PLASMA APPLICATION FOR SYNTHESIS OF NOVEL NANOSTRUCTURES AND

MECHANICAL PROPERTIES INVESTIGATION

By

Alborz Izadi

A DISSERTATION

Michigan State University

in partial fulﬁllment of the requirements

Submitted to

for the degree of

Mechanical Engineering – Doctor of Philosophy

2020

PLASMA APPLICATION FOR SYNTHESIS OF NOVEL NANOSTRUCTURES AND

MECHANICAL PROPERTIES INVESTIGATION

ABSTRACT

By

Alborz Izadi

Today, nanomaterials are receiving increasing attention, as they exhibit exciting and useful at-
tributes that can be employed across applications and disciplines range from electronics to drug
delivery. This dissertation focuses on two disparate projects, both related to plasma-synthesized
nanocrystals. The two projects highlight the unique properties of nanocrystals, and their sustain-
able synthesis using gas phase approaches. These methods apply non-thermal plasma reactors for
nanocrystal synthesis. First, we introduce a plasma-based gas phase method for Gold nanoparticle
(AuNPs) synthesis. Second, we use plasma-synthesized Silicon nanocrystals (SiNCs) in bilayer
and composite structures with a commonly used elastomer, to investigate the mechanical behavior
of the structures in a ﬁrst-of-its kind investigation.

There are many applications for AuNPs due to their interesting optoelectronic properties, such
as tunable optical absorption and plasmonic resonance behavior. While synthesis and stabilization
of colloidal AuNPs is well-established, new synthesis routes can lead to enhanced versatility
of applications for AuNPs, particularly if the methods allow avoidance of solution processes or
surfactants. In Chapter 2, we introduce a plasma-based synthesis of AuNPs, using a consumable
gold wire and a radiofrequency power source. The AuNPs are monodisperse, with an average
diameter of 4 nm. While production yield is low, the narrow size distribution of the AuNPs and the
avoidance of solution processing in this method are promising for future syntheses of metal NPs
based on plasmas.

Next, a comprehensive analysis of the mechanical behavior of SiNC/PDMS systems, using
plasma-produced SiNCs has been performed. Chapter 3 details our experimental methods com-
bined with modeling to estimate the mechanical behavior of thin layers of SiNCs on PDMS, as
deposited directly onto the PDMS from non-thermal plasma reactor. For the ﬁrst time we estimated
the mechanical behavior of thin ﬁlms of SiNCs by using the onset of bifurcations as an indicator of

their modulus. Next, reaching towards luminescent nanocomposites for applications in luminescent
devices, we investigated the optical and mechanical behavior of blended SiNC/PDMS nanocom-
posites. The results from these investigations, reported in Chapter 3 and Chapter 4, have shed light
for the ﬁrst time on the interactions between SiNCs and PDMS both in bilayer and composite struc-
tures, pointing to future optoelectronic and opto-mechanical device applications with predictable
properties. Finally, in Chapter 5, we share ongoing projects which will be ﬁnalized soon, as well as
detailing future work surrounding these ideas. We share our parametric study to uncover the various
e ects of surface functionality, SiNC layer thickness, and PDMS modulus on the resulting SiNC
thin ﬁlm. Concurrently, we probed the formation of wrinkles and cracks on SiNC ﬁlm surfaces,
which were deposited on pre-stretched PDMS using a ﬁnite bending conﬁguration, to examine how
instabilities on the thin ﬁlms can be predicted.

Copyright by
ALBORZ IZADI
2020

To my mother, Shahla, my father, Hossein, and my sister, Dena

v

ACKNOWLEDGEMENTS

I would like to thank all the people who contributed in some way to the work described in this thesis.
First of all, I would like to express my sincere gratitude to my advisor, the most wonderful woman I
have met in the U.S., professor Rebecca Anthony, for the continuous support of my PhD research, her
patience, enthusiasm, and knowledge . When I joined her research group in 2014, I had absolutely
no experience in research. Without her guidance and support, ﬁnishing this dissertation would not
have been possible. In addition, I thank our collaborator, professor Sara Roccabianca. Without
her e orts my job would have undoubtedly been more di cult. I am thankful to my committee
members, professor Anctil and professor Yeom, whose comments have always been helpful. Every
result described in my thesis was accomplished with the help and support of fellow lab mates and
collaborators from Plasma and Nanomaterial Laboratory (PNL). I am especially thankful to Kyle
O’Shea, who was my closest friend during the last three years of graduate school, Alex Ho, the
most reliable lab-mate during my research , Dr. Gabriel Ceriotti, who always had useful ideas
and suggestions for me, Geena Du , Brant Toback, Michael Bigelow and many others who are all
my best friends. I am very grateful for the modeling data shared with me from Dr. Roccabianc’s
lab. I also would like to thank all members of Mechanical Engineering, Chemical Engineering and
Material Sciences Departments at MSU, specially Dr. Askeland, professor Naguib and professor
Wichman. I thank all the journals which gave me the permission to include copyrighted pictures
as part of my dissertation. I am also thankful for the funding sources that allowed me to pursue my
graduate school studies including National Science Foundation. I am thankful to all the members
of Advanced Microscopy Center Facility at MSU, who helped and mentored me through my sample
preparation process. I would like to acknowledge all my friends, professor Babak Borhan, Dr. Reza
Loloee, professor Maghrebi, Dr. Mehdi Nasr, Mohammad, Fariborz and Alireza who supported
me during my stay at MSU and made my time a lot more fun.

Finally, I would like to thank my mom, Shahla, my dad, Hossein, and my sister, Dena, for their
constant love and support. My deepest appreciation goes to my parents, who dedicated their own

vi

time and energy to support me in all aspects of my career and acadamic desires. Their video calls
always motivated me to stay strong.
I have a wish to always have their back and give my best
support to them. I will never forget their sacriﬁces in tough moments of my journey. Without their
support, help and dedication, it would be absolutely impossible for me to pursue graduate studies
abroad and be where I am today. Another deep appreciation from the bottom of my heart goes to
my sister, Dena, who supported me through moments of hopelessness and played the role of my
family during my stay in East Lansing.

vii

TABLE OF CONTENTS

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x

xi

1.1
1.2 Gold Nanoparticles .

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INTRODUCTION .

LIST OF TABLES .
LIST OF FIGURES .
CHAPTER 1

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
Introduction to Non-thermal Plasma . . . . . . . . . . . . . . . . . . . . . . . . .
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
1.2.1 History of AuNP synthesis . . . . . . . . . . . . . . . . . . . . . . . . . .
1.2.2 AuNP formation in the gas phase . . . . . . . . . . . . . . . . . . . . . . .
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1
1
4
5
5
8
1.3.1
SiNCs on Flexible Substrates: Polydimethylsiloxane (PDMS) . . . . . . . . 11
1.3.2 Mechanical properties of SiNCs, deposited as thin ﬁlms . . . . . . . . . . . 12
CHAPTER 2 AUNP SYNTHESIS: A PLASMA-BASED APPROACH . . . . . . . . . . . 14
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15
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. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 27

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2.1
2.2 Experimental Details
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2.3 Results and Discussion .
2.4 Conclusion .
.

1.3 Silicon Nanoparticles .

Introduction .

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CHAPTER 3 ESTIMATING THE MECHANICAL PROPERTIES OF SINC LAYERS

3.1
3.2 Methods .

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Introduction .
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ON PDMS .
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3.2.1
3.2.2
3.2.3
3.2.4 Uniaxial tensile test .
3.2.5
3.2.6 Modeling bifurcation in elastic layered structures

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 30
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 32
PDMS fabrication . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33
Plasma synthesis of SiNCs . . . . . . . . . . . . . . . . . . . . . . . . . . 33
Silicon nanocrystal ﬁlm characterization . . . . . . . . . . . . . . . . . . . 34
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36
. . . . . . . . . . . . . . . . . . . . . . . . . . 37
. . . . . . . . . . . . . . 38
Finite bending . . . . . . . . . . . . . . . . . . . . . . . . . . . 39
. . . . . . . . . . . . . . . . . . . . . . . . 40
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 41
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45

3.2.6.1
3.2.6.2 Bifurcation analysis

Finite bending experiment

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3.3 Results and Discussion .
3.4 Conclusion .
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CHAPTER 4 OPTICAL AND MECHANICAL BEHAVIOR OF SINC-EMBEDDED

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Introduction .

PDMS SLABS .
4.1
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4.2 Experimental Work .

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 49
.
Plasma synthesis of SiNCs . . . . . . . . . . . . . . . . . . . . . . . . . . 49
4.2.1
Functionalization of SiNCs . . . . . . . . . . . . . . . . . . . . . . . . . . 50
4.2.2
4.2.3
PDMS preparation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50
4.2.4 Uniaxial tensile testing . . . . . . . . . . . . . . . . . . . . . . . . . . . . 51

viii

.

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BIBLIOGRAPHY .

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4.4 Conclusions .

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CHAPTER 5 FUTURE .

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4.3 Results .

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 51
4.3.1
SiNC characterization and nanocomposite formation . . . . . . . . . . . . 51
4.3.2 Nanocomposite characterization . . . . . . . . . . . . . . . . . . . . . . . 53
Photoluminescence . . . . . . . . . . . . . . . . . . . . . . . . . 53
4.3.2.1
4.3.2.2 Microscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 53
4.3.2.3 Mechanical properties of the composite . . . . . . . . . . . . . . 57
4.3.2.4 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 58
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 61
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62
. . . . . . . . . . . . . 63
5.1 AuNP Application for Silicon NanoRods (SiNRs) Sythesis
5.2 Modulus of SiNC Layers on PDMS: Parametric Investigation . . . . . . . . . . . . 68
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 68
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 71
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75
.
77
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 77
SiNC deposition . . . . . . . . . . . . . . . . . . . . . . . . . . 77
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78
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. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 89

Introduction .
5.3.1.1
5.3.2 Results
.
5.3.3 Conclusion .
Future work .
5.3.4
.
.

5.2.1
.
5.2.2 Experimental work .
5.2.3 Results
.
5.2.4 Conclusions .

5.3 Unbending- Another Finite Bending Condition for Mechanical Property Estimation

Introduction .

5.3.1

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ix

LIST OF TABLES

Table 2.1: Reaction parameters for plasma-based AuNP synthesis . . . . . . . . . . . . . . 18

Table 2.2: Comparing plasma-based and hotwire methods . . . . . . . . . . . . . . . . . . 25

Table 3.1: Critical angles of bifurcation for di erent slenderness ratios, expressed in

degrees. Angles are reported for ⇤ = 2, 3 and 5. . . . . . . . . . . . . . . . . . . 42

Table 5.1: List of operation used for making surface instabilities in SiNC plasma reactor

.

. .

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at PNL .

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.

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 84

x

LIST OF FIGURES

Figure 1.1: Hot wire synthesis of AuNPs as developed by Boies et al.

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

7

Figure 2.1: An overview of gold nanoparticle syhthesis . . . . . . . . . . . . . . . . . . . . 15

Figure 2.2: Photograph and schematic of plasma-based AuNP synthesis.

Insets show

TEM images of AuNPs.

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

Figure 2.3: Schematic and photograph of hot-wire AuNP synthesis, including inset table

of reaction parameters.

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 19

Figure 2.4:

Imaging of AuNPs using transmission electron microscopy (TEM). (a) Bright
ﬁeld and (b) dark ﬁeld low-magniﬁcation images of AuNPs collected on
carbon ﬁber TEM grid; (c) higher-magniﬁcation TEM image of single Au NP
with [2 0 0] crystallographic plane identiﬁed as based on (d) FFT of the same
AuNP.

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 19

. .

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Figure 2.5:

(a)STEM imaging conﬁrms crystalline AuNPs, with inset showing the scan-
ning line for EDS analysis. (b) EDS conﬁrms the composition of the AuNP.

. . 20

Figure 2.6: AuNP average diameter and standard deviation of various samples from Ta-

ble.2.1, as measured from TEM images . . . . . . . . . . . . . . . . . . . . . . 21

Figure 2.7:

Figure 2.8:

(a) SEM image of the consumable Au wire surface before plasma exposure. (b) SEM
image of the Au wire surface after plasma exposure (conditions as for Sample 2 in
Table2.1), showing noticeable roughness. (c) Higher magniﬁcation of SEM image
of the Au wire after the plasma process.

. . . . . . . . . . . . . . . . . . . . . . . 22

(a,b) TEM images of AuNPs as synthesized using plasma-based method (Ar: 900
sccm, pressure: 180 Torr) and the hot-wire method (Ar: 900, DC power: 12 W)
respectively. (c) Average AuNP diameter and standard deviation for the plasma-based
method (circular symbols, Samples 1 and 2 from Fig.2.1) and the hot-wire method
(square symbols, Samples 3 and 4 from Fig.2.3) as measured using TEM image
analysis. AuNP, gold nanoparticle; DC, direct current; TEM, transmission electron
microscopy .

