The more pure sp3 the sample is the sharper the peak (small full width

half maximum FWHM). Growing with large amounts of CH4 in the gas results in more sp2 bonds

forming at ~1550 cm-1

. PECVD synthetic diamond beyond the sp3 peak at 1330 you also see

peaks at 1550 and 3300. Both of these peaks are not ideal for device fabrication. The 1550 peak

is an indication of sp2 bonding in the diamond lattice, while the 3300 peak is from carbon

hydrogen bonds in the lattice. This happens because the sample is synthetic PECVD grown most

likely in a hydrogen dense environment. The HPHT synthetic diamond in blue shows only a

strong peak at 1330 and so does the NSCD (natural single crystal diamond) meaning the lattice is

nearly pure sp3 bonding. This makes the HPHT sample ideal for growth and electronic

applications which would be affected by the imperfect PECVD substrate. For boron doped

diamond boron concentration must exceed the Mott transition density and the raman signal will

begin to show weak signals around 610 cm−1

, 925 cm−1

, 1045 cm−1

, 1375 cm−1

, and 1470 cm−1

and a downshift zone bond center in the raman peak at wave numbers 500, 1225, 1230, 1320 cm

−1 and 1332 cm

−1 [153]. For the diamond samples grown the doping density was kept well below

the Mott transition therefore many of these mentioned changes in the raman signal do not appear.

The next characterization technique that is used is XRD (x-ray diffraction). With this

method one can learn about the crystal structure of the diamond substrates and characterize the

growth as well. The bulk is characterized by using transmission (Laue) geometry. The x-ray

diffractes in different directions and an image is formed from the wave fields interfering with

one another. The diffracted waves are observed on a position sensitive detector and the intensity

and shape of the diffracted spot is related to the quality of the crystal. In a perfect crystal

diffracted waves will produce Laue diffraction sports of uniform intensity. The position of the

Laue diffraction spot is given by the Bragg equation:

where d is the

101

spacing of the crystal planes and the angle theta is known as the Bragg angle. For single crystal

diamond the diffracted spot represents a spatial distribution of intensity due to the variation in

the lattice spacing. This allows for crystal defects such as dislocations, platelets, growth sections,

and stacking faults to be observed. There plane of interest is (100), excessive asymmetric shape

of the peak means that there are nitrogen platelets in the crystal. Diffusion streaks along the

along the 100 direction for 400 Bragg relation in boron doped crystal implies that boron atoms

likely precipitate on the 100 plane too [154]. Figure 4.12 compares the XRD patterns for

different types of diamond.

Figure 4.12 Comparison of three types of synthetic diamond. The first left image is synthetic

PECVD diamond, the middle left is synthetic PECVD diamond with nitrogen incorporation, the

middle right is boron doped PECVD grown on a synthetic PECVD substrate, and the right image

is a heavily boron doped synthetic diamond sample.

102

A figure of merit for the quality of the diamond is the FWHM of the 100 diffraction peak. For

the samples shown in Figure 4.12 FWHM were 0.0017deg and for the boron doped sample it was

0.0022deg. This means the crystal quality is good and close to the ideal lattice spacing for single

crystal diamond. One noticeable effect is once nitrogen or boron atoms are added to the diamond

films the XRD peak stretches along the 100 direction. It is most notable in the commercial

nitrogen doped sample. These results suggest that the boron atoms precipitate on the (100) plane

and expand the lattice planes. Overall samples were repeatedly grown and show consistency and

uniformity of the grown samples.

The next characterization technique that was used was SIMS. The principle behind SIMS

is high energy ions like Cs+ or O- are shot at the sample and the material gets sputtered of the

sample substrate. This then gets analyzed to determine the composition of the material. For these

experiments Cs beam was used to detect hydrogen, nitrogen, silicon, and carbon while an O2

beam was used to detect boron and molybdenum. Figure 4.13 shows several SIMS profiles.

103

Figure 4.13 SIMS profile for boron for (a) PECVD grown samples (b) and diffusion doped

sample (c)shows the profile for additional materials that get incorporated into the film during

growth which include Si , N , O, and H.

What is immediately observed is that the diffusion doping technique will only allow the boron to

diffuse ~50nm while the PECVD growth method can allow a variety of doping profiles and

depths. In addition to the boron incorporated into the diamond for hydrogen rich PECVD

growths a great deal of hydrogen is incorporated into the film. This results shows up as a Raman

peak at 3300 cm-1

and indicates that the film may have a high conductivity hydrogen layer on the

surface. This is significant because the holes that the boron is adding to the diamond film will be

104

compensated from the free electrons generate at the surface making the material seem metallic in

nature [155]. As deposited films prepared by chemical vapor deposition have a high conductivity

semiconducting layer near the surface which must be removed by oxidation in either

concentrated acid or oxygen ambient annealing. Hydrogen atoms terminate the dangling bonds

and passivate shallow and deep levels giving diamonds surface an negative electron affinity.

Looking at figure 4.5 one can see the band diagram for diamond with a negative electron affinity

(NEA) which is hydrogen surface terminated and a surface with positive electron affinity (PEA)

with is oxygen surface terminated. The behavior of the semiconductor is very different as a

result. For NEA the hydrogen makes the surface negatively charged and this induces a hole

accumulation layer near the surface causing band bending and high conductivity due to the hole

accumulation at the bent band. Another effect of the NEA is it exhibits little temperature

dependence of the hole concentration between 120K to 400K. If the diamond was oxidized

properly the film would have a strong temperature dependence with activation energies roughly

0.37eV.

105

Figure 4.14 Shows the effect on the XPS data as a reulst of the high conductively layer and a

fabricated device using the HCL as a diode at room temperature (green), 100oC (blue), and

200oC (red).

Figure 4.14 shows the results of a fabricated device using the HCL [156]. To verify that the

results were truly coming from the high conductivity layer XPS was performed on the surface of

the hydrogenated diamond. In this process x-rays interact with the surface material and ionize the

surface. Photoelectrons are generated from the core level and move through the surface. The

identification of chemical composition is done by measuring the energy of the photoelectron that

is released. This energy is called the binding energy as is the different between the initial and

final states of the atom at the surface. Depending on the material will deflect at a certain position

on a detector. The peak of interest is the 1s C peak at ~284 eV this is a good indicator of single

crystal diamond the broader the peak implies other forms of carbon like graphite may be present

in the grown film. This peak for natural diamond remains at ~284 eV but after the introduction of

the HCL the peak shifts to higher energies because of the negative electron affinity. Figure 4.14

shows the peak has shifted to ~287 eV. This verifies that the HCL layer is present and influences

the behavior at the surface of the diamond diode sample. This means that the rectification is

occurring in the top 20nm of the material and depleting this region and the minimal temperature

dependence also verifies this [157].

The next analysis technique used was determining if the growth recipes used produced

smooth usable surfaces. The ideal growth would leave the (100) surface with the 2x1

reconstruction. This type of surface termination has a crosshatch pattern on the surface as can be

seen in the profilomtery scanning images. Using a variety of growth conditions [149, 158-160]

that all have low methane concentrations of <0.15% and TMB concentrations <10000ppm

106

smooth boron doped diamond samples of growth rates ~100nm/hour can be repeatedly grown.

The profilometer scans are 0.5mm x 0.5mm with all of them having nearly 2nm RMS surface

roughness for samples that started off with RMS roughness >30nm. Not only were the samples

grown but they improved the final surface finish while doping. The growth process is a balance

between deposition and etching. Too low of a CH4 concentration etching will take over but to

high the surface will grow hillocks and other undesirable features. The key to smooth surface

growth (Rrms < 2nm ) a balance has to be found between the two competing processes. Further

introducing a dopant during the growth will influence where this ideal point is:

Acid Cleaning

1. Sodium Hydroxide [NaOH] : Hydrogen Peroxide[H2O2] (5:3) @ 60oC for

20min

2. Sulfuric Acid [H2SO4] :Nitric Acid [HNO3] : Perchloric Acid [HClO4]

(3:4:1) @ 250oC for 30min

3. Hydrochloric Acid [HCl]: Nitric Acid [HNO3] (6:1) @ 80oC for 30min

4. Ammonium Hydroxide [NH4OH]: Hydrogen Peroxide [H2O2]: H2O

(1:1:5) @ 75-85oC for 10min

5. Hydrofluoric acid [HF] : Nitric Acid [HNO3] (1:1) @ 20oC for 10min

6. Deionized Water (DI) 100oC for 10min

Surface Preparation in H2 plasma for 30 mins

Undoped Diamond Growth

Doped Diamond Growth (TMB, trimethylboron)

Post Undopded growth

Post H2 plamsa anneal

107

Anneal and Cool in H2 ambient.

Figure 4.15 Optical profilometry of the diamond samples showing very smooth (100) surfaces

with roughness RMS values <5nm.

Varying the concentration of CH4 and the TMB in the gas dramatically affects the smoothness of

the final surface. Another factor is doping efficiency, despite 1600ppm of TMB in the chamber

during the doped growth only a small fraction actually gets incorporated into the crystal. An

perfect tool to determine the doping efficiency in single crystal diamond is Fourier transform

infrared spectroscopy (FTIR) [161, 162]. The main peak observed in FTIR is the intrinsic 2-

phonon absorption bands at ~2000 cm-1

. Since the films being grown are very thin significant

color changes from the incorporation of boron is difficult to see by eye. For certain doping

108

ranges FTIR is an ideal tool to characterize the diamond and doping efficiency. Figure 4.16

shows the FTIR spectra from a variety of diamond samples. Besides the main peak at ~2000 cm-1

there is an additional peak at ~1290 cm-1

and on other at 3900-2650cm-1

known as the two

phonon absorption, one phonon absorption, and the three phonon absorption [163]. The

hydrogen-impurity-related absorptions are at 4495, 3107, 2786 and 1405 cm-1

while he weak

absorptions at 1330, 1172 and 1013 cm-1

are due to B-nitrogen complex [164-166]. Another

challenge is determine which of the incorporated atoms are electrically active and which are not

active. Uncompensated boron concentration is typically calculated using the area or the height of

the peak located around ~2800cm-1

when the concentration is low. For high concentration the

absorption in around 2800cm-1

it to high making measurement unreliable, to avoid this issue the

uncompensated boron concentration can be evaluated using another peak in the one phonon

region at 1290cm-1

using the following [162]:

where α is the absorption coefficient. For

the system being used the power densities of the plasma are ~40W/cm3 for such lower power

densities the doping efficiency was found to be very low [167]. Ideally higher power densities

>100W/cm3

are needed to increase the doping efficiency otherwise it remains <1%. What is clear

from the Figure 4.16 FTIR data is the starting FTIR spectra for different types of diamond is

different. The main feature different between the natural diamond and synthetic diamonds is the

lack of a strong one phonon absorption peak. This peak is also associated with nitrogen vacancy

complexes for synthetic diamonds this is not present until the nitrogen content in the diamond

films is much higher. For diamond films that don’t have an initial strong single phonon peak it’s

difficult to use the 1290 cm-1

to determine the doping concentration. Instead the SIMS technique

and other electrical characterization will be needed. This tells us a very interesting fact the rough

109

estimate of the vacancies present in the diamond films. Thie

Figure 4.16 FTIR spectra of several diamond samples. This compares the natural diamond to

the synthetic diamond spectra. For the natural diamonds the one phonon absorption peak as well

as the two phonon absorption peak is present. For the synthetic diamonds only the two phonon

absorption peak is present.

110

Using this knowledge of the presence of vacancies in the natural diamond lattice another

technique to develop diamond diodes was attempted. Using the SiNM technique as shown in

Figure 4.10 one can diffusion dope the natural diamond. The process can be described by

substitutional doping of boron atoms into the diamond surface. Boron atoms reside in Si lattice

sites and the SiNM is heavily boron doped across its entire thickness before thermal diffusion.

Boron atoms diffuse into diamond and replace some carbon atoms at the top region of diamond

after thermal diffusion. There is an exchence of the vacancies in the diamond with the dopant

SiNM that allows for the PI structure to form in the diamond. Figure 4.17 shows the results of

the diamond diode based on the SiNM doping technique. The XPS results show there is a

addition of Silicon and nitrogen into the top layer in significant amounts roughly ~4 atomic

percent. This means several materials have also diffused into the surface as well as the boron

which will affect the performance of diode devices. The IV data for the diode device is shown in

Figure 4.17 as well as shows large leakage current during reverse bias for very low voltages.

This implies that the rectification in the diamond sample is not ideal and the addition of Si and N

into the materials compensates and makes generation and recombination centers in the material

reducing its performance. It seems in the end the best method for diamond growth is growth in a

PECVD reactor using a doping gas on HPHT Ia substrates. The next section will detail the

growth of PECVD diamond for radiation hard applications.

111

Figure 4.17 Diffusion doped diamond diode with XPS data. Shows the the SiNM also diffuses

nitrogen and silicon in addition to the boron. Great deal of leakage current as a result of this.

The smaller peaks to the right ~105eV and ~160eV correspond to Si incorporation into the to

layers of the lattice from diffusion ~4% in the lattice. Also nitrogen is also incorporated at

~408eV.

4.4 Fabrication of PI Diodes for high sensitivity Temperature Sensors

Of the several methods available the PECVD growth technique has the most flexibility

for growing semiconducting diamond samples. The diffusion process though promising is

limited in the depth of doped material that can be grown also the addition of undesirable

elements in the films makes it non ideal for radiation applications where some materials are more

112

sensitive to irradiation. In Figure 4.18 the effect of radiation on Boron, the dopant in the diamond

behaves.

Figure 4.18 Boron has two naturally occurring and stable isotopes, 11

B (80.1%) and 10

B (19.9%)

- 10

B is used in boron neutron capture therapy. The carbon in diamond is nearly all 12

C

Lithium has two stable isotopes, 6Li (7.59%) and

7Li (92.4%) – the nuclear cross section of

6Li

940 barns while 7Li is 45mbarns [ 1 barn = 10

-28 m

2 ] making

7Li less affected by neutron

irradiation [KSU (P. Ugorowski) ]

Of all the materials in the periodic table two of the most radiation resistant are carbon 12 and

boron. Both these materials have a very large cross section meaning they are nearly transparent

to high energy radiation. If the diffusions process is used the addition of silicon and other

materials with smaller cross sections will introduce decay into other undesirable materials in the

region where rectification occurs degrading the long term performance of the device at high

radiation and high temperatures. In nature there are two isotopes of boron 11 and boron 10 and

113

they occur 80% and 20% respectively. For carbon there is carbon 12 and carbon 13 and in nature

nearly all carbon is carbon 12. So if the device can be made of pulley carbon 12 and boron atoms

the device will have a very minimal interaction with the irradiating source [168]. To test the

effect of irradiation on single crystal diamond samples two samples were placed into the UW

research reactor as shown is Figure 4.19 and irradiated for 15 minutes. One of the samples was

natural single crystal diamond and the second sample was chemical vapor deposited single

crystal diamond. The characteristic peak position of diamond is ~1300 cm-1

that has a full width

half maximum of ~2-3 cm-1

. The vibrational motion associated with this Raman signal involves

the stretching of the two atom basis with all unit cells moving in phase. The more pure sp3 the

sample is the sharper the peak. An illustration of this is shown in Figure 4.19.