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24

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Figure 2.9: OES test for Plasma-Based AuNP synthesis, 340-1050 nm spectrometer with

Hydrogen(red) and without Hydrogen(blue)

. . . . . . . . . . . . . . . . . . . 28

xi

Figure 3.1: Steps from plasma deposition to ﬁnite bending. SiNCs are deposited through
a slit-shaped oriﬁce onto PDMS, which is then bent to observe formation
of wrinkles on the SiNC surface.
(b)
Schematic diagram of reactor and SiNC synthesis. (c) SiNC/PDMS sample
illuminated with UV light, demonstrating photoluminescence.
(d) Finite
bending applied with a lab-designed apparatus. . . . . . . . . . . . . . . . . . . 32

(a) Photograph of plasma reactor.

Figure 3.2: SEM and FIB cross-sectioning of the SiNC layer. The thickness is estimated
to be (⇠4.50 µm). (a) SEM image showing the surface of the SiNC layer on
PDMS, including the FIB-milled area (inset). (b) Higher magniﬁcation SEM
image of the FIB-milled area demonstrating the thickness of the SiNC layer.

. . 35

Figure 3.3:

(a) X-ray di raction conﬁrmed formation of SiNCs. XRD for SiNCs shows
crystal reﬂections at 2✓ values 28.3, 47.3 and 56.1 degrees consistent with
the (111), (220), and (311) planes of silicon.
(b) TEM image conﬁrming
the formation of SiNCs: black circles highlight individual SiNCs. The inset
shows FFT of the correspondig SiNCs of the TEM image.

. . . . . . . . . . . . 36

Figure 3.4: The stress machine loaded with a colored PDMS and dots depicted on the

PDMS, dislocation of dots will be recorded by a digital camera . . . . . . . . . 37

Figure 3.5:

(a and b)A schematic diagram of proposed aparatus for ﬁnite bending mea-
surement desgned by Solid Work software and (c) Alluminum cast aparatus
designed at PNL .

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 37

.

.

Figure 3.6: PDMS sample in the ﬁnite bending apparatus. (a) Side view, showing sam-
(b) Top view showing

ple (with handles) in the ﬁnite bending apparatus.
formation of bifurcations.

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38

Figure 3.7: Reference and deformed conﬁgurations for the PDMS-SiNC bilayer. The
bilayer deforms from the Cartesian coordinate system in the reference conﬁg-
uration (e1, e2, e3) to a cylindrical coordinate system in the deformed conﬁg-
uration (er, e✓, ez). Speciﬁcally, a plane at constant x0
1 transforms to a plane
at constant r (dashed line), and a plane at constant x0
2 transforms to a circular
arc at constant ✓ (dashed-dotted line). Since the out of plane deformation is
taken to be zero, we have x0

. . . . . . . . . . . . . . . . . . . . . . . . . 39

3 = z.

xii

Figure 3.8: Critical angles of bifurcation evaluated as a function of ratio of material
properties for a PDMS-SiNC bilayer sample. The ratio of the neo-Hookean
material parameters is deﬁned as ¯µ0 = µNC
0 . All curves have been
calculated for a thickness ratio H included in the range 0.0013–0.0018. Each
line represent a di erent value of slenderness ratio, ⇤ = 2, black solid line; ⇤
= 3, black dashed line; ⇤ = 4, dark gray solid line; ⇤ = 5, dark gray dashed
line; ⇤ = 6, light gray solid line. The insert represents a schematic of the
bilayer and deﬁnes the slenderness ratio.

0 /µP

. . . . . . . . . . . . . . . . . . . . . 42

Figure 3.9: Plotting points from the experiments on graph obtained by numerical cal-
culations. The ratio of the neo-Hookean material parameters is deﬁned as
¯µ0 = µNC
0 . Shown here are points obtained for ⇤ = 2, 3 and 5. Insert
shows an instability on the surface of a sample with ⇤ = 2.

0 /µP

. . . . . . . . . . . 43

Figure 4.1: SiNC Characterization: XRD (a) shows crystallographic reﬂections for nano-
sized silicon. Gray lines are the actual data and the black trace is the peak
ﬁt. TEM imaging (b) indicates spherical nanocrystals and lattice spacing
consistent with crystalline silicon. Inset shows FFT of image.

. . . . . . . . . . 52

Figure 4.2: SiNCs in solution, (a) photographed in room light and b) under UV excitation
Shown are hexane, NF, and F samples from left to right. (c) PL spectroscopy
shows much brighter luminescence for surface-functionalized (F) SiNCs as
compared to nonfunctionalized (NF) SiNCs in solution.

. . . . . . . . . . . . . 53

Figure 4.3:

(a) Photographs of SiNC/PDMS nanocomposites in room light and (b) under
UV excitation (c) shows the average PL QY of SiNCs in hexane and in
the nanocomposites, with error bars reﬂecting the standard deviation of the
measurements.

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 54

.

.

.

Figure 4.4:

(a) For reference during microscopy, we assigned each nanocomposite the
same coordinate system, (b) A single CLSM image of a slice is shown and
(c) the depth-resolved composite image (scale inset), demonstrates the distri-
bution of small clusters within the sample slab.

. . . . . . . . . . . . . . . . . 55

Figure 4.5: Representative CLSM images of the SiNC/PDMS nanocomposites

. . . . . . . 56

Figure 4.6: SiNC cluster size statistics based on image processing from the CLSM images.
3 separate images were analyzed for each sample type. A key to the data
representative is shown as inset.

. . . . . . . . . . . . . . . . . . . . . . . . . 56

Figure 4.7: Mechanical properties of the nanocomposites. (a) shows a representative plot
of Cauchy stress vs. stretch data for each sample type. From this data, (b)
we extracted the neo-Hookean coe cient for each sample, showing a distinct
reduction for the higher-concentration sample.
(c) the rupture stretch was
una ected by the SiNC inclusions.

. . . . . . . . . . . . . . . . . . . . . . . . 58

xiii

Figure 5.1:

(a) Schematic Diagram for SiNRs synthesis (b) early stage experimental set
up with hot wire method (Courtesy R.J.Anthony) . . . . . . . . . . . . . . . . . 64

Figure 5.2: High ﬂowrate of AuNP-Ar from the ﬁrst step(600-1200sccm) and low ﬂowrate

of SiH4 from the side arm did not mix properly in the second step of the reactor

65

Figure 5.3: Modiﬁed hotwire AuNP reactor for sythesis SiNRs . . . . . . . . . . . . . . . . 65

Figure 5.4:

(a) Filament formation associated with high plasma density (b) High plasma
density resulted in wide size range of Silicon nanoparticles with both crys-
talline and amorphous structures. (c) Controlling the SiH4 ﬂowrate and RF
power stoped ﬁllament and resulted in narrow size distribution for Silicon
nanoparticles .

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 67

.

.

.

.

Figure 5.5: Schematic of dual-plasma reactor for AuNPs-SINRs simultaneous synthesis

. . 68

Figure 5.6: Schematic diagram for investigating the impact of ﬁlm thickness, surface
functionality of SiNCs via gas phase plasma and shear modulus of the PDMS
substrate on mechanical properties of SiNC thin ﬁlm . . . . . . . . . . . . . . . 71

Figure 5.7: SEM images for PDMS with 3:1 slenderness ratio as deﬁned in chapter 3
with rastering deposition time of 5 mins (Sample 1), 10 mins (Sample 2)
and 30 mins (Sample 3) and averaged measured thickness of the ﬁlm with
corresponding standard deviation in comparison with the sample discussed
with same slenderness ratio and 20 mins (in red)of rastering deposition time
from previous chapter .

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 72

Figure 5.8: The methodology introduced in Chapter 3, employed to estimate neo-Hookean
coe cient for SiNC layars corresponds to parametric variation. The asterisks
indicate samples with signiﬁcant deviation from value we abrtained in Chapter 3 73

Figure 5.9: Neo-Hookean coe cient of all sample introduced in this section.The sam-
ples with asterisks, present the data with siginﬁcant deviation from samples
reported in Chapter 3 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 74

Figure 5.10: The surface functionalization applied in gas phase during plasma processing
and sythesis of SiNCs. The PL and QY measured. This results coupled with
mechanical property estimation of corresponding samples presented in this
Figure. The results here, conﬁrm that improving the oprical properties and
QY do not interfere with mechanical property of the SiNC thin ﬁlm.

. . . . . . 75

Figure 5.11: Sample holder for various stretching ratios (left), sample holder for ﬁnite

bending around pushing rod (right) . . . . . . . . . . . . . . . . . . . . . . . . 78

Figure 5.12: Above view of sample, relaxed after deposition

. . . . . . . . . . . . . . . . . 78

xiv

Figure 5.13: SEM image of deposited area with 20 percent pre-stretching condition, Col-
umn A3-B3-C3 (left), Columns A2-B2-C2 and A4-B4-C4 (middle), Columns
A1-B1-C1 and A5-B5-C5 (right)

. . . . . . . . . . . . . . . . . . . . . . . . . 79

Figure 5.14: SEM image of deposited area with 20 percent pre-stretching condition, Col-
umn A3-B3-C3 (left), Columns A2-B2-C2 and A4-B4-C4 (middle), Columns
A1-B1-C1 and A5-B5-C5 (right)

. . . . . . . . . . . . . . . . . . . . . . . . . 79

Figure 5.15: pre-stretched before deposition or during deposition . . . . . . . . . . . . . . . 80

Figure 5.16: Early design for ﬁnite bending surface instabilities by wrapping PDMS around
the rod, the picture on right is the microscope image of the sample after
unwrapped with the scale bar of 10 micrometer . . . . . . . . . . . . . . . . . . 82

Figure 5.17: A new method for pure bending surface instabilities, the mathematical cor-
relation and assumptions for incompressible, Neo-Hookean material (left)
.

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 83

.

.

.

.

.

.

.

.

.

.

.

Figure 5.18: A new method for pure bending surface instabilities, the mathematical cor-
relation and assumptions for incompressible, Neo-Hookean material (left)
.

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 83

.

.

.

.

.

.

.

.

.

.

.

Figure 5.19: XRD for SiNCs deposited on PDMS, 30 Watt(Left), 25 Watt (right) . . . . . . . 85

Figure 5.20: Lambda ( ) measurement for surface instabilities, 10 round rotations in the
reactor(up-left), 20 round rotations(up-right) and 30 round rotations (bottom),
numbers are corresponding to the cells that measurement done with speciﬁed
standard deviation for each averaged size letter corresponding to the cell
which wrinkle instabilities observed . . . . . . . . . . . . . . . . . . . . . . . 86

Figure 5.21: A sample of PDMS deposited with SiNCs under pre-stretched condition for
10 round rotation(left), deﬁnition of Lambda ( ) and thickness of the SiNCs
ﬁlm(right)

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 86

. .

.

.

.

.

Figure 5.22: Atomic Forced Microscopy (AFM), AFM scanned a line across a 20 by 20
micrometer area and provided a 3D model and a cross-sectional view for
investigating the amplitude of these surface instabilities . . . . . . . . . . . . . 87

xv

CHAPTER 1

INTRODUCTION

The focus of this dissertation is synthesis and use of nontoxic nanomaterials for a variety of
applications. Nanomaterials such as nanocrystals are critical for a range of novel applications
ranging from biological treatments and diagnostic procedures to devices such as light-emitting
devices (LEDs), solar photovoltaics, and wearable sensors. There are two main thrusts to the
research: one focuses on new gas-phase routes to freestanding, surfactant-free gold nanoparticles,
including a novel plasma reactor. The other explores the optical and mechanical behavior of
plasma-synthesized silicon nanocrystals. This introductory chapter sets the stage for discussing
the subsequent research by providing a background on plasmas, the history and use of gold
nanoparticles, and the prior work on studying the mechanical behavior of thin ﬁlms of silicon
nanocrystals.