114

Figure 4.19 The samples were irradiated with an average fast flux of ~ 2.63E+12 [n/cm2

s] and

a flux greater than 2.9eV of ~ 6.511E+11[n/cm2

s] for 15 minutes. This time attempts to replicate

the conditions the samples will experience during real operation

Figure 4.19 shows the Raman mapping results for the natural and single crystal diamond after

irradiation. Pre-irradiation Raman scans were also collected. The two maps show the peak

location for the characteristic C-C sp3 bond in diamond at ~1330 cm

-1. The magnitude of the full

width half maximum remains the same for both the samples. The two scale bars for Figure 4.19

are slightly different causing the difference in color. The constant color across the samples

indicates that the FWHM remains the same over very large regions of the diamond sample. The

width of the peak for the samples is ~6cm-1

making the FWHM ~3cm-1

, which is characteristic of

high quality single crystal diamond. The large neutron cross section of diamond and B10 helps

prevent significant interaction with the lattice after the 15 minutes of irradiation. No significant

changes in the structure of the single crystal diamond were observed. This is consistent with the

experimental data shown for diamond.

The next step for using the diamond as a temperature sensor is developing a reliable PI

junction. Schottky diodes have previously been demonstrated as viable temperature sensors but

are limited to lower temperatures up to 400oC [132]. The Schottky diodes suffer from reliability

issues as well as high leakage currents at elevated temperatures. PI junctions would eliminate

many of these issues and allow the junction to detect up to 800oC. Additionally using such a

wide band gap material like diamond allows for the device to be very sensitive. Sensitivity is a

figure of merit for temperature sensors. It determines how much of a voltage change occurs for a

change in degrees Celsius. The higher the sensitivity the quicker more detail temperature profiles

115

can be created. This will allow for more detail temperature measurement in extreme

environments. In Figure 4.20 the derivation of the sensitivity starting from the ideal diode

equation is shown and displays an exponential dependence. Even as the temperature increases

above 500oC diamond is able to maintain a high sensitivity while still being able to withstand the

high radiation in the environment.

2 lnln2

1 )1`(`

...

lnln

)ln()ln(

24`

`

1

C

1

0

0

0

2/3**

3

2

0

0

0

0

nifNNq

k

qWA

I

q

k

T

V

nifeq

TkBBTE

T

V

IIk

qVT

IIq

kTV

mmh

kB

eTBI

eII

Vea

kT

E

ga

a

a

he

kT

E

nkT

qV

g

g

a

Proc. Roy. Soc. (London) A277 (1964)

Figure 4.20 The left shows the ideal diode equation after some algebra extracting the sensitivity

which has the materials band gap in the exponent. The right plot shows how diamonds band gap

changes over a wide temperature range (<1%) meaning the sensitivity will stay the same even as

the environment changes [13].

To make the diamond samples a PECVD (Plasma Enhanced Chemical Deposition

System) system was used. The growth process is a balance between deposition and etching. Too

low of a CH4 concentration etching will take over but to high the surface will grow hillocks and

116

other undesirable features. The key to smooth surface growth (Rrms < 2nm ) a balance has to be

found between the two competing processes. The smooth surface is critical for additional device

processing. Further introducing a dopant during the growth will influence where this ideal point

is. Using this technique the diode structure can be grown repeatedly and reliably. In selecting a

design for the diode the effects of irradiation must be taken into account.

Figure 4.21 The structure of the PECVD grown diamond and the respective IV curve from the

devices. The IV shows little leakage current while having ideality factors close to one.

The energy required to produce one electron-hole pair is 13.2 eV, so the charge collected on the

electrodes will be increase as more energetic particles interact with it. To minimize this effect the

thickness of the diode is kept thin. The thinner the layer the less cross section is available for the

high energy photons to interact with the material, reducing the effect of irradiation creating extra

current. The current samples are <500nm in height which ensures that the effect is small enough

117

the measurement equipment tool will only be recording the effects of the temperature change and

not the changes in the flux of the reactor.

Figure 4.22 The proposed design for the capillary for insertion into the reactor. The diode will

be inside the capillary and placed next to the fuel and the two connectors will be treaded through

an insulating material.

One critical step in realizing this instrument is creating a holder to allow for reproducible

results. We plan on coordinating with the INL HTTL (Idaho National Lab High Temperature

Thermocouple Lab) group on developing a holder for the diamond sample. We are currently

working with the INL group to determine the most effective way to attach the lead wires and

package the device. Once this is completed out of pile test can continue at UW-Madison. By

using molybdenum wires and a capillary the device can be probed and remain small helping to

118

reduce the effects of irradiation and allowing for quick installation of the device into the out of

pile testing structure. Figure 4.22 shows the proposed plan.

4.5 Future Work & Summary

Future work includes characterizing the Current vs. Voltage (IV) properties of the

diamond diode samples post irradiation. This will provide another method for monitoring the

damage introduced to the lattice by the irradiation. The IV technique in addition to raman

mapping will be a convenient way to track the changes to the sample after irradiation and high

temperature measurements. Other future plans include refining the diode processing steps to

accommodate the test configurations for the UW and INL TREAT reactor. This may include

changing the metal pad size and increasing the doping levels to get stronger rectification from

the diodes.

119

Chapter 5 Electrical Artificial Human Eye Photo-detector Array

5.1 Introduction

The human eye is a remarkable feat of evolutionary engineering. This organ has inspired

many biomimicry projects for numerous applications and in a variety of areas [169]. Human eyes

can be conceptualized as single-lens optical system, and one planoconvex lens has a parabolic

focal plane, which means it cannot provide a large focus region on a planar array of light

detecting devices. Retina naturally adopts the curvilinear shape which approximates the focal

plane of lens such that human eyes have large view field and supreme capability of focusing,

which illuminates the shape design of photodetectors. Semiconductor photodiodes are common

optoelectronic devices used for commercial light detecting applications [170], but the concept of

human eyes is difficult to implement for photodiodes due to the intrinsically planar nature of

semiconductor fabrication techniques, such as patterning, etching, ion implantation, material

deposition and growth, and etc [171]. Here we report strategies to overcome these limitations and

implement them to fabricate a concave curvilinear photodiode array with single-crystalline

silicon. The approach fabricates photodiodes on semiconductor membrane transferred to a

designed planar flexible substrate, and then mechanically mounts the array onto a concave

fixture for final implementation. It simplifies the process flow of curvilinear semiconductor

devices by arranging the deformation mounting to the last step and thus preserving the feasibility

of most semiconductor fabrication techniques

120

5.2 Background of Artificial Human Eye

Electronic and optoelectronic semiconductor devices on complex curvilinear surfaces are

versatile in various areas due to their new degree of design freedom and biomimicry merits[172],

such as conformal flexible electronics [173-175], artificial eyes [176-179], and skins [180, 181].

However, fabricating devices on non-planar surfaces can be major challenge because of the

planar nature of most fabrication techniques when they were developed on planar wafer or plate

materials in semiconductor industry[182]. The available techniques for direct fabrication on non-

planar surface, such as soft lithography and mechanical molding[183-185] and lens-assisted

lithography[186], are consequently complicated and expensive, and imposing specific

requirements. Encouraged by the promising prosperity in non-planar devices, strategies are being

investigated to circumvent the limits set by non-planar surfaces and meanwhile utilize mature

semiconductor fabrication techniques as much as possible for economic consideration. One

straightforward strategy is to fabricate devices on planar substrate and then mechanically

transform them to required surfaces. Despite the simplicity of the concept, the implementation

requires comprehensive design and consideration to ensure acceptable yield. In this chapter, we

present our fabrication flow implementing this strategy.

The process flow of fabricating flexible silicon photodiode array is described next. The

flow starts from silicon-on-insulator (SOI) material on which ion-implantation and thermal

furnace annealing were performed to form p-n junction. Then an array of holes was made by

photolithography and reactive-ion etching (RIE) so that the buried oxide (BOX) can be etched

away by immersing the material in HF solution. Once the BOX was fully undercut, the top

silicon membrane sank to the handling silicon substrate. Membrane transfer was performed so

121

that the silicon membrane was attached to PI (polyimide) flexible substrate with SU-8 as

adhesive layer. Following active region pattern by RIE to isolate each pixel and metal (Ti/Au)

interconnection network conclude the fabrication. SU-8 was repeatedly used as dielectric and

passivation material due to its capability of insulation and absorbing stress. Microscope picture

of the doped membrane is shown in Figure 5.1. Because of the dopants, different colors are

observed in regions with different types of doping. Holes for undercutting are also in the picture.

Typical shapes of photodiode are marked in the picture, including regular hexagon and pentagon

and irregular pentagon. The idea of choosing these shapes in the design is from insect compound

eye that ensures array consisting of these shapes is able to cover a hemispherical surface without

any gap. Finished photodiodes are shown in Figure 5.1. Two metal layers were fabricated to

form interconnection and an SU-8/SiO2 layer was used as insulation.

Figure 5.1 (a) Microscope picture of the doped membrane with etching holes. Different colors

indicate two types of doping. Shapes of each photodiode are marked out. (b), Microscope picture

for the finished silicon photodiode. Two metal layers clearly form interdigitated connection.

122

5.3 Electrical Characterization and Image Acquisition

To demonstrate the functionality, this concave photodiode array was tested in an optical

system with only one planoconvex lens (from Edmunds Optics, 10mm diameter and 10mm focal

length), as shown in Figure 5.2. It also demonstrates that the focal plane of a single lens is

parabolic surface. The radius of curvature of the fixture, which also decides the shape of the

photodiode array and the PI, is designed to approximate the simulated lens focal plane so that the

area of the focused region will be maximized. Within this measurement setup, an object of

hollow “W” made by carving through a cardboard is illuminated by the expanded green laser

beam, and placed in front of the lens. The photodiode array on PCB is placed behind the lens at

the focusing position and connected to electrical measurement equipment. Due to the limited

number of photodiodes (276 photodiodes in total), the image acquired looks mosaic. So a simple

technique is utilized to improve resolution. The object “W” is placed on a rotary stage and

multiple images are acquired with different rotation angle, and then these images are rotated

back to a determined angle and overlapped to derive a refined image. The image shown in Figure

5.2c and d is such a refined image with acceptable resolution. 6 images are used with rotation

angles of 0, 12, 24, 36, 48, 60 degrees, and they are rotated back to 0 degree respectively and

overlapped. This technique is also used for the array on the fixture with larger curvature and it

not only improves the resolution but also eliminates the “blind spots”.

123

Figure 5.2 (a) shows the optical setup for image creation (b) shows the convave photo detector

array (c) and (d) show the collected image using Labview and Matlab to extract and process the

collected IV data from the pixels.

5.4 Summary

In conclusion, we demonstrate the procedure to fabricate photodiode array on concave

curvilinear surface and verify its functionality within a single-lens image system that has

practical applications such as endoscope. This procedure is compatible with planar

semiconductor fabrication techniques, so that not only photodiodes but all other semiconductor

devices, such as MOSFETs, BJTs, LEDs, etc, are available to integrate onto curvilinear surfaces.

124

On the other hand, little requirement is imposed on the curvature of the surface since imaging

techniques are available to eliminates “blind spots” and improves resolution. With these merits

of this procedure, an artificial eye device with similar size of human eye is realizable with a

tunable lens and controllable fixture. Generally, this procedure enlarges the device design

freedom and will make numerous ideas practical in various areas.

125

Chapter 6 Conclusion and Future Work

6.1 Conclusions

In this these three main topics were covered the first was graphene syntheses and the

applications the material can be used in. Some of which include transparent electrodes, bang gap

modification in bilayer graphene, and transistor devices. As the processing challenges of current

silicon technology continues alternative like graphene will have to be considered to continue

gaining performance enhancements and lower power operation. Besides this the application of

transparent electrodes allows for graphene to be used in another unrelated area giving it the

opportunity to expand beyond the device section and into the energy harvesting and display

technologies. The next material was diamond this material has the potential to change the

capabilities of next generation power electronic devices. As the proliferation of electric vehicles

increases and the development of a new generation of nuclear reactor starts materials that have

properties that are sutitable for multiple fields are become very attractive. Current growth

technology has found ways to grow diamond in large quantities and with a variety of doping

concentrations for use in a variety of application ranging from temperature sensors to high power

rectifiers. As the demands on current wide bandgap materials continues alternative like diamond

will continue to gain popularity and find its place in the family of useful semiconductors.

6.2 Future Work

Future work will include developing additional straining techniques for graphene and

other two dimensional films. The development of controllable strain will allow for materials like

bilayer graphene to begin being using for active electronic devices. The main challenge will be

126

ensuring that the band gaps remain uniform in time and throughout the surfaces to ensure

consistent device performance. Additional project will be creating optoelectronic devices that

can utilize the unique optical and electronic properties of graphene on the FIR and terahertz

regimes.

Future work for diamond devices will include developing high power transistors using

the PECVD doped diamond samples. In the coming years as the world looks to green energy to

replace traditional coal and natural gas nuclear energy is an ideal candidate. New reactors are

being designed with more safety features and smaller in size. One issue is developing a new

generation of monitoring tools to accommodate the new reactors different and new operating

regimes. Sensors made of synthetic diamond can be used for countaig high energy particles as

well as creating diodes for use as temperature sensors Beyond the radiation environment

diamond host of promising properties make diamond an ideal material for many high

performance applications. The future demand and the ease of growth of diamond make the future

of diamond very promising.

6.3 References

1. Freitag, M., Graphene: Trilayers unravelled. Nature Physics, 2011. 7(8): p. 596-597.

2. Rao, C.N.R. and A.K. Sood, Graphene: synthesis, properties, and phenomena. 2013:

John Wiley & Sons.

3. Henriksen, E., et al., Cyclotron resonance in bilayer graphene. Physical review letters,

2008. 100(8): p. 087403.

127

4. Zhang, Y., et al., Direct observation of a widely tunable bandgap in bilayer graphene.

nature, 2009. 459(7248): p. 820-823.

5. Choi, S.-M., S.-H. Jhi, and Y.-W. Son, Controlling energy gap of bilayer graphene by

strain. Nano letters, 2010. 10(9): p. 3486-3489.

6. Verberck, B., et al., Strain-induced band gaps in bilayer graphene. Physical Review B,

2012. 85(12): p. 125403.