1.1

Introduction to Non-thermal Plasma

Today, many research studies have been dedicated to investigating the synthesis and use
of nanocrystal-based materials, including those processed in the liquid phase [83, 122, 164].
While liquid-phase synthesis and/or surface modiﬁcation has yielded exceptional results for some
nanocrystalline materials, there are limitations imposed by solution processing. For example,
the temperature of liquid phase synthesis methods are limited by the boiling points of solvents
[83]. This can hamper the formation of some nanocrystals due to the required high crystalliza-
tion temperatures, though some approaches have been successful for materials such as silicon
nanocrystals [36, 63, 64, 189]. Additionally, liquid-phase synthesis and/or surface modiﬁcation
processes necessitate disposal procedures for the solution wastes and by-products which are time-
consuming and are not environmentally sustainable.
In contrast, gas phase approaches do not
require organic solvents during the synthesis process and are capable of achieving nanocrystals
from high-temperature-crystallization materials. There are two classes of gas phase synthesis

1

methods, di erentiated according to the thermal equilibrium during the process. Thermal equi-
librium syntheses include ﬂame synthesis [83, 158], thermal pyrolysis in furnaces [129], laser
pyrolysis [83, 91] and thermal plasmas [142]. In this class of gas phase synthesis, the entire system
achieves a relatively high temperature (e.g. gas species, excited species, and synthesized nano-
materials). During many thermal-equilibrium processes, the produced nanoparticles have reduced
charging and a large degree of agglomeration. Some liquid phase post-processing steps are often be
required in this regard. Also, higher temperature thermal plasma increases the chance of broad size
distribution [83]. Gas-phase processes that are not in thermal equilibrium (such as non-thermal
plasmas) are the alternate class of method for nanoparticle synthesis. In non-thermal equilibrium
we have heavy gas species and ions that are close to room temperature, while free plasma electrons
(which have negligible mass in comparison) can achieve temperature between 1000 to 50000 K
[83, 107]. Collisions between plasma electrons and molecules in gas ﬂow dissociate and ionize
gaseous precursors, producing highly reactive components [83]. Non-thermal plasmas have several
of advantages including ligand-free, high-purity ﬁnal products. High melting point materials can
be synthesized due to highly exothermic reactions just at the surfaces of the forming nanoparticles.
In non-thermal plasmas, synthesis of very small nanoparticles (with sizes below 10 nm and size
distributions around 10-20% std. dev.) are possible. Di usion loss to the walls, together with
nanoparticle agglomeration, are negligible due to negative charges of surface of the reactor walls
as well as at the surfaces of nanoparticles - both of these are the result of the high kinetic energy
and collision frequency of electrons with surfaces in the plasma. Additionally, nanoparticle growth
occurs over a particle residence time on the ms time scale [3, 83, 107]. These attributes make
nonthermal plasmas an interesting substitute for liquid phase synthesis.

In general, nonthermal plasmas are classiﬁed by low ionization fraction.

In radiofrequency
(RF) plasmas, such as the ones used in this dissertation, ions are generated by the applied rf ﬁeld.
This ﬁeld causes the free electrons to accelerate to high speed. The number of free electrons
exponentially increases in the plasma and few of these electrons are able to exchange energy with
heavier gas species, such as precursor molecules [83]. Most of the energy provided to the plasma

2

also absorbed with electrons since they have higher mobility compared to ions. The densities of
electrons and ions remain approximately equal due to the principle of quasineutrality in the plasma,
except near the walls (this small region of non-neutrality is called the plasma sheath) [18, 92]. The
di erence between speed and temperature of electrons and ions in the plasma is what establishes
the condition of non-thermal equilibrium, and is also the cause of the negative charging of surfaces
(i.e. walls and nanoparticles) in the plasma [25, 83].

In the non-thermal plasma, nanoparticles will be charged to negative potentials in the order of
kBTe/2e. By considering the typical temperature of electrons, individual nanoparticles will have
a charge of up to few volts. As an example, a typical particle ﬂowing in the plasma with diameter
of around 1 nm carrying a single electron will have a negative potential of -1.4 V [83, 136]. In
order to validate a charge balance between electron and ion currents reaching the surface of the
nanoparticles, it is essential to note that nanoparticles of di erent sizes will charge to approximately
the same steady state potential.

Bouchoule et al. [17, 83] proposed a three-steps mechanism for nanoparticle growth in non-
thermal plasmas. Although the repulsion between negative charges prevents nanoparticles from
agglomeration, the smaller particles at earlier stage of growth (nanoparticles with size less than 1 nm)
usually have no charge and their aggolomeration is one of the main mechanisms for nanoparticles
growth. First, initial clusters of precursor fragments generate at nucleation stage, forming proto-
nanoparticles. Second, since the nanoparticles emerging from the ﬁrst stage are mostly less than
1nm in size and has no charge, larger particles grow quickly through the coagulation of initial
neutral small clusters [83]. This step continues up to point that all of nanoparticles have single
charge. Finally, the growth of nanoparticles continues by reaction of precursor on their surface.

Non-thermal plasma reactors can operate under reduced pressure or atmospheric pressure.
These reactors can include various components and characterization tools including the reaction
chamber, vessels for precursor injection [83], gas processing pipelines, analysis ports, and the
nanomaterial collection system (often by impaction or di usion). The reaction chamber can be
modiﬁed over a vast range for ﬂow-through and batch processing requirements.

3

This dissertation research employs plasma reactors for nanomaterials synthesis in two separate
projects. A background on alternate methods provides useful context. In the next section we brieﬂy
review prior methods for fabrication of noble metal nanoparticles, pertinent to the chapter on
gold nanoparticle synthesis using a novel non-thermal plasma reactor. Then we brieﬂy talk about
silicon based- nanomaterials for electronic, sensor, and optoelectronic applications, especially
nanoparticles produced by non-thermal plasmas. In addition, we introduce poly(dimethylsiloxane)
(PDMS) as a stretchable substrate with some unique physical properties that make it suitable for
probing the mechanical behavior of silicon nanoparticle layers. Finally, we review some existing
mechanical models for estimating the mechanical properties of a thin ﬁlm of semiconductor layers
on a ﬂexible substrate.

1.2 Gold Nanoparticles

Gold nanoparticles (AuNPs) are one of the oldest forms of nanostructures, with medical appli-
cations dating back to 1618 [15]. Today scientists are especially interested in Localized Surface
Plasmon Resonance (LSPR), a phenomenon of unique optical shape-dependent property with quan-
tum features, which results from the small size of these structures in less than 10nm. LSPR is
an optical property which occurs when a noble metal nanoparticle interacts with a photon[43].
Depending on the metal, particle size, particle shape, and photon energy, the free electrons in the
nanoparticle can undergo resonant oscillations and amplify or otherwise modify the electromag-
netic ﬁeld nearby. LSPR has two main consequences. First, the electric ﬁeld near the particle’s
surface is greatly enhanced; second, the nanoparticle’s optical absorption has a maximum at the
plasmon resonant frequency (which is dependent on metal, size, and shape). The enhancement
decreases quickly with increasing distance from surface of the particle. LSPR occurs at visible
wavelengths for the noble metal nanoparticles including AuNPs [21, 71, 90, 98, 101, 147, 198].
The LSPR e ects (both for emission and absorption) from AuNPs can be used in many applications
including as biological labels, therapeutic cancer treatments, solar photovoltaics, near-ﬁeld and
electron microscopy, and more [47, 70].

4

In addition to uses centered around their LSPR, metal nanoparticles can be used to catalyze the
growth of anisotropic semiconductor nanostructures such as silicon nanorods (SiNR) [63]. These
anisotropic structures o er novel and shape/size-dependent optoelectronic properties, and so the
catalysis o ered by AuNPs and others adds to their attractive properties.

1.2.1 History of AuNP synthesis

Among the synthesis methods for AuNPs production, many involve wet chemistry-based ap-
proaches, which require ligands and other solution-phase modiﬁers to the naoparticle surfaces
[101, 121]. Solution based syntheses are among the earliest methods for formation of gold nanos-
tructures and are still used today[20, 35, 39, 49, 54, 101, 171, 197]. In general, AuNPs in solution
phase produced by reduction of chloroauric acid. In these methods the naoparticle growth happens
step by step from molecular precursors (also called bottom-up), leading to better control on the
shape and size distribution of AuNPs in the ﬁnal product. In the solution-based routes, the AuNPs
end up solubilized in water or other solvents. We endeavored to eliminate the possibility of con-
tamination from these liquids and extend the uses of AuNPs by synthesizing the AuNPs in the gas
phase. This will open up a new approach for AuNPs synthesis and reduce both product contami-
nation as well as reduce the need for solution disposal after-process. Our research builds on prior
work to synthesize AuNPs in the gas phase, particularly aiming towards a synthesis method that
permits easy integration of the AuNP product into subsequent reaction steps without interruption
of process parameters. We review these pre-existing methods below, and in Chapter 2 we introduce
our new plasma-based method.

1.2.2 AuNP formation in the gas phase

Other work has led to gas-phase, solvent-free synthesis routes to AuNPs. There are several
options for synthesis in the gas-phase, including inert gas condensation of Au vapor in a chamber,
spark-ignition AuNP generation, sputtering techniques, and hot-wire synthesis. Here we detail
the hot-wire synthesis method, which has demonstrated promise in multi-step reaction procedures.

5

[15]. This approach is similar to other gas-phase methods in that it requires a gas ﬂow and a
heat source for generating the AuNPs. Since the metal wire is heated via electrical resistance,
the average background gas temperature will not increase except in the area locally close to the
resistive element. Under the right conditions of gas ﬂow, pressure, and suplied electrical current,
the evaporated metal atoms reach supersaturation upon leaving the metal wire surface, and rapidly
nucleation of nanoparticles begins. Gas drag then carries the AuNPs downstream for collection
or further use. The big advantage of the hot wire method is that according to previous studies,
background gas has small e ect on property of metal nanoparticles. Peineke et al. described that
the hot wire method is suitable for all metals which have vapor pressure of higher than (3.75*10-5
Torr) at the melting point. According to this guideline, gold (with a vapor pressure of 1.5*10-5
Torr) does not meet this requirement. However, as Boies et al. reported [15], gold can be used
if the source is an electroplated wire structure. For example, they had success using a Platinum
(Pt) (99.5%, trace metal, Aldrich) wire with 0.5 mm in diameter as the core. The gold as the outer
shell was deposited via electroplating by the ProPlate Company in Minnesota. Studies have shown
that hot wire method with a coaxial gold source can produce AuNPs with less than 10 nm average
diameter [15, 16]. Figure1.1 depicts the schematic diagram of hot wire gold nanoparticle formation
method as developed by Boies et al. This method is promising as it results in AuNPs entrained into
the gas ﬂow – but had only been reported at atmospheric pressure prior to our work. For coupling
into a low-pressure reactor for possible multi-step reactions including non-thermal plasmas, we
were interested to explore this method as an AuNP catalyst source at lower pressure. We were also
interested in other methods for AuNP synthesis at lower pressure.

Our work on a reduced-pressure approach to the hot wire synthesis of AuNPs is detailed in
Chapter 2, alongside our introduction of a plasma-based synthesis of AuNPs, using a consumable
gold wire and a radiofrequency power source. In our plasma method, the AuNPs are mono-disperse,
with an average diameter of 4 nm, o ering some improved properties as compared to our attempt
at a low-pressure hot wire synthesis route. Notabely, this size range is one of the most interesting
for optoelectronic and LSPR functionality for AuNPs. As shown in that chatper, although the

6

Figure 1.1: Hot wire synthesis of AuNPs as developed by Boies et al.

production yield is low, the narrow size distribution of the AuNPs and the avoidance of solution
processing in this method are promising for future syntheses of metal NPs based on plasmas.

7

1.3 Silicon Nanoparticles

The other focus of this dissertation is on deployment of silicon nanoparticles onto and within
polymer substrates for evaluation of optical and mechanical properties of these bilayered and hybrid
macroscale structures. Silicon has been a dominant material in the ﬁeld of electronics over the
last half century or more, due to its semiconducting electronic behavior, optical band gap at 1.1
eV, its natural abundance, its low cost, and its relative nontoxicity. Despite the long history of
use in transistors, traditional solar photovoltaics, bulk silicon has poor emission behavior. Due
to its indirect band gap, silicon requires that photon emission and absorption processes involve a
momentum-balancing phonon [107]. This disadvantage limits the applications of bulk silicon as
a potential material for host of optoelectronic devices including emissive applications, although
doping, alloying, and combination of bulk silicon into hybrid structures with other direct band gap
semiconductors has allowed some device development in the area of optoelectronics. [75].

However, since the late 1980s, new discoveries on the behavior of nanoscale silicon have
revealed that optical processes become much more e cient when the silicon becomes quantum-
conﬁned, meaning its size becomes small enough that electron-hole combination processes are
a ected.
In particular, when the silicon nanostructure size shrinks to approximately the Bohr
exciton radius (around 5 nm), the electron and hole wavefunctions can exceed the nanostructure’s
physical size. This widens the band gap, relieves some of the quantum selection rules governing
the recombination processes, and improves the e ciency of photon processes for the relaxation of
the exciton. This is manifested in size-tunable photon energy and increased photolumienscence
(PL) e ciency. Many studies have been probed the PL and other optical properties of silicon
nanocrystals (SiNCs), leading to their application in optoelectronic devices, microelectronics, and
solar photovoltaics [5, 26, 32–34, 39, 107, 139, 167].

In the last few decades many advancements on inﬂuence of surface conditions on the lumi-
In particular,
nescence properties of quantum conﬁned SiNCs have emerged [107, 176, 199].
passivating the SiNC surfaces with organic ligands has been demonstrated to improve their PLQY.
There are various approaches to this, including solution-phase thermal hydrosilylation, UV- and

8

white-light-initiated hydrosilylation, and gas-phase methods including using furnaces and low-
energy plasmas [5, 37, 76, 105, 106, 150, 189].

The wide range of applications for luminescent SiNCs had been limited by the lack of a simple,
fast, and inexpensive synthesis approach [107]. Liquid-phase approaches, including synthesis in
inverse micelles [187] synthesis in high temperature supercritical solutions,[42] the oxidation of
metal silicide,[135] and the reduction of Silicon tetrahalides and other alkylsilicon halides,[8] are
mostley time consuming processes. These methods usually have low production rates. Gas-phase
methods including the pyrolysis of silane in furnace ﬂow reactors, [94, 107, 127] the decom-
position of silane through laser light irradiation (photolysis),[12] and laser pyrolysis using high
power infrared lasers [107], all may yield faster processing times. Although these methods have
faster production time, are prone to particle agglomeration. In order to deagglomerate the ﬁnal
clusters,further post processing is necessary. These post processing steps might diminish quan-
tum conﬁnement properties [107]. However, non-thermal plasma processes as among the most
promising, due to the rapid, facile synthesis and tight control over SiNC properties correspond to
the easily adjustable reactor parameters such as pressure, reactor dimensions, supplied power, and
gas species.

SiNCs experience an interesting journey during synthesis in a non-thermal plasma reactor. One
of the earliest reports regarding the particle heating in plasmas dealt with micron-sized particles
that were injected into argon plasmas [38]. For 0.5-3 µm, they measured particle temperatures that
were 75 k above the gas temperature [38]. The propsed model is that the heating of the particles
are due to electron-ion recombination and the main cooling event is due to radiation and coduction
at the surface of the particles.

Despite a low overall gas temperature, SiNCs can reach elevated temperatures,as pointed out
by Daugherty et al. and Mangolini et al. [38, 107]. These investigations considered the e ects
of electron-ion recombination at SiNC surfaces, other reactions between neutral gas species with
SiNC surfaces, radiative cooling processes, and gas conduction cooling processes. Results indicate
that the SiNC temperatures can spike to several hundred Kelvin higher than the gas temperature.

9

These spikes are responsible for the crystallization of the SiNCs. These works also revealed that
bigger nanoparticles require higher temperatures for their crystallization. Signiﬁcantly, smaller
particles will gain higher energy (corresponding to higher spike temperature) due to exothermic
recombination events on their surfaces[107]. Kramer et al. [84, 85] studied the di erences between
nanoparticle heating and cooling events at atmospheric pressure and low pressure respectively.
Since gas conduction is faster in atmospheric pressure, SiNCs produced at atmospheric pressure
experienced higher spikes of temperature and fast cooling successively. However, in low pressure
environment, cooling is much slower and SiNCs experience higher average temperature after
successive heating and cooling events.

One of the major attractions of nonthermal plasmas compared to other gas-phase methods
is their ability to produce quantum-conﬁned, luminescent SiNCs that are smaller than 5 nm [83].
Studies have shown that decreasing residence time in the reactor produces smaller particles [55, 83].
Synthesis of SiNCs with tunable sizes between 3 nm and 6 nm and residence time of 2 ms and 7 ms,
respectively, in a capacitively coupled plasma (CCP) has already been reported more than a decade
ago [107]. One important consideration is that the large surface-to-volume ratio for nanoparticles of
all types means that e cient luminescence processes require a surface with few defects. For many
materials this is accommodated by ligand attachment at the surfaces, a process which is intrinsic to
their synthesis in the solution phase. With plasma-produced SiNCs, as-produced hydrogen-coated
nanoparticles have a high density of defects and thus a low PL yield. With air exposure, an oxide
layer at the shell of these plasma-produced SiNCs enhances the luminescence e ciencies. Even
more improved luminescence for plasma-produced SiNCs can be achieved via passivation of the
SiNCs by surface-functionalization with organic compounds, such as alkene organic ligands. This
covalent bonding at the surfaces of the SiNCs helps to achieve solution PL quantum yields (QYs) as
high as 60% [79, 105]. Our work on using plasma-produced SiNCs in bilayer and hybrid polymer
structures (3) includes surface functionalization using a corresponding procedure, and application
of these surface-modiﬁed SiNCs.

Andthony et al. [7] pointed out that the injection of hydrogen H2 downstream of the synthesis

10

plasma also improves the quantum yield of the resulting SiNCs [83]. The improvement has been
conﬁrmed by electron paramagnetic resonance (EPR), with lower surface defects in SiNCs produced
using H2 injection. [133]. These prior discoveries built the foundation of our approach to synthesis
and surface-modiﬁcation of SiNCs, as will be discussed in subsequent chapters.

1.3.1 SiNCs on Flexible Substrates: Polydimethylsiloxane (PDMS)

In Chapter 3 and Chapter 4, we introduce methods for investigating the SiNCs both on top of and
embedded within polydimethysiloxane (PDMS), a commonly-used ﬂexible/stretchable elastomer.
We anticipate these structures to ﬁnd applications in the next generation of gas sensors, solar panel
devices beyond traditional architectures, LEDs, and other ﬂexible microelectronics[26]. Previously,
luminescent SiNCs have been used exclusively in traditional rigid devices - but ﬂexible, stretchable
devices o er exciting opportunities. Recently, the Plasmas and Nanomaterials Laboratory (PNL)
performed a study on the aging properties of SiNCs on PDMS [104]. Keeping the synthesis
parameters constant, SiNCs were deposited directly out of the plasma reactor into thin-ﬁlm layers
on the PDMS and onto bulk Si wafer substrates and/or glass using inertial impaction through a
slit-shaped oriﬁce [3, 5, 68]. This impaction results in a di erent thin ﬁlm morphology on the rigid
substrate as compared to the PDMS. Speciﬁcally, we observed the SiNC ﬁlm to be less dense on the
rigid substrate, as observed using electron microscopy [104]. The reason behind this morphological
di erence is under investigation, however, is likely attributed to the di erence in elastic properties
of the PDMS as compared to the rigid substrate. In chapter 4 we introduce a in situ method for
measuring mechanical properties of SiNCs and PDMS as a bilayer system.

PDMS belongs to a class of polymeric compounds that are commonly referred to as silicones.
PDMS is by far the most widely used organosilicon compound polymer [1]. PDMS also has
distinctive ﬂow properties. PDMS is inert, non-toxic and non-ﬂammable. Its applications span a
wide range, including in contact lenses and medical devices, additives in beauty products, food-
related industries as a foam prevention agent, adhesives, replica molding, micro-nanoﬂuidic and
membranes, and more [24]. Moreover, PDMS is a useful functional polymer material in academic

11

experimental research and industrial applications due to its unique properties including reasonable
thermal stability, high transparency, low surface free energy and excellent biocompatibility which
provides a wide range of other applications for medical device development and health monitoring
sensors [24].

PDMS is a viscoelastic polymer. Its viscosity is a function of temperature. During long ﬂow
times and high temperatures, it acts like liquids with high viscosity, however, in low temperature
it acts like a rubber-like material. Viscoelasticity is a common form of nonlinear elasticity for a
wide range of polymers [46]. The elastic behavior of PDMS mostly depends on the concentration
of cross-linking agent used during the curing process. PDMS can conform to various types of
surfaces during the curing process. Obviously this feature makes PDMS a tunable material for
a wide range of experimental applications including, microelectromechanical systems [111, 145].
Mostly, the cross-linked cured version of PDMS as a rubber-like material, has a solidiﬁed form
and is extensively used for stretched, bent, compressed and other type of stress-based deformation
experiments [185]. Generally, PDMS has a low elastic modulus which enables it to be easily
deformed and results in the a rubber-like behavior [78, 96]. The shear modulus of PDMS changes
with preparation conditions, and consequently varies between the range of 100 kPa to 3 MPa. This
topic will be discussed further in chapter 3 and chapter 4.

1.3.2 Mechanical properties of SiNCs, deposited as thin ﬁlms

To date, nanoparticles have been used for variety of applications. However, there is a noticeable
lack of knowledge regarding the mechanical behavior of layers of nanoparticles, especially the
layer’s ability to bend, stretch and deform when deposited on a ﬂexible substrate. The mechanical
behavior upon deformation of a substrate will cause instabilities in deposited thin ﬁlm structures,
including in layers of nanoscale particles. Additionally, when nanoparticles are mixed within
polymers, they can a ect the nanocomposite’s mechanical behavior. Combining the mechanical
properties of a polymer-based materials such as PDMS, with the tunable optical, electronic, and
estimated mechanical properties of SiNCs could lead to the introduction of a new class of highly

12

ﬂexible technologies utilizing multi-layer and composite architectures including wearable sensors
and foldable LEDs, among others. [10, 11, 156]. In order to be able to practically design devices
with such a new level of ﬂexibility, the mechanical properties of the substrate and thin ﬁlm layers
should be investigated individually and together as a system.

Large deformations such as uniaxial compression, equibiaxial compression, pure shear, ﬁnite
bending, and unbending are attracting interest as methods to estimate the mechanical behavior of
materials [14, 40, 143, 144, 153]. In addition, uniaxial tensile testing can reveal information about
the mechanical behavior of nanocomposites. Chapter 3, Chapter 4, and Chapter 5 discuss our
e orts to apply uniaxial tensile testing, bifurcation analysis, and theoretical modeling frameworks
to estimate the mechanical properties of nanocomposites based on SiNCs and of SiNC thin ﬁlms
deposited on PDMS directly from the gas phase.

In chapter 3 we discuss a novel technique for measuring the mechanical properties of thin ﬁlm
of nanocrystals, especially SiNCs, due to their vast ranges of application and growing market for
future device development, deposited on PDMS as a ﬂexible substrate. This method further can be
applied for other nanostructure thin ﬁlms and ﬂexible substrates. In Chapter 4 we investigate the
mechanical properties of a composite structure, SiNC clusters embedded in PDMS for luminescence
solar concentrator and other ﬂexible optical devices. In Chapter 5, we introduce a parametric study
to investigate the dependence of various parameters, including, thickness of the ﬁlm, modulus of
the substrate and surface functionality of SiNCs on the mechanical property of the thin ﬁlm in a
bilayer system. Later in Chapter 5 we also report our study of another class of surface instability,
which arises as a result of unbending of SiNC layers as deposited on an intially pre-stretched PDMS
system. In summary, these results shed light on the mechancial behavior of SiNCs, dependent on
their surface functionality, both in nanocomposite and in bilayer structures. This work contributes
to the future applications of SiNCs in ﬂexible, bendable, and stretchable manifestations.

13

CHAPTER 2

AUNP SYNTHESIS: A PLASMA-BASED APPROACH

Portions of this work were published in the journal "Plasma Processes and Polymers" in July 2019
and reproduced here. This article was also featured on the front cover of the July 2019 issue of the
journal.

14

Figure 2.1: An overview of gold nanoparticle syhthesis

2.1

Introduction

Metallic nanostructures have proven to be among the most ﬂexible nanomaterials owing to
the synthetic control of their size, shape, composition, as well as the resulting tunable optical
properties[72]. Gold nanoparticles (AuNPs) are among the oldest forms of engineered nanostruc-

15

tures, with medical applications dating back to the early 17th century. AuNPs are broadly attractive
across disciplines, and are used for applications such as medical theranostics, Surface Enhanced
Raman Spectroscopy (SERS), optoelectronic device enhancements, and as seeding sites for other
materials such as silicon nanowires[43, 61, 72, 97, 159, 170, 190].

Most synthesis methods of AuNP production involve wet chemistry-based approaches[20, 35,
49, 54, 73, 101, 121, 171, 197]. In solution-based routes, the AuNPs typically end up solubilized
in water or other solvents using surfactants, which is useful for some applications – but this can
also increase contamination for many post-processing purposes. To eliminate the possibility of
contamination in post-synthesis reactions and applications, there is interest in routes for synthesizing
the AuNPs in the gas phase. Pre-existing methods in gas phase include inert gas condensation,
spark ignition, plasma sputtering, and hot wire synthesis[15, 81, 117, 132, 137, 161]. Recently,
several groups have turned to ﬂow-through plasma synthesis of metallic NPs and nanostructured
ﬁlms[28, 29, 86, 188, 197]. The advantages to these plasma-based methods are a relatively narrow
size distribution and continuous synthesis. Here we introduce a new method for making AuNPs
based on a radiofrequency (rf) plasma together with a consumable Au electrode for synthesis of
monodisperse AuNPs with an average diamter of ⇠4nm. We compare this method with a low
pressure hot-wire technique for AuNP synthesis, ﬁnding that while the hot-wire method has a
higher yield of AuNPs, the plasma-based route produces NPs with a smaller size dispersity. This
approach adds to the versatility of plasma reactors for nanomaterials synthesis.