7. Spieler, H., Radiation detectors and signal processing. Introduction: Vertex Detection in

High Energy Physics.-2001, Univ. Heidelberg, 2001: p. 28-34.

8. Willatzen, M., M. Cardona, and N. Christensen, Linear muffin-tin-orbital and k⋅ p

calculations of effective masses and band structure of semiconducting diamond. Physical

Review B, 1994. 50(24): p. 18054.

9. Neves, A. and M.H. Nazaré, Properties, growth and applications of diamond. 2001: IET.

10. Massarani, B., J. Bourgoin, and R. Chrenko, Hopping conduction in semiconducting

diamond. Physical Review B, 1978. 17(4): p. 1758.

11. Bundy, F.P., The P, T phase and reaction diagram for elemental carbon, 1979. Journal of

Geophysical Research: Solid Earth (1978–2012), 1980. 85(B12): p. 6930-6936.

12. Mankelevich, Y.A. and P. May, New insights into the mechanism of CVD diamond

growth: Single crystal diamond in MW PECVD reactors. Diamond and Related

Materials, 2008. 17(7): p. 1021-1028.

13. Evans, T. and P. James. A study of the transformation of diamond to graphite. in

Proceedings of the Royal Society of London A: Mathematical, Physical and Engineering

Sciences. 1964: The Royal Society.

128

14. Ellinger, F., et al. Review of advanced and beyond CMOS FET technologies for radio

frequency circuit design. in Microwave & Optoelectronics Conference (IMOC), 2011

SBMO/IEEE MTT-S International. 2011: IEEE.

15. Novoselov, K.S., et al., Electric field effect in atomically thin carbon films. science, 2004.

306(5696): p. 666-669.

16. Katsnelson, M.I., Graphene: carbon in two dimensions. Materials today, 2007. 10(1): p.

20-27.

17. Aurich, H., J. Trbovic, and C. Schonenberger, Fabrication and Electrical

Characterization of Graphene. 2007, Master thesis by Hagen Aurich University of Basel

Department of Physics.

18. Lee, J., et al., Ab initio study of adsorption properties of hazardous organic molecules on

graphene: Phenol, phenyl azide, and phenylnitrene. Chemical Physics Letters, 2015. 618:

p. 57-62.

19. McCann, E. and M. Koshino, The electronic properties of bilayer graphene. Reports on

Progress in Physics, 2013. 76(5): p. 056503.

20. Neto, A.C., et al., The electronic properties of graphene. Reviews of modern physics,

2009. 81(1): p. 109.

21. Novoselov, K., et al., Two-dimensional gas of massless Dirac fermions in graphene.

nature, 2005. 438(7065): p. 197-200.

22. Zhang, Y., et al., Experimental observation of the quantum Hall effect and Berry's phase

in graphene. nature, 2005. 438(7065): p. 201-204.

23. Bolotin, K.I., et al., Ultrahigh electron mobility in suspended graphene. Solid State

Communications, 2008. 146(9): p. 351-355.

129

24. Du, X., et al., Approaching ballistic transport in suspended graphene. Nature

nanotechnology, 2008. 3(8): p. 491-495.

25. Dean, C., et al., Boron nitride substrates for high-quality graphene electronics. Nature


26. Elias, D., et al., Dirac cones reshaped by interaction effects in suspended graphene.

Nature Physics, 2011. 7(9): p. 701-704.

27. Mayorov, A., et al., Interaction-driven spectrum reconstruction in bilayer graphene.

science, 2011. 333(6044): p. 860-863.

28. Chen, Z., et al., Graphene nano-ribbon electronics. Physica E: Low-dimensional Systems

and Nanostructures, 2007. 40(2): p. 228-232.

29. Han, M.Y., et al., Energy band-gap engineering of graphene nanoribbons. Physical

review letters, 2007. 98(20): p. 206805.

30. Li, X., et al., Chemically derived, ultrasmooth graphene nanoribbon semiconductors.

science, 2008. 319(5867): p. 1229-1232.

31. Wang, X., et al., Room-temperature all-semiconducting sub-10-nm graphene nanoribbon

field-effect transistors. Physical review letters, 2008. 100(20): p. 206803.

32. Ritter, K.A. and J.W. Lyding, The influence of edge structure on the electronic properties

of graphene quantum dots and nanoribbons. Nature materials, 2009. 8(3): p. 235-242.

33. Pedersen, T.G., et al., Graphene antidot lattices: designed defects and spin qubits.

Physical review letters, 2008. 100(13): p. 136804.

34. Bai, J., et al., Graphene nanomesh. Nature nanotechnology, 2010. 5(3): p. 190-194.

35. Yan, J.-A., L. Xian, and M. Chou, Structural and electronic properties of oxidized

graphene. Physical review letters, 2009. 103(8): p. 086802.

130

36. Gilje, S., et al., A chemical route to graphene for device applications. Nano letters, 2007.

7(11): p. 3394-3398.

37. Sofo, J.O., A.S. Chaudhari, and G.D. Barber, Graphane: a two-dimensional

hydrocarbon. Physical Review B, 2007. 75(15): p. 153401.

38. Balog, R., et al., Bandgap opening in graphene induced by patterned hydrogen

adsorption. Nature materials, 2010. 9(4): p. 315-319.

39. Bonaccorso, F., et al., Graphene photonics and optoelectronics. Nature photonics, 2010.

4(9): p. 611-622.

40. Bae, S., et al., Roll-to-roll production of 30-inch graphene films for transparent

electrodes. Nature nanotechnology, 2010. 5(8): p. 574-578.

41. Kim, K.S., et al., Large-scale pattern growth of graphene films for stretchable

transparent electrodes. nature, 2009. 457(7230): p. 706-710.

42. Berciaud, S., et al., Probing the intrinsic properties of exfoliated graphene: Raman

spectroscopy of free-standing monolayers. Nano letters, 2008. 9(1): p. 346-352.

43. Ferrari, A., et al., Raman spectrum of graphene and graphene layers. Physical review

letters, 2006. 97(18): p. 187401.

44. Graf, D., et al., Spatially resolved Raman spectroscopy of single-and few-layer graphene.

Nano letters, 2007. 7(2): p. 238-242.

45. Ferrari, A.C., Raman spectroscopy of graphene and graphite: disorder, electron–phonon

coupling, doping and nonadiabatic effects. Solid State Communications, 2007. 143(1): p.

47-57.

46. Das, A., et al., Monitoring dopants by Raman scattering in an electrochemically top-

gated graphene transistor. Nature nanotechnology, 2008. 3(4): p. 210-215.

131

47. Huang, M., et al., Phonon softening and crystallographic orientation of strained

graphene studied by Raman spectroscopy. Proceedings of the National Academy of

Sciences, 2009. 106(18): p. 7304-7308.

48. Frank, O., et al., Raman 2D-band splitting in graphene: theory and experiment. Acs

Nano, 2011. 5(3): p. 2231-2239.

49. Li, Z., et al., Band structure asymmetry of bilayer graphene revealed by infrared

spectroscopy. Physical review letters, 2009. 102(3): p. 037403.

50. Mak, K.F., et al., Measurement of the optical conductivity of graphene. Physical review

letters, 2008. 101(19): p. 196405.

51. Elias, D., et al., Control of graphene's properties by reversible hydrogenation: evidence

for graphane. science, 2009. 323(5914): p. 610-613.

52. Pereira, V.M., A.C. Neto, and N. Peres, Tight-binding approach to uniaxial strain in

graphene. Physical Review B, 2009. 80(4): p. 045401.