2.2 Experimental Details

We used a non-thermal radiofrequency (RF) plasma reactor for AuNP synthesis. Figure2.2
shows a schematic view of this reactor. The plasma is initiated via RF power (13.56 MHz) supplied
via a matching network to a copper coil around a quartz tube (2.54 cm O.D. and 20 cm length). The
grounded electrode is a pure Au (99.99) wire threaded coaxially within the tube, which is also the
source of Au for the AuNPs. We ﬂowed Argon (Ar) gas through the reactor at ﬂowrates between
600-1200 standard cubic centimeters per minute (sccm) and used oriﬁce plates downstream of the

16

Figure 2.2: Photograph and schematic of plasma-based AuNP synthesis. Insets show TEM images
of AuNPs.

reactor to modulate the pressure between 60-220 Torr. The pressure beneath the oriﬁce was <10
Torr.

We varied the power to the external coil between 5-15W (nominal), upon which the gas
surrounding the Au wire began to glow. To conﬁrm that the Au wire acts as the ground, we
attempted to operate the reactor without the coaxial wire – with no success at initiating a plasma.
We placed a TEM grid (Lacey Carbon Film Cu 300 mesh from TedPella®) at the exit of oriﬁce
for collecting and analyzing the AuNPs. Figure 1 shows a schematic diagram of our process, and
Fig2.1 shows a list of samples and corresponding parameters we explored for the plasma method,
including pressure, power, and discharge length.

We sought to compare our plasma method with an alternate technique, similar to the previously-
reported hot-wire AuNP synthesis technique at atmospheric pressure reported by Boies et al. [15],
but operated under reduced pressure of 10 Torr. Our hot-wire setup consisted of a glass tube with a
Platinum (Pt) wire electroplated with Au (0.5 mm thickness) and connected to a DC power supply.
The electroplated Pt-Au wire was threaded via an electrical vacuum coupler feedthrough into a

17

Table 2.1: Reaction parameters for plasma-based AuNP synthesis

sealed glass vacuum reactor tube and connected externally to a DC power supply. The glass reactor
tube was a three-way heavy-duty quartz glass (2.54 cm O.D. and 12.7 cm length) which was fused
to a ﬂexible stainless steel ﬂange to prevent it from breaking during experiments. Ar was used as
a background gas for carrying AuNPs towards the reactor exit, with ﬂowrates between 300-1200
sccm and reactor pressure from 2-30 Torr, depending on gas ﬂowrate and oriﬁce size. DC power
was supplied to the Pt-Au wire, resistively heating the gold at the surface and causing evaporation.
The gold atoms entered the gas phase and condensed to form AuNPs. Figure 2 shows a schematic
view of this reactor.

2.3 Results and Discussion

We measured the diameter of the plasma-produced AuNPs with Transmission Electron Mi-
croscopy (TEM) and Scanning TEM (STEM). TEM images showed individual AuNPs deposited
onto the grid at the exit of the reactor (Fig.2.4a-d). Both bright ﬁeld and dark ﬁeld images (Fig2.4)
show that the collection grid is covered with individual AuNPs. A higher-magniﬁcation image
shows lattice spacing, with the [2 0 0] plane identiﬁed (Fig.2.4c). We performed Fast Fourier
Transform (FFT) analysis of an individual AuNP to measure the crystal lattice spacing (Fig.2.4d).

18

Figure 2.3: Schematic and photograph of hot-wire AuNP synthesis, including inset table of reaction
parameters.

Figure 2.4: Imaging of AuNPs using transmission electron microscopy (TEM). (a) Bright ﬁeld
and (b) dark ﬁeld low-magniﬁcation images of AuNPs collected on carbon ﬁber TEM grid; (c)
higher-magniﬁcation TEM image of single Au NP with [2 0 0] crystallographic plane identiﬁed as
based on (d) FFT of the same AuNP.

We performed STEM combined with Energy Dispersive X-ray Spectroscopy (EDS) to further
conﬁrm the presence of Au and to measure possible contamination – which we did not observe

19

Figure 2.5: (a)STEM imaging conﬁrms crystalline AuNPs, with inset showing the scanning line
for EDS analysis. (b) EDS conﬁrms the composition of the AuNP.

Fig.2.5[188, 190, 193]. The EDS peaks aside from Au are from Copper (Cu) and Carbon (C), both
likely from the carbon-coated copper TEM grid we used. We did measure some trace Oxygen (O),
which is likely the result of normal atmospheric contamination during sample loading.

We used high resolution TEM (HRTEM) to measure the average diameter of the AuNPs.
Figure2.6 shows a chart of AuNP diameter (error bars reﬂect standard deviation) that we found for
various reactor conditions as listed in Figure 2.1.

The AuNPs have narrow size distribution with average diameter below 5 nm, regardless of
the reactor conditions we tested (including pressure, RF power, and reactor volume). While the
parameters we tested did not yield a major change in AuNP diameter, the small NP size is within a
promising range for use as plasmonic resonators and in other applications[45, 60, 175].

20

Figure 2.6: AuNP average diameter and standard deviation of various samples from Table.2.1, as
measured from TEM images

To understand the physical process during AuNP synthesis, we used SEM to image the gold
wire before and after the plasma synthesis under the conditions for Sample 2 (900 sccm Ar, 10 W
supplied RF power). Figure.2.7 shows the SEM images. Roughness at surface of the Au wire after
the plasma indicates physical destruction via the plasma process.

21

Figure 2.7: (a) SEM image of the consumable Au wire surface before plasma exposure. (b) SEM image
of the Au wire surface after plasma exposure (conditions as for Sample 2 in Table2.1), showing noticeable
roughness. (c) Higher magniﬁcation of SEM image of the Au wire after the plasma process.

We hypothesize that the cause of AuNP production is related to the increased wire temperature
during the plasma process. During the process, the wire glows as though it is heated, and when we
increased the power to the plasma we saw eventual melting of the wire. By inserting a thermocouple
lead into the plasma during the reaction, we were able to measure the gas temperature within a
few seconds of extinguishing the plasma (RF interference prohibited direct measurement during

22

the reaction), with temperatures around 400K. This is far from the melting poing of gold (1337 K),
consistent with the nonthermal nature of this plasma. However, resistive heating from the supplied
RF as well as electron/ion recombination at the wire could lead to increased wire temperature.
Due to the higher pressure of our reaction (100-300 Torr) operated at RF we do not suspect ion
sputtering to be a major cause of AuNP formation, since any ions produced are likely to have low
energies.

While the plasma-based method was promising regarding the AuNP size distribution, the mass
yields were low – in fact while collection of NPs on a TEM grid was facile, we were unable to
collect enough AuNPs for an accurate weight measurement. For comparison, we also performed
AuNP synthesis using the hot-wire method under low-pressure conditions (1-10 Torr), as described
in the Experimental Details section. Here we compare these two methods. In the hot wire method,
we applied a narrow range of DC power to the Au-Pt wire (10-15 W) and the system pressure was
lower than plasma-based method (several Torr versus 100-200 Torr for the plasma-based method).
The rate of production was lower for the plasma-based method, as conﬁrmed with the TEM results
from each method under the same collection time. The hot-wire method produced AuNPs with a
broad size distribution and larger average diameter (see Fig.2.8). By comparison, the plasma-based
method yields AuNPs with a smaller average diameter and narrower standard deviation, which is
attractive as many applications for these NPs rely on size-dependent properties.

23

Figure 2.8:
(a,b) TEM images of AuNPs as synthesized using plasma-based method (Ar: 900 sccm,
pressure: 180 Torr) and the hot-wire method (Ar: 900, DC power: 12 W) respectively. (c) Average AuNP
diameter and standard deviation for the plasma-based method (circular symbols, Samples 1 and 2 from
Fig.2.1) and the hot-wire method (square symbols, Samples 3 and 4 from Fig.2.3) as measured using TEM
image analysis. AuNP, gold nanoparticle; DC, direct current; TEM, transmission electron microscopy

The lack of size control for the plasma method deserves comment. We explored a wide range of
plasma powers and gas ﬂowrates, and did not see a pronounced change in average AuNP diameter.
In line with our hypothesis that the plasma process for AuNP synthesis is thermally driven, we did
not see any size change as a function of power. At low powers, the wire was not heated su ciently
to lead to AuNP production. However, the thin wire equilibrated in temperature rapidly – so once
we supplied enough power for AuNP synthesis, increasing the power led to the wire melting, not
to increased AuNP vapor and size: the available power range was small due to this, and did not

24

allow for size control. We also did not see a ﬂowrate dependence. In other ﬂow-through plasma
synthesis reactors, the size of the resulting nanoparticles is dictated by the residence time of the
nucleates/particles in the plasma[28, 58, 87, 107, 197]. However, in our process, the pressure
and ﬂowrate were not decoupled (due to using oriﬁce-based pressure control) and any increase in
ﬂowrate caused also an increase in pressure, reducing the e ect of ﬂowrate on the residence time.
Furthermore, the long residence times for our process (due to the higher pressure) meant that we
were not dynamically altering AuNP formation by changing the gas ﬂowrate. When we tried to
reduce the pressure and thus the residence time, we were not able to synthesize AuNPs.

We also calculated the residence time (RT) of the AuNPs in the reactor. This calculation is
based on Equation 2.1 below. The RT (described below) roughly estimates the time each particle
travels in the plasma and thus can give an idea of how long the particle has to grow. This equation
accounts for ﬂuid velocity and steady ﬂow conditions through the oriﬁce. This equation shows that
pressure and ﬂowrate have a direct e ect on a particle’s RT. These results show a loose correlation
between the size of the AuNPs and their RT, as we dicussed above in this section. Table 2.2 shows
calculated RT for particles characterized in Table 2.1.

ResidenceTime(RT) = (((⇡ ⇤ R2)⇤ L ⇤ ((P/Patm))))/((F(sccm)) ⇤ 60)

(2.1)

Table 2.2: Comparing plasma-based and hotwire methods

Despite the success at synthesizing AuNPs using the hot-wire method and some control over
size via process parameters, we were discouraged by material costs (namely, the Pt core wire),

25

experimental constraints, and di culties associated with the reactor geometry, as well as the wide
size distribution of the AuNPs. The plasma technique is a more direct, lower-cost method for AuNP
synthesis and gas-entrainment, although the production yield is also low enough to prompt future
studies on how to improve on this method.

Here, we presented a gas phase method which is based on an RF plasma to synthesis AuNPs.
We also produced AuNPs by hotwire method under vacuum and compared these two methods. For
either method, it is critical to understand the conditions during AuNP growth so that the method can
be extended to other noble metals, and to aid future processes based on speciﬁc AuNP properties
(for example, Si nanorod growth based on supersaturation of a eutectic alloy of Au and Si using the
AuNPs as seeds). Reaching toward future experiments using these combined gas-phase processes,
we performed preliminary experiments to analyze the AuNP temperature during formation, the gas
temperature during the plasma process, and employed methods to analyze the gas composition with
future use in estimating the electron temperature in the plasma. Those results are summarized in
this section.

Based on results presented earlier in this Chapter, future experiments on anisotropic semicon-
ductor nanostructure growth in the PNL will use the plasma-based AuNP synthesis route. As a
preliminary step towards this objective, we performed experiments to determine the exact nature
of AuNP production in this method as we brieﬂy discussed earlier [83, 84]. In one experiment,
we attempted to able to measure the ion density in the plasma using an electrical probe which is
installed in the plasma reactor. This theory has been applied in experiment as well [19, 84, 109].
A function generator is connected to a variable frequency power supply through a capacitor. The
probe is powered at RF frequencies (< 13.56 MHz) and chopped – during the “signal-o ” phase,
the capacitor is discharged by the plasma ion ﬂux to the probe. The rate of discharge is proportional
to the ion density (which is equal to the electron density under the assumption of quasineutrality in
the plasma) [19, 83, 84, 134]. Our experiments with electrical probe measurements are ongoing,
due to complications with the lower-frequency power supply which led to many false starts. While
our probe measurements are on hold until we can obtain more robust and reliable equipment, we

26

plan to use this method in the future as a means to estimate the plasma characteristics during AuNP
synthesis.

We also used Optical Emission Spectroscopy (OES) to investigate di erent excited species and
their relative intensities in the plasma system which helps to calculate species concentrations and can
give an idea of the electron temperature during the reaction. In OES measurements, a spectrometer
collects all emitted light from the plasma (ultraviolet and visible light). The lines are indicative of
various atomic transitions from excited-state species in the plasma. The ultimate objective in this
study was to then use a computer code to estimate the concentrations of various species and the
electron temperature during the process, based on comparing these characteristic line intensities.
Figure 2.9 shows a sample of OES from the Ar plasma with the gold central ground electrode,
including a trace addition of H2 gas for use in calculating electron temperature[19, 84, 109]. The
data in fact was inconclusive. We hoped to see a line from Au, if in fact we had generated any
excited-state Au atoms during our process. However, either the process was not energetic enough
to excite the Au atoms/clusters that coagulated during AuNP synthesis, or else the densities were
too low to observe. Characteristic lines from Au are expected at 406.511 nm, 460.750 nm, 479.273
nm, 627.817 nm and 751.073 nm (from NIST Atomic Spectra Database Lines Data website).
Additionally, we hoped to use the trace H2 that we added to the plasma as a means to estimate the
electron temperature, but we were also not able to clearly see evidence of the H2 optical emission,
making that calculation impossible. It is interesting that the two spectra with and without H2 (red
and blue, respectively) are qualitatively di erent from one another - and so the addition of H2 even
at low concentration (1 sccm) was su cient to shift the emission from the Ar plasma.

In Figure 2.9 we used a 340-1050 nm Ocean Optics USB spectrometer to run the AuNPs plasma

reactor with using H2.
2.4 Conclusion

We presented a new gas phase method which is based on an RF plasma to produce AuNPs in
a more sustainable path. For comparison, we also produced AuNPs by hot wire method at low

27

Figure 2.9: OES test for Plasma-Based AuNP synthesis, 340-1050 nm spectrometer with Hydro-
gen(red) and without Hydrogen(blue)

pressure. This new plasma-based approach results in a smaller average size and smaller overall
distribution of sizes per sample, as compared to the hot wire method. Among the challenges
presented by this method are a low production yield and invariant particle diameter, based on
the parameters we tested, although the average AuNP diameter of 4 nm is suitable for many
applications. Future work will explore methods to increase the production.yield as well as control
the AuNP diameter using the plasma method, along with e orts to employ this facile technique to
synthesis of other metallic NPs.

28

ESTIMATING THE MECHANICAL PROPERTIES OF SINC LAYERS ON PDMS

CHAPTER 3

Portions of this work have been published in the journal "Nanoscale" in March 2019 and are
reproduced here.

29

3.1

Introduction

As ﬂexible devices become more prevalent, it is increasingly important to measure and predict
the mechanical properties of active device layers, including thin ﬁlms [9–11]. Nanoscale com-
ponents, such as nanocrystals, can be incorporated into mechanically ﬂexible devices.[10, 152]
Silicon in its nanocrystal form has semiconducting and photoluminescence properties which make
it a promising material for technologies like microelectronics[103, 140, 163], photonics[50, 103],
solar-photovoltaics[95, 166], memory devices[146], biosensors, and light emitting devices[6, 7], as
well as other applications involving polymer matrices.[115, 173] Using silicon nanocrystals in these
applications is particularly attractive considering the natural abundance and non-toxicity of silicon.
In most of the aforementioned applications, silicon nanocrystals (SiNCs) are deposited on rigid
substrates. Depositing a SiNC ﬁlm onto stretchable or ﬂexible substrates opens up possibilities for
new device development and modiﬁcations[179]. While the opportunity of designing these devices
is exciting, it also highlights the need for an innovative way to estimate mechanical properties for
these systems.

Substantial work has been done by various groups on studying the mechanical properties of
nanocrystalline materials, both in the elastic [57, 191] and plastic [62, 172, 180, 181, 184] regimes.
Some works also discuss the signiﬁcant di erences in mechanical properties of nanocrystalline
material from their bulk counterparts [57, 62, 172, 181, 182, 184, 191, 196]. These testing
techniques, while powerful in estimating the material properties for a large variety of nanocrystalline
materials, inherently require the destruction of the samples or the plastic deformation of the material.
This is in conﬂict with the scope of this work, which is focusing on non-destructive techniques to
estimate elastic behavior of thin layers of nanocrystals on deformable substrates. These combined
systems limit in-situ measurement of the layers’ mechanical properties such as elastic modulus.
This necessitates ﬁnding an appropriate method to understand and predict the response of not just
the nanocrystal layers during deformation, but of the whole elastomer-nanocrystal layer system.
Indeed, in-situ evaluation of mechanical behavior is crucial for thin layers, for which interactions
between the substrate and the thin ﬁlm can be as important as the properties of the thin ﬁlm itself. To

30

our knowledge, such in-situ evaluation of the mechanical properties of SiNC layers on elastomers
has not been previously reported.

Prior in situ studies on nanoscale materials have employed microscopy during uniaxial tensile
testing to observe deformation under applied loads [181, 183, 184]. Barring the specialty equipment
necessary for these experiments, a promising option that has been used in the last two decades to
estimate material properties of thin ﬁlms is to measure the onset of instability patterns on the
surface of the thin ﬁlm due to di erent loading conditions [30]. This technique has the advantage
of being simple, inexpensive, and non-destructive; furthermore it can be applied to any type of
thin layer deposited on a deformable substrate. Because of its ﬂexibility, this method has been
employed to estimate mechanical properties of several isolated thin ﬁlms [157, 168, 178] or multi-
layered structures [77, 102, 125, 162]. All of these studies focused on applying homogeneous
deformations, namely uniaxial tension or compression, to the system. However, other types of non-
homogeneous loading conditions, such as bending, can also lead to the formation of instabilities.
In the past, several studies have focused on developing a theoretical framework to describe the
onset of instabilities on the surface of rectangular blocks subject to ﬁnite bending[31, 40, 41, 44,
51, 52, 143, 144, 153, 169]. The framework developed in these works focused only on ﬁnite
bending deformation for a homogeneous block. These early predictions have also been recently
conﬁrmed by experimental work [51, 52]. In the last decade, a few studies focused on predicting
bifurcations during ﬁnite bending of multi-laminated structures, speciﬁcally in bilayered systems,
both theoretically and experimentally[143, 144].

A recent study on the optical and structural properties of SiNCs on polydimethylsiloxane
(PDMS) performed at PNL and demonstrated the formation of instabilities on a nanocrystal ﬁlm
deposited on a prestretched elastomer [104]. Combining the onset of instabilities with the predic-
tions from the theoretical modeling framework described, we evaluate the mechanical properties of
the SiNC thin ﬁlm deposited on an elastomeric substrate. This theoretical framework can further
be applied to di erent combinations of nanocrystals and elastomers, for in situ measurement of
mechanical properties of such bilayered systems.

31

Figure 3.1: Steps from plasma deposition to ﬁnite bending. SiNCs are deposited through a slit-
shaped oriﬁce onto PDMS, which is then bent to observe formation of wrinkles on the SiNC
surface. (a) Photograph of plasma reactor. (b) Schematic diagram of reactor and SiNC synthesis.
(c) SiNC/PDMS sample illuminated with UV light, demonstrating photoluminescence. (d) Finite
bending applied with a lab-designed apparatus.

3.2 Methods

Here we take a novel approach to evaluate the mechanical properties of thin ﬁlms of nanocrystals
on an elastomeric substrate. A PDMS-SiNC bilayer was used as a model to present the framework.
First, we inertially impacted a uniform ﬁlm of SiNCs on PDMS. Then, we applied a ﬁnite bending
deformation to the system until instabilities formed on the surface coated with the SiNCs, noting
the critical angle at which the instabilities appeared. Finally, we applied a theoretical modeling
framework to describe the bilayer system. See Figure. 3.1 for an overview of our process.

32

3.2.1 PDMS fabrication

To prepare the PDMS as substrates for this study, a Dow Corning heat-cured prepolymer PDMS
184-kit Sylgard Elastomer with curing agent catalyst was used with a 10:1 weight ratio of base to
curing agent catalyst. Changing this mass ratio will result in PDMS substrates in di erent sti ness.
In other words, curing agent ratio will determine the mechanical properties such as shear modulus
of the PDMS polymer after curing process. This topic will be discusssed in Chapter 5. This mixture
was stirred vigorously with a mechanical stirrer until air bubbles formed, and placed under vacuum
for 10-15 minutes to remove the trapped air. The mixture was then weighed and poured into a
petri dish of 85 mm diameter and was cured for 120 minutes at 60oC by placing on a hot plate to
form PDMS with an average thickness 2.5 mm. When PDMS cured, the thickness at the edge of
the container is alittle bit higher, due to surface tenssion of ﬂuid during the curing process. After
curing, the PDMS was cut in rectangular sections according to our requirements and then attached
to two metallic handles with Gorilla® Super Glue as shown in Figure.3.1b and d.

3.2.2 Plasma synthesis of SiNCs

We produced the SiNCs using a nonthermal plasma reactor as has been reported previously[76, 80,
83, 104, 107]. This reactor, as presented in Figure 3.1,consists of a pyrex tube through which we ﬂow
94 sccm (standard cubic centimeters per minute) of Argon (Ar) and 16 sccm of silane (SiH4, 5%
in Ar), as well as 50 sccm of H2 through the sidearm for surface passivation. Radiofrequency (RF)
power at 13.56 MHz was supplied to dual ring electrodes encircling the reactor, and the pressure was
kept constant in the range of 3.75 to 3.85 Torr using a slit-shaped oriﬁce. Obviously changing the
size of the oriﬁce, will change the operating pressure of this reactor. The stando  distance between
the oriﬁce and the substrate was ⇠4 mm. In this study, we keep the stando  distance constant to
presevre same impaction for SiNCs deposited on the ﬂexible substrate. The downstream pressure
was ⇠0.500 Torr. We also mesured the pressure in the middle of the reactor, using a side arm in the
same height as for the side arm desgined for Hydrogen line. This middle pressure was ⇠1.250 Torr.
The nanocrystals were deposited on thin ﬁlms by inertial impaction directly out of the reactor onto

33

the substrate[6, 48, 69, 124]. One of the assets of this synthesis method is the room-temperature
process and ability to directly deposit SiNCs without post-processing steps or solvents, even on
temperature-sensitive substrates such as PDMS. The PDMS samples were made (as described in
Sec. 3.2.1), attached to a pushrod, and placed under the oriﬁce downstream of the reactor. The
SiNCs were deposited by rastering the substrate back and forth beneath the oriﬁce to achieve a
uniform ﬁlm of SiNCs on the PDMS. Figure 3.1b shows a work-ﬂow schematic of the reactor and
deposition of SiNCs on the PDMS. The rastering duration was 20 minutes with each back and forth
movement lasting 3 seconds, resulting in a SiNC layer thickness of ⇠4.5 µm.

3.2.3 Silicon nanocrystal ﬁlm characterization

The thickness of the sample was measured by scanning electron microscope (SEM) imaging of
a cross-section of the sample generated using a focused ion beam (FIB). Figure.3.2a shows an
SEM image of the sample, and Figure.3.2b shows the FIB-generated cross section for thickness
estimation. To decrease charging e ects and improve the image quality of the SiNC ﬁlm, a thin
layer of Platinum (Pt) (⇠6 nm) was deposited via Ar plasma sputtering prior to imaging.To perform
a reasonable FIB cross sectioning, we need to roughly estimate the thickness of the ﬁlm prior to
milling and ﬁnal imaging. As we mentioned in previous section, we started our FIB milling with
an estimation of 4.5 µm. The speed of FIB cross sectioning depends on the thin ﬁlm structure and
its thickness. The FIB image presented in Figure 3.2 took around 10 mins. Another challenge in
this thickness measurement was due to similarity of contrast of PDMS and SiNCs. In other word,
due to the existing Silicon in the structure of PDMS and SINCs, the contrast at the interface of
PDMS, as the substrate and SiNC layers, as the thin ﬁlm, is negligible and need high resolutiom
for conﬁrming proper estimation of the interface as presented in Figure 3.2. X-ray di raction
(XRD), shown in Figure.3.3a, conﬁrms the formation of crystalline nanoparticles. Transmission
electron microscopy (TEM) was performed on the SiNCs using a JEOL 2200 instrument, showing
individual NCs (Figure.3.3b). The TEM sample was prepared by tapping the TEM grid on the
surface of the sample to collect some of the SiNCs. The size of the SiNCs was measured using

34

TEM image analysis, giving an average SiNC diameter of ⇠3.7 nm.

Figure 3.2: SEM and FIB cross-sectioning of the SiNC layer. The thickness is estimated to be
(⇠4.50 µm). (a) SEM image showing the surface of the SiNC layer on PDMS, including the FIB-
milled area (inset). (b) Higher magniﬁcation SEM image of the FIB-milled area demonstrating the
thickness of the SiNC layer.