53. McCann, E. and V.I. Fal’ko, Landau-level degeneracy and quantum Hall effect in a

graphite bilayer. Physical review letters, 2006. 96(8): p. 086805.

54. Traversi, F., V. Russo, and R. Sordan, Integrated complementary graphene inverter.

arXiv preprint arXiv:0904.2745, 2009.

55. Li, S.-L., et al., Enhanced logic performance with semiconducting bilayer graphene

channels. Acs Nano, 2010. 5(1): p. 500-506.

56. Li, S.-L., et al., Low operating bias and matched input− output characteristics in

graphene logic inverters. Nano letters, 2010. 10(7): p. 2357-2362.

57. Harada, N., et al., A polarity-controllable graphene inverter. Applied Physics Letters,

2010. 96(1): p. 012102.

132

58. Li, S.L., et al., Complementary‐Like Graphene Logic Gates Controlled by Electrostatic

Doping. Small, 2011. 7(11): p. 1552-1556.

59. Geim, A.K. and K.S. Novoselov, The rise of graphene. Nature materials, 2007. 6(3): p.

183-191.

60. Berger, C., et al., Ultrathin epitaxial graphite: 2D electron gas properties and a route

toward graphene-based nanoelectronics. The Journal of Physical Chemistry B, 2004.

108(52): p. 19912-19916.

61. Berger, C., et al., Electronic confinement and coherence in patterned epitaxial graphene.

science, 2006. 312(5777): p. 1191-1196.

62. Reina, A., et al., Large area, few-layer graphene films on arbitrary substrates by

chemical vapor deposition. Nano letters, 2008. 9(1): p. 30-35.

63. Sutter, P.W., J.-I. Flege, and E.A. Sutter, Epitaxial graphene on ruthenium. Nature

materials, 2008. 7(5): p. 406-411.

64. Coraux, J., et al., Structural coherency of graphene on Ir (111). Nano letters, 2008. 8(2):

p. 565-570.

65. Li, X., et al., Large-area synthesis of high-quality and uniform graphene films on copper

foils. science, 2009. 324(5932): p. 1312-1314.

66. Bhaviripudi, S., et al., Role of kinetic factors in chemical vapor deposition synthesis of

uniform large area graphene using copper catalyst. Nano letters, 2010. 10(10): p. 4128-

4133.

67. Li, X., et al., Graphene films with large domain size by a two-step chemical vapor

deposition process. Nano letters, 2010. 10(11): p. 4328-4334.

133

68. Borgström, M.T., et al., Synergetic nanowire growth. Nature nanotechnology, 2007. 2(9):

p. 541-544.

69. Zhang, W., et al., First-principles thermodynamics of graphene growth on Cu surfaces.

The Journal of Physical Chemistry C, 2011. 115(36): p. 17782-17787.

70. Wu, P., et al., Mechanisms of Graphene Growth on Metal Surfaces: Theoretical

Perspectives. Small, 2014. 10(11): p. 2136-2150.

71. Rodat, S., et al., Kinetic modelling of methane decomposition in a tubular solar reactor.

Chemical Engineering Journal, 2009. 146(1): p. 120-127.

72. Mi, X., et al., Facet-insensitive graphene growth on copper. Physical Review B, 2012.

85(15): p. 155436.

73. Smith, D.L., et al., Mechanism of SiN x H y Deposition from NH 3‐SiH4 Plasma. Journal

of the Electrochemical Society, 1990. 137(2): p. 614-623.

74. Noskov, A., et al., Correlation between stress and structure in chemically vapour

deposited silicon nitride films. Thin Solid Films, 1988. 162: p. 129-143.

75. Lennard-Jones, J., On the forces between atoms and ions. Proceedings of the Royal

Society of London. Series A, Containing Papers of a Mathematical and Physical

Character, 1925: p. 584-597.

76. Liu, Z., et al., Interlayer shear strength of single crystalline graphite. Acta Mechanica

Sinica, 2012. 28(4): p. 978-982.

77. Nagashio, K., et al., Electrical transport properties of graphene on SiO2 with specific

surface structures. Journal of Applied Physics, 2011. 110(2): p. 024513.

78. Cancado, L., et al., Influence of the atomic structure on the Raman spectra of graphite

edges. Physical review letters, 2004. 93(24): p. 247401.

134

79. Casiraghi, C., et al., Raman fingerprint of charged impurities in graphene. Applied

Physics Letters, 2007. 91(23): p. 233108.

80. Balandin, A.A., et al., Superior thermal conductivity of single-layer graphene. Nano

letters, 2008. 8(3): p. 902-907.

81. Mohiuddin, T., et al., Uniaxial strain in graphene by Raman spectroscopy: G peak

splitting, Grüneisen parameters, and sample orientation. Physical Review B, 2009.

79(20): p. 205433.

82. Robinson, J.A., et al., Raman topography and strain uniformity of large-area epitaxial

graphene. Nano letters, 2009. 9(3): p. 964-968.

83. Yoon, D., Y.-W. Son, and H. Cheong, Strain-dependent splitting of the double-resonance

Raman scattering band in graphene. Physical review letters, 2011. 106(15): p. 155502.

84. Grimvall, G., Thermophysical properties of materials. 1999: Elsevier.

85. Frank, O., et al., Compression behavior of single-layer graphenes. Acs Nano, 2010. 4(6):

p. 3131-3138.

86. Pan, W., et al., Biaxial compressive strain engineering in graphene/boron nitride

heterostructures. Scientific reports, 2012. 2.

87. Zabel, J., et al., Raman spectroscopy of graphene and bilayer under biaxial strain:

bubbles and balloons. Nano letters, 2012. 12(2): p. 617-621.

88. Lee, C., et al., Measurement of the elastic properties and intrinsic strength of monolayer

graphene. science, 2008. 321(5887): p. 385-388.

89. Koenig, S.P., et al., Ultrastrong adhesion of graphene membranes. Nature


135

90. Ni, Z.H., et al., Uniaxial strain on graphene: Raman spectroscopy study and band-gap

opening. Acs Nano, 2008. 2(11): p. 2301-2305.

91. Kuzmenko, A., et al., Determination of the gate-tunable band gap and tight-binding

parameters in bilayer graphene using infrared spectroscopy. Physical Review B, 2009.

80(16): p. 165406.

92. Blakslee, O., et al., Elastic constants of compression‐annealed pyrolytic graphite. Journal

of Applied Physics, 1970. 41(8): p. 3373-3382.

93. Ding, X., et al., Measurement of shear strength for HOPG with scanning tunneling

microscopy by thermal excitation method. Ultramicroscopy, 2012. 115: p. 1-6.

94. Guo, Y., W. Guo, and C. Chen, Modifying atomic-scale friction between two graphene

sheets: A molecular-force-field study. Physical Review B, 2007. 76(15): p. 155429.

95. Zhang, L., et al., Determination of the electronic structure of bilayer graphene from

infrared spectroscopy. Physical Review B, 2008. 78(23): p. 235408.

96. Nicol, E. and J. Carbotte, Optical conductivity of bilayer graphene with and without an

asymmetry gap. arXiv preprint arXiv:0801.1836, 2008.

97. Kuzmenko, A., et al., Infrared spectroscopy of electronic bands in bilayer graphene.

Physical Review B, 2009. 79(11): p. 115441.

98. Mathkar, A., et al., Controlled, stepwise reduction and band gap manipulation of

graphene oxide. The Journal of Physical Chemistry Letters, 2012. 3(8): p. 986-991.