35

Figure 3.3: (a) X-ray di raction conﬁrmed formation of SiNCs. XRD for SiNCs shows crystal
reﬂections at 2✓ values 28.3, 47.3 and 56.1 degrees consistent with the (111), (220), and (311)
planes of silicon.
(b) TEM image conﬁrming the formation of SiNCs: black circles highlight
individual SiNCs. The inset shows FFT of the correspondig SiNCs of the TEM image.

3.2.4 Uniaxial tensile test

We applied a standard uniaxial tensile test to evaluate the mechanical properties of the PDMS.
A rectangular sample of PDMS was mounted on a custom built uniaxial tensile test machine and
four black dots were applied on it as ﬁducial markers. Figure 3.4 shows a PDMS mounted to the
load cell for performing unaxial tensile test. The sample was then stretched cyclically for 5 cycles
from a conﬁguration with a 10g pre-load ( reference conﬁguration) until a 20% deformation was
achieved. The sample was stretched between two clamps, one of which was attached to a load cell
to measure the force on line with the applied deformation. The strain rate used during the tensile
test was 0.1 mm/s. The force from the load cell was converted to engineering stress by calculating
the average cross-section area of the sample in the reference conﬁguration. A Hitachi KP-M2AN
CCD camera was used to capture and feed the image of the ﬁducial markers during uniaxial stretch
into LabVIEW.

36

Figure 3.4: The stress machine loaded with a colored PDMS and dots depicted on the PDMS,
dislocation of dots will be recorded by a digital camera

Figure 3.5: (a and b)A schematic diagram of proposed aparatus for ﬁnite bending measurement
desgned by Solid Work software and (c) Alluminum cast aparatus designed at PNL

3.2.5 Finite bending experiment

We designed a ﬁnite bending apparatus similar to the one used previously by Roccabianca
et al.(2010)[143]. A schematic diagram of this aparatus shown in Figure 3.5. Two mainpa-
rameters in this design are slenderness ratio (⇤ = l0/h0) which is the ratio of width to thickness
of the sample in cross section, and thickness ratio (H = hNC
0 ) which is the ratio of thickness
of the SiNC layer by that of the PDMS before ﬁnite bending (see Fig. 3.6a for reference). The
third dimension is taken to be 8-10 times larger than l0 to achieve a plane-strain condition on the
cross section.In other words, due to the size ratio of third dimension to the other two dimensions of
the substrate,the third dimension does not have an impact on the formation of surface instabilities.
Here (and in all following equations) the superscript ‘P’ refers to the PDMS and ‘NC’ refers to the
SiNCs.

0 /hP

37

Figure 3.6: PDMS sample in the ﬁnite bending apparatus. (a) Side view, showing sample (with
handles) in the ﬁnite bending apparatus. (b) Top view showing formation of bifurcations.

The bilayer PDMS/SiNC system was put in the ﬁnite bending apparatus (see Fig.3.6) and
progressively bent until the SiNC ﬁlm underwent bifurcation, at which point the critical angle was
recorded. Due to higher critical angles of bifurcation observed for samples with higher slenderness
ratios, deformation was applied by hand instead of using the apparatus. The onset of bifurcations
was detected by direct visual inspection, and the critical angle was measured using analysis of
photographs of the side view of the sample. Each measurement was performed three times to
ensure accurate and reproducible readings.

3.2.6 Modeling bifurcation in elastic layered structures

The following sections pertaining to theoretical modeling of the behavior of the bilayer system
were largely executed by M. Sinha and Dr. S. Roccabianca, our collaborators on this project. For
completeness they are included here.

Roccabianca et al. (2010), formulated a theoretical framework to study incremental bifrucations
due to ﬁnite bending of an elastic incompressible multilayered block[143]. This theory was
developed for an N-layered system for Mooney-Rivlin type of material. The authors also performed
some experiments in bilayered structures that conﬁrmed their theoretical predictions. Here, we
employ a similar theoretical modeling framework for the bilayered system to estimate the ratio of
mechanical properties of SiNC thin ﬁlm to that of PDMS.

38

3.2.6.1 Finite bending

The ﬁnite bending deformation is prescribed such that the reference conﬁguration is rectangular
and described in a Cartesian coordinate system, and the deformed conﬁguration is a sector of ring
described in a cylindrical coordinate system (see Fig. 3.7).

Figure 3.7: Reference and deformed conﬁgurations for the PDMS-SiNC bilayer. The bilayer de-
forms from the Cartesian coordinate system in the reference conﬁguration (e1, e2, e3) to a cylindrical
coordinate system in the deformed conﬁguration (er, e✓, ez). Speciﬁcally, a plane at constant x0
1
transforms to a plane at constant r (dashed line), and a plane at constant x0
2 transforms to a circular
arc at constant ✓ (dashed-dotted line). Since the out of plane deformation is taken to be zero, we
have x0

3 = z.

Brieﬂy, the ﬁnite bending deformation can be described, within both the PDMS and the SiNC

layers, by the following deformation gradient F[143]

F =

l0
2 ¯✓r

er ⌦ e0
1 +

2 ¯✓r
l0

39

e✓ ⌦ e0

2 + ez ⌦ e0
3,

(3.1)

where l0 is the overall width of the sample and 2 ¯✓ is the angular deformation of the sample.
We consider the volume of the system to be conserved and the two layers to be perfectly bonded.
Therefore, we can apply the interface condition r NC
i are the outer radius
of the SiNCs layer and the inner radius of the PDMS, respectively. This allows us to write,

i , where r NC
o

o = rP

and rP

r NC
o

i  hP
0 , hNC,↵  = rP
 hNC
0 , hP,↵ ,

(3.2)

where hm

0 and hm represent the thickness of each layer before and after bending (m = NC, P),
and ↵ = 2 ¯✓/l0. To predict all geometrical characteristics of the deformed state, we consider all the
layers to be equilibrated. Brieﬂy, we deﬁne the Cauchy stress in each layer as,

T =  ⇡I + F

dW
dC

FT,

(3.3)

where C = FTF and I are the right Cauchy-Green deformation tensor and the identity tensor,
respectively, and W and ⇡ are the strain energy function and the Lagrange multiplier for each layer,
respectively. We have employed the neo-Hookean strain energy function to describe both layers,
µ0
deﬁned as W =
2 (trC   3). Note that we are deﬁning two constitutive equations, one for the
PDMS layer as a function of the material parameter µP
0 , and one for the SiNC layer as a function
of the material parameter µNC

.

0

For the layers to be equilibrated, the equilibrium equation in absence of body forces is required
to be satisﬁed within each layer. Furthermore, boundary conditions of zero traction are applied at
the innermost and outermost surfaces, as well as an interface condition of perfect bonding between
the two layers.

3.2.6.2 Bifurcation analysis

The loss of uniqueness of the plane-strain incremental boundary value problem has been investi-
gated for multi-layered structures in previously published work[143]. Here we have adapted the
formulation to ﬁt our application for a bilayered system of PDMS and SiNCs. Brieﬂy, for each layer,
we seek to solve the incremental counterpart of the equilibrium, div(⌃) = 0, where ⌃ = €SFT is the

40

incremental ﬁrst Piola-Kirchho  stress tensor, and the ﬁrst Piola-Kirchho  stress and the Cauchy
stress tensors are related by S = T F-T for incompressible materials. Here, and in the following, we
use a superimposed dot to represent incremental quantities. The linearized constitutive equation
for each layer is given by

⌃ =   €⇡I + CL

(3.4)

where C is the fourth-order tensor of instantaneous elastic moduli and L = grad u is the gradient
of incremental displacement u(x). We consider bifurcations to be represented within each layer
by an incremental displacement ﬁeld in a separable-variables form, as described by Roccabianca
et al. (2010). For incompressible isotropic elastic materials, one can write C as a function of two
incremental moduli, namely µ and µ⇤. For neo-Hookean materials subject to ﬁnite bending in
plane strain condition, the elastic moduli can be written as

µ = µ⇤ =

µ0

2 h(↵r)2 + (↵r) 2i

(3.5)

Finally, the incremental boundary value problem is completed by a set of boundary and interface
conditions. Speciﬁcally, zero normal and shear incremental stresses at the innermost and outermost
layer, zero incremental shear stress and incremental radial displacement at the boundary ✓ = ± ¯✓,
and continuity of incremental stresses and displacements across the interface. To calculate the
angle that satisﬁes the bifurcation condition, ✓cr, we employed the compound matrix method, as
described in a previously published study [144].

3.3 Results and Discussion

Experimentally, the critical angles of bifurcation were measured for nine samples by performing
the ﬁnite bending experiment described in Sec. 3.2.5. The values, reported in Table 3.1, have been
obtained for three experimental groups, characterized by a slenderness ratio ⇤ = 2, 3 and 5, each
group formed by a total of three samples (see Fig. 3.8).

41

Figure 3.8: Critical angles of bifurcation evaluated as a function of ratio of material properties
for a PDMS-SiNC bilayer sample. The ratio of the neo-Hookean material parameters is deﬁned
as ¯µ0 = µNC
0 . All curves have been calculated for a thickness ratio H included in the range
0.0013–0.0018. Each line represent a di erent value of slenderness ratio, ⇤ = 2, black solid line;
⇤ = 3, black dashed line; ⇤ = 4, dark gray solid line; ⇤ = 5, dark gray dashed line; ⇤ = 6, light gray
solid line. The insert represents a schematic of the bilayer and deﬁnes the slenderness ratio.

0 /µP

Slenderness ratio, ⇤ Critical angle, ✓cr Material parameters ratio, ¯µ0
2
3
5

47.33 ± 3.21
71.33 ± 1.53
130.67 ± 3.21

1.90 ± 0.13
1.85 ± 0.05
1.68 ± 0.04

Table 3.1: Critical angles of bifurcation for di erent slenderness ratios, expressed in degrees.
Angles are reported for ⇤ = 2, 3 and 5.

Employing the model described in Sec. 3.2.6, we evaluated the critical angle of bifurcation
for the bilayered structure, shown in Fig. 3.8 and Fig. 3.9. Speciﬁcally, after ﬁxing the overall
geometry of the sample, namely the thickness ratio H and the slenderness ratio ⇤, we calculated the
value of the angle representing the onset of bifurcation for each value of the neo-Hookean material
parameters ratio, deﬁned as ¯µ0 = µNC
In both Figures, we considered ¯µ0 to be included

0 /µP
0 .

42

Figure 3.9: Plotting points from the experiments on graph obtained by numerical calculations. The
ratio of the neo-Hookean material parameters is deﬁned as ¯µ0 = µNC
0 . Shown here are points
obtained for ⇤ = 2, 3 and 5. Insert shows an instability on the surface of a sample with ⇤ = 2.

0 /µP

between 0 and 100, however a pilot study showed that the value of the critical angles plateaus for
values of ¯µ0 included between 102 and 105 (data not shown). Furthermore, the same pilot study
showed that the critical angle of bifurcation is independent of the thickness ratio H, if H is included
in the range 0.0013 – 0.0018. The experimental value measured in this study is included within
that range, speciﬁcally H = 0.0017 on average for the nine samples.

Figure 3.8 shows the variation of the critical angle of bifurcation for varying values of the
slenderness ratio, speciﬁcally for ⇤ included in the range 2 to 6. For a ﬁxed value of ¯µ0, which
is represented by a vertical line in the graph, an increase in the slenderness ratio results in a later
appearance of the bifurcation, namely a higher value for the critical angle. This trend is consistent
with previously published results[143, 144].