99. Ci, L., et al., Atomic layers of hybridized boron nitride and graphene domains. Nature

materials, 2010. 9(5): p. 430-435.

100. Tauc, J., Optical properties and electronic structure of amorphous Ge and Si. Materials

Research Bulletin, 1968. 3(1): p. 37-46.

136

101. Metzger, C., et al., Biaxial strain in graphene adhered to shallow depressions. Nano

letters, 2009. 10(1): p. 6-10.

102. Ding, F., et al., Stretchable graphene: a close look at fundamental parameters through

biaxial straining. Nano letters, 2010. 10(9): p. 3453-3458.

103. Gui, G., et al., Local strain effect on the band gap engineering of graphene by a first-

principles study. Applied Physics Letters, 2015. 106(5): p. 053113.

104. Shioya, H., et al., Straining graphene using thin film shrinkage methods. Nano letters,

2014. 14(3): p. 1158-1163.

105. Carozo, V., et al., Raman signature of graphene superlattices. Nano letters, 2011. 11(11):

p. 4527-4534.

106. Lu, C.-C., et al., Twisting bilayer graphene superlattices. Acs Nano, 2013. 7(3): p. 2587-

2594.

107. Cadelano, E., et al., Nonlinear elasticity of monolayer graphene. Physical review letters,

2009. 102(23): p. 235502.

108. Garaj, S., W. Hubbard, and J. Golovchenko, Graphene synthesis by ion implantation.

Applied Physics Letters, 2010. 97(18): p. 183103.

109. Rius, G., M. Yoshimura, and N. Mestres, Synthesis of patterned nanographene on

insulators from focused ion beam induced deposition of carbon. Journal of Vacuum

Science & Technology B, 2012. 30(3): p. 03D113.

110. Baraton, L., et al., Synthesis of few-layered graphene by ion implantation of carbon in

nickel thin films. Nanotechnology, 2011. 22(8): p. 085601.

111. Gutierrez, G., et al., Multi-layer graphene obtained by high temperature carbon

implantation into nickel films. Carbon, 2014. 66: p. 1-10.

137

112. Mi, H., et al., Creating periodic local strain in monolayer graphene with nanopillars

patterned by self-assembled block copolymer. Applied Physics Letters, 2015. 107(14): p.

143107.

113. Fenno, L., O. Yizhar, and K. Deisseroth, The development and application of

optogenetics. Annual review of neuroscience, 2011. 34: p. 389-412.

114. Thongpang, S., et al., A micro-electrocorticography platform and deployment strategies

for chronic BCI applications. Clinical EEG and neuroscience, 2011. 42(4): p. 259-265.

115. Henle, C., et al., First long term in vivo study on subdurally implanted micro-ECoG

electrodes, manufactured with a novel laser technology. Biomedical microdevices, 2011.

13(1): p. 59-68.

116. Viventi, J., et al., Flexible, foldable, actively multiplexed, high-density electrode array for

mapping brain activity in vivo. Nature neuroscience, 2011. 14(12): p. 1599-1605.

117. Rubehn, B., et al., A MEMS-based flexible multichannel ECoG-electrode array. Journal

of neural engineering, 2009. 6(3): p. 036003.

118. Richner, T.J., et al., Optogenetic micro-electrocorticography for modulating and

localizing cerebral cortex activity. Journal of neural engineering, 2014. 11(1): p. 016010.

119. Schendel, A.A., et al., A cranial window imaging method for monitoring vascular growth

around chronically implanted micro-ECoG devices. Journal of neuroscience methods,

2013. 218(1): p. 121-130.

120. Ledochowitsch, P., et al. A transparent μECoG array for simultaneous recording and

optogenetic stimulation. in Engineering in Medicine and Biology Society, EMBC, 2011

Annual International Conference of the IEEE. 2011: IEEE.

138

121. Minami, T., Substitution of transparent conducting oxide thin films for indium tin oxide

transparent electrode applications. Thin Solid Films, 2008. 516(7): p. 1314-1321.

122. Noh, H.-s., et al., Wafer bonding using microwave heating of parylene intermediate

layers. Journal of micromechanics and microengineering, 2004. 14(4): p. 625.

123. Granqvist, C. and A. Hultåker, Transparent and conducting ITO films: new developments

and applications. Thin Solid Films, 2002. 411(1): p. 1-5.

124. Guillen, C. and J. Herrero, Comparison study of ITO thin films deposited by sputtering at

room temperature onto polymer and glass substrates. Thin Solid Films, 2005. 480: p.

129-132.

125. Li, X., et al., Transfer of large-area graphene films for high-performance transparent

conductive electrodes. Nano letters, 2009. 9(12): p. 4359-4363.

126. Chen, H., et al., Mechanically strong, electrically conductive, and biocompatible

graphene paper. Adv. Mater, 2008. 20(18): p. 3557-3561.

127. Li, N., et al., The promotion of neurite sprouting and outgrowth of mouse hippocampal

cells in culture by graphene substrates. Biomaterials, 2011. 32(35): p. 9374-9382.

128. Williams, J.C., et al., Complex impedance spectroscopy for monitoring tissue responses

to inserted neural implants. Journal of neural engineering, 2007. 4(4): p. 410.

129. Norlin, P., et al., A 32-site neural recording probe fabricated by DRIE of SOI substrates.

Journal of micromechanics and microengineering, 2002. 12(4): p. 414.

130. Gross, G.W., Simultaneous single unit recording in vitro with a photoetched laser

deinsulated gold multimicroelectrode surface. Biomedical Engineering, IEEE

Transactions on, 1979(5): p. 273-279.

139

131. Yuen, T. and W. Agnew, Histological evaluation of polyesterimide-insulated gold wires

in brain. Biomaterials, 1995. 16(12): p. 951-956.

132. Zhang, N., et al., Temperature sensor based on 4H-silicon carbide pn diode operational

from 20 C to 600 C. Applied Physics Letters, 2014. 104(7): p. 073504.

133. Meier, D., CVD diamond sensors for particle detection and tracking. 1999, CERN

Geneva (Switzerland).

134. Kittel, C., Introduction to solid state physics. 2005: Wiley.

135. Alianelli, L., et al., A planar refractive x-ray lens made of nanocrystalline diamond.

Journal of Applied Physics, 2010. 108(12): p. 123107.

136. Mildren, R. and A. Sabella, Highly efficient diamond Raman laser. Optics letters, 2009.

34(18): p. 2811-2813.

137. May, P.W., Diamond thin films: a 21st-century material. Philosophical Transactions of

the Royal Society of London A: Mathematical, Physical and Engineering Sciences, 2000.

358(1766): p. 473-495.

138. Porro, S., et al., Nanocrystalline diamond coating of fusion plasma facing components.

Diamond and Related Materials, 2009. 18(5): p. 740-744.

139. Nebel, C.E., et al., Diamond for bio-sensor applications. Journal of Physics D: Applied

Physics, 2007. 40(20): p. 6443.

140. Bergonzo, P., et al., 3D shaped mechanically flexible diamond microelectrode arrays for

eye implant applications: The MEDINAS project. Irbm, 2011. 32(2): p. 91-94.

141. Clark, C., P. Dean, and P. Harris. Intrinsic edge absorption in diamond. in Proceedings of

the Royal Society of London A: Mathematical, Physical and Engineering Sciences. 1964:

The Royal Society.

140

142. Sokoloff, J., Electron localization in crystals with quasiperiodic lattice potentials.

Physical Review B, 1980. 22(12): p. 5823.

143. Calzaferri, G. and R. Rytz, The band structure of diamond. The Journal of Physical

Chemistry, 1996. 100(26): p. 11122-11124.

144. Thonke, K., The boron acceptor in diamond. Semiconductor science and technology,

2003. 18(3): p. S20.

145. Inushima, T., et al., Hopping conduction via the excited states of boron in p-type

diamond. Diamond and Related Materials, 2000. 9(3): p. 1066-1070.

146. Bovenkerk, H., et al., Preparation of diamond. 1959.

147. Ece, M., B. Oral, and J. Patscheider, Nucleation and growth of diamond films on Mo and

Cu substrates. Diamond and Related Materials, 1996. 5(3): p. 211-216.

148. Yamada, H., et al., Numerical analyses of a microwave plasma chemical vapor

deposition reactor for thick diamond syntheses. Diamond and Related Materials, 2006.

15(9): p. 1389-1394.

149. Watanabe, H., et al., Homoepitaxial diamond film with an atomically flat surface over a

large area. Diamond and Related Materials, 1999. 8(7): p. 1272-1276.

150. Okushi, H., et al., Device-grade homoepitaxial diamond film growth. Journal of crystal

growth, 2002. 237: p. 1269-1276.

151. Bachmann, P.K., D. Leers, and D.U. Wiechert, Diamond Thin Films: Preparation,

Characterization and Selected Applications Progress Report. Berichte der

Bunsengesellschaft für physikalische Chemie, 1991. 95(11): p. 1390-1400.

152. Goodwin, D., et al., Handbook of industrial diamonds and diamond films. editor Prelas

MA, New York, Dekker, 1998.

141

153. Mermoux, M., et al., Raman characterization of boron-doped 111 homoepitaxial

diamond layers. Diamond and Related Materials, 2006. 15(4): p. 572-576.

154. Maeta, H., et al., The characterization of synthetic and natural single crystal diamonds by

x-ray diffraction. Physica B: Condensed Matter, 2006. 376: p. 283-287.

155. Hayashi, K., et al., Investigation of the effect of hydrogen on electrical and optical

properties in chemical vapor deposited on homoepitaxial diamond films. Journal of

Applied Physics, 1997. 81(2): p. 744-753.

156. Yamanaka, S., et al., Electrical conduction of high-conductivity layers near the surfaces

in hydrogenated homoepitaxial diamond films. Applied surface science, 2000. 159: p.

567-571.

157. Tsubota, T., et al., Surface morphology and electrical properties of boron-doped diamond

films synthesized by microwave-assisted chemical vapor deposition using trimethylboron

on diamond (100) substrate. Diamond and Related Materials, 2000. 9(7): p. 1362-1368.

158. Tokuda, N., et al., Hillock-free heavily boron-doped homoepitaxial diamond films on

misoriented (001) substrates. Japanese journal of applied physics, 2007. 46(4R): p. 1469.

159. Takeuchi, D., et al., Device Grade B‐Doped Homoepitaxial Diamond Thin Films. physica

status solidi (a), 2001. 186(2): p. 269-280.

160. Takeuchi, D., et al., High quality homoepitaxial diamond thin film synthesis with high

growth rate by a two-step growth method. Diamond and Related Materials, 1999. 8(6): p.

1046-1049.

161. Tallaire, A., et al., Growth of large size diamond single crystals by plasma assisted

chemical vapour deposition: Recent achievements and remaining challenges. Comptes

Rendus Physique, 2013. 14(2): p. 169-184.

142

162. Issaoui, R., et al., Evaluation of freestanding boron-doped diamond grown by chemical

vapour deposition as substrates for vertical power electronic devices. Applied Physics

Letters, 2012. 100(12): p. 122109.

163. Field, J.E., The properties of natural and synthetic diamond. 1992: Academic Press

London.

164. Thongnopkun, P. and S. Ekgasit, FTIR Spectra of faceted diamonds and diamond

simulants. Diamond and Related Materials, 2005. 14(10): p. 1592-1599.

165. Jenkins, A.L. and R.A. Larsen, Gemstone identification using Raman spectroscopy.

disclosure, 2004. 7: p. 9.

166. Iakoubovskii, K. and G. Adriaenssens, Optical characterization of natural Argyle

diamonds. Diamond and Related Materials, 2002. 11(1): p. 125-131.

167. Issaoui, R., et al., Growth of thick heavily boron-doped diamond single crystals: Effect of

microwave power density. Applied Physics Letters, 2010. 97(18): p. 182101.

168. Tsubota, M., et al., High-temperature characteristics of charge collection efficiency using

single CVD diamond detectors. Nuclear Instruments and Methods in Physics Research

Section A: Accelerators, Spectrometers, Detectors and Associated Equipment, 2015. 789:

p. 50-56.

169. Atchison, D.A. and G. Smith, Optics of the human eye. 2000.

170. Sze, S.M. and K.K. Ng, Physics of semiconductor devices. 2006: John Wiley & Sons.

171. Chen, W.-K., VLSI technology. Vol. 8. 2003: CRC Press.

172. Rim, S.-B., et al., The optical advantages of curved focal plane arrays. Opt. Express,

2008. 16(7): p. 4965-4971.

143

173. Kim, D.-H., et al., Stretchable and foldable silicon integrated circuits. science, 2008.

320(5875): p. 507-511.

174. Kim, D.H. and J.A. Rogers, Stretchable electronics: materials strategies and devices.

Advanced Materials, 2008. 20(24): p. 4887-4892.

175. Yuan, H.-C., et al., Flexible photodetectors on plastic substrates by use of printing

transferred single-crystal germanium membranes. Applied Physics Letters, 2009. 94(1):

p. 13102.

176. Ko, H.C., et al., A hemispherical electronic eye camera based on compressible silicon

optoelectronics. nature, 2008. 454(7205): p. 748-753.

177. Jung, I., et al., Paraboloid electronic eye cameras using deformable arrays of

photodetectors in hexagonal mesh layouts. Applied Physics Letters, 2010. 96(2): p.

021110.

178. Jeong, K.-H., J. Kim, and L.P. Lee, Biologically inspired artificial compound eyes.

science, 2006. 312(5773): p. 557-561.

179. Lee, L.P. and R. Szema, Inspirations from biological optics for advanced photonic

systems. science, 2005. 310(5751): p. 1148-1150.

180. Mannsfeld, S.C., et al., Highly sensitive flexible pressure sensors with microstructured

rubber dielectric layers. Nature materials, 2010. 9(10): p. 859-864.

181. Takei, K., et al., Nanowire active-matrix circuitry for low-voltage macroscale artificial

skin. Nature materials, 2010. 9(10): p. 821-826.

182. Chen, W.-K., VLSI technology 2003: CRC Press.

144

183. Heung Cho, K., et al., A hemispherical electronic eye camera based on compressible

silicon optoelectronics. Nature, 2008. 454(Copyright 2008, The Institution of

Engineering and Technology): p. 748-53.

184. Dinyari, R., et al., Curving monolithic silicon for nonplanar focal plane array

applications. Applied Physics Letters, 2008. 92(9): p. 091114-3.

185. Jin, H.-C., et al., Soft lithographic fabrication of an image sensor array on a curved

substrate. Journal of Vacuum Science & Technology B: Microelectronics and Nanometer

Structures, 2004. 22(5): p. 2548-2551.

186. Jeong, K.-H., J. Kim, and L.P. Lee, Biologically inspired artificial compound eyes.

Science, 2006. 312(Compendex): p. 557-561.

carbon material growth, characterization, and device

Documents