Figure3.9 shows how we employed the onset of bifurcation to estimate the mechanical properties

43

of the SiNC layer. Brieﬂy, for each value of slenderness ratio considered in the experiments, we
calculated the correspondent theoretical sets of bifurcation angles. Then, the critical angle measured
experimentally was employed to uniquely identify the corresponding ¯µ0. The experimental value of
the critical angle ﬁxed the abscissa of the point in the ✓cr – ¯µ0 plane, represented in Fig. 3.9. Then,
we imposed that the experimental point belonged on the bifurcation threshold curve calculated for
the corresponding slenderness ratio, which in turns ﬁxed the ordinate of the point which is ¯µ0. The
value estimated for each experimental group, namely for each value of slenderness ratio considered,
is reported in Table 3.1. A vertical dashed line in Fig. 3.9 represents the average value of ¯µ0 for
all the experimental data, which is 1.81 ± 0.12. We then estimated the value of µP
0 = 190 kPa
by ﬁtting the data collected from the uniaxial tensile test, as described in Sec. 3.2.4. This gave
us a value for the neo-Hookean material coe cient for the SiNC layer of µNC
0 345 ± 23 kPa. The
neo-Hookean material constant is analogous to the shear modulus for low values of applied strain.
This ﬁnding is a clear indicator that the mechanical properties of this materials system are
determined by the interactions between SiNCs within the NC layer and between the NC layer and
the PDMS. Direct measurements of the mechanical properties of individual silicon nanocrystals
via microscope-guided nanoindentation have yielded elastic moduli around 170 GPa,[53] several
orders of magnitude higher than our measurement of the shear modulus for the SiNC layer on
PDMS (assuming incompressible behavior, shear modulus is ⇠1/3 of the elastic modulus). Others
have produced elastic/Young’s modulus estimates for single SiNCs between 70-170 GPa.[82, 119,
149] This range of values (⇠20-60 GPa for shear modulus) is also consistent with mechanical
measurements of thin ribbons of crystalline silicon on PDMS [82]. By contrast, our value of 345
kPa for shear modulus is far lower, and is consistent with the higher porosity of our nanopowder
layer as compared to bulk silicon. For example, Gross et al. showed that increasing porosity of
powder hydroxyapatite coatings leads to decreasing elastic modulus - further, regardless of porosity,
the coatings exhibited values of elastic modulus more than 35x smaller than the modulus of single-
crystal hydroxyapatite. [59, 192] Similarly, the Young’s modulus of aluminum foams is between
10 1 and 10 3 the value of Young’s modulus of solid aluminum, depending on porosity.[2] Finally,

44

previous results have shown that increased porosity signiﬁcantly decreases the Young’s modulus of
polymer thin ﬁlms deposited on PDMS, when measured by employing mechanical instabilities[155].
The porosity of our layers was di cult to ascertain with high accuracy. However, a preliminary
density measurement provided us with a value near the limit for close-packed randomly oriented
rigid spheres (volume fraction 0.64). While we did not probe the e ects of porosity in this
proof-of-concept investigation, the porosity of the SiNC layer is tunable by adjusting reactor
properties,[48, 69] and future work will be devoted to determining the relationship between layer
porosity and mechanical properties. This versatile method is broadly applicable to extract the
mechanical properties of nanoparticle layers on elastomer substrates regardless of material, layer
thickness, and density. Our results also o er a unique opportunity in the future to investigate
how the NC surface functionality (organic ligand coating, inorganic shell, etc.)
inﬂuences the
mechanical behavior of the NC ﬁlms.

3.4 Conclusion

In conclusion, we have used a new method based on the onset of bifurcation to estimate
the mechanical properties of thin layers of nanocrystals on PDMS by coupling the experiment
and theoretical modeling.
In general this method, opening up new possibilities for estimating,
and controlling the mechanical behavior of nanoparticle layers on stretchable or bendable device
technologies, leveraging both the unique optoelectronic properties of nanoparticles together with
the mechanical versatility o ered by these ﬁlms.

Finally, the results obtained from this experiment are exciting since previous attempts in in-
vestigating instability formation due to ﬁnite bending in a multilayered structure were limited to
structures with elastomeric layers only. Here we have reliably shown formation of instabilities in
thin SiNC ﬁlms deposited on PDMS as a ﬂexible substrate. This brings opportunities to study
and develop constitutive models which are better suited to such constituents, in addition to other
experimental methods that can perform non-destructive and in-situ measurements of this class of
multilayer systems. This versatile method is broadly applicable to extract the mechanical proper-

45

ties of nanoparticle layers on elastomer substrates. In Chapter 5 we apply same methodology to
investigate a parametric study regarding the inﬂuence of a few properties of the SiNC layer and the
PDMS and how they a ect the mechanical properties of the SiNC thin ﬁlms. Also this method,
regardless of material, layer thickness, and density can be applied for a range of nanomaterial ﬁlms
and substrates.

46

OPTICAL AND MECHANICAL BEHAVIOR OF SINC-EMBEDDED PDMS SLABS

CHAPTER 4

47

4.1

Introduction

Luminescent nanocrystal inclusions in polymer hosts have received considerable interest re-
cently, in particular due to their promise in novel optoelectronic device applications such as pho-
tovoltaics based on luminescent solar concentrators (LSCs), electrochromic windows, and more
[23, 65, 89, 115, 116, 141, 165, 174, 186]. The combination elastomeric behavior of polymers
together with the tunable optical absorption and emission from these nanocomposites also opens
avenues to stretchable and ﬂexible manifestations of these same applications, as well as wearable
electronics such as biosensors for health monitoring systems, and more. While many groups
have studied the optical behavior of these nanocomposites, their mechanical behavior is not well
understood, although a few recent studies have made advances in uncovering the e ects of the
nanocrystal inclusions [23, 66]. Modiﬁcations to the mechanical behavior of the polymers upon
inclusion of the nanocrystals is especially important for elastomeric polymers with intended use in
ﬂexible or stretchable applications. The inclusions can change the overall mechanical properties of
the composites, depending on a range of factors such as the nanocrystal concentration, dispersion
within the host polymer, the nanocrystal surface properties, and more. Going beyond informing
the deformability of the composites, tests on the mechanical properties of the composites can also
provide information about the interaction between the inclusions and the polymer, and the ways in
which applied stretching or bending can a ect the other physical properties of each.

A particularly compelling materials system is silicon nanocrystals included in polydimethyl-
siloxane (PDMS). Silicon nanocrystals (SiNCs) are made from an abundant and nontoxic material,
and exhibit exciting tunable optical properties such as photoluminescence[3, 5, 27, 36, 107, 113, 114,
154, 160, 194]. PDMS is a commonly-used silicone-based elastomer used in biomedical, industrial,
microﬂuidic, and optoelectronic applications [27, 88, 99, 110, 112, 123, 131, 148, 151, 177, 195].
As a host for nanocrystals, PDMS o ers easy fabrication, tunable mechanical behavior, biocom-
patibility, and high translucency in ultraviolet through near infrared spectral range [93]. Recently
SiNCs embedded in polymers such as polymethyl(methacrylate) (PMMA) and PDMS have been
proposed as agents for LSCs, solar cells, UV-blocking layers, and more due to the NCs’ absorbance

48

in the ultraviolet (UV) – blue spectral range and e cient emission at red-near infrared (NIR) wave-
lengths. Understanding the inﬂuence of the SiNCs on the mechanical behavior of the composite
NC/polymer system as a whole is important for designing future applications based on this materials
system, as well as other NC/polymer combinations.

A recent report by Chen et al.indicated a modiﬁed rupture stretch and mechanical properties
(tensile strength and Young’s modulus) for SiNC-infused PDMS, depending on the surface func-
tionality of the SiNCs. That group also demonstrated the UV-blocking ability of the SiNC/PDMS
composite system, highlighting the multifunctionality of these materials. We sought to further
characterize the inﬂuence of surface termination on the mechanical properties of SiNC/PDMS com-
posites by using uniaxial tensile testing to evaluate the neo-Hookean coe cient of the composite
system, depending on both nanoparticle concentration within the nanocomposite and nanoparticle
surface functionality as modiﬁed using a solution-phase thermal hydrosilylation technique. The
neo-Hookean coe cient is analogous to the shear modulus of a material, at low values of applied
strain. Our results indicate that, while all nanocomposites exhibited photoluminescence (PL), the
SiNC concentration and surface functionality play a profound role in the mechanical properties of
the composite.

4.2 Experimental Work

4.2.1 Plasma synthesis of SiNCs

SiNCs were synthesized with a non-thermal low-pressure plasma reactor as described previously
[7, 13, 22, 56, 104, 107, 118, 120, 126]. The reactor consists of a pyrex tube with dimensions of
0.25” OD at the reaction zone, expanding to 1” OD in the plasma e uent. Dual ring electrodes
bearing radiofrequency (RF) power at 13.56 MHz encircled the tube, with power delivered at 30W
via a lab-built impedance matching network. Argon (Ar, 96 standard cubic centimeters per minute,
sccm) and silane (SiH4, 5% in Ar, 16 sccm) are ﬂown through the reactor at a pressure of 3.6-3.8
Torr as controlled using an oriﬁce near the exhaust. Hydrogen gas (H2, 50 sccm) is injected just
below the onset of the expansion region in order to provide more thorough surface passivation

49

[7]. SiNCs are nucleated and grown during the reaction and collected at the reactor exhaust onto
stainless steel mesh ﬁlters. The ﬁlters were weighed before and after the reaction to measure the
weight of SiNCs collected, which was on average 3.2mg for the 12 minutes collection time used for
each batch. A section of the mesh ﬁlter was removed for FTIR analysis, before the remaining mesh
ﬁlter loaded with SiNCs was placed into a test tube, capped with a rubber septum, and ﬂushed with
N2 to ensure an inert atmosphere.

4.2.2 Functionalization of SiNCs

Thermal hydrosilylation was performed for all F samples using a nitrogen-purged Schlenk line
[106]. The SiNCs were mixed with a dried and degassed solution of 5:1 v/v mesitylene and
1-dodecene (Aldrich) and sonicated for 5 minutes. The resulting cloudy yellow solution was
transferred via cannula to a round-bottom ﬂask with a reﬂuxer and heated in a sandbath heater to
205-210°C for 2-hours. Following this process, the SiNC solution became clear indicating surface
stabilization with the dodecyl ligands. The resulting solution was dried under N2/vacuum for 24
hours at a temperature of 60°C to evaporate the residual solvent mixture. 12 mL hexane was added
to the round bottom ﬂask and sonicated for 5 minutes to ensure proper redispersion. Next, 2 mL
of the solution was placed in a vial for PL spectroscopy and the remaining 10 mL was ready for
nanocomposite formation.

4.2.3 PDMS preparation

We made the PDMS using a previously documented method [110]. A Dow Corning heat-cured 184
Sylgard Elastomer PDMS kit was used with a 10:1 w/w ratio of base to curing agent. The SiNCs
with known mass and concentration in hexane were added to the PDMS pre-polymer solution.
The mixture was stirred vigorously and then placed in a vacuum desiccator to remove any air
from the mixture and most of the hexane evaporated during this process. The resulting composite
mixture was poured into an aluminum mold with a rectangular shape and deﬁned dimensions. The

50

PDMS-SiNCs composite was cured on a hot plate at 60 °C, typically reaching ﬁnal cure in 3 hours.
Once the samples were cured they removed using tweezers.

4.2.4 Uniaxial tensile testing

We performed a standard uniaxial tensile test as reported previously to evaluate the mechanical
properties of the plain as well as the nanocrystal infused PDMS [154] A rectangular sample was
mounted on a custom built uniaxial tensile test machine and four black dots were applied on
it as ﬁducial markers. The sample was then stretched cyclically for 5 cycles from a reference
conﬁguration, until a stretch of 1.2 was achieved [154]. This stretch is equivalent to a tensile
strain of 20%. The sample was stretched between two clamps, one of which was attached to
a force transducer to measure the force on line with the applied deformation; simultaneously,
the deformation was measure by tracking the ﬁducial markers using a CCD camera (Hitachi KP-
M2AN). The strain rate used during the tensile test was 1 mm/s. The force was converted to Cauchy
stress by evaluating the average cross-section area of the sample in the reference conﬁguration, and
the change in cross section for the deformed conﬁguration. The samples were then stretched in the
same setup at a strain rate of 1 mm/s until rupture.

4.3 Results

4.3.1 SiNC characterization and nanocomposite formation

SiNCs were produced using a nonthermal plasma technique, as reported previously and describeed
in the Experimental Details section below [5, 104, 107, 154]. We characterized the SiNCs using x-
ray di raction (XRD) and transmission electron microscopy (TEM) (Fig. 4.1). The average size of
a single SiNC was estimated using TEM image analysis to be⇠4 nm in diameter and conﬁrmed to be
around⇠3.7 nm using Scherrer broadening analysis of the XRD peaks [56, 154]. In formation of the
nanocomposites, we considered two separate surface treatments: non-functionalized SiNCs (NF,
which have a native hydride coating from the plasma synthesis process), and functionalized SiNCs
(F, which have alkylated surfaces following a thermal hydrosilylation process with 1-dodecene as

51

Figure 4.1: SiNC Characterization: XRD (a) shows crystallographic reﬂections for nanosized
silicon. Gray lines are the actual data and the black trace is the peak ﬁt. TEM imaging (b) indicates
spherical nanocrystals and lattice spacing consistent with crystalline silicon. Inset shows FFT of
image.

functionalizing ligand and then dried). Photographs of the NF and F samples are shown in Figure
4.2a-b under room light and ultraviolet (UV) excitation, demonstrating visible red luminescence. PL
spectra from solutions of NF SiNCs and F SiNCs are presented in Figure 4.2. Solution PL quantum
yields from these samples were 1.6% and 54%, respectively, as measured using an Ocean Optics,
Inc. USB2000 spectrometer and integrating sphere and excited using a 395nm LED excitation
source. The NF and F samples were mixed into a small volume of hexane (12 mL, of which 10 mL
was used for fabricating the nanocomposites) and then added to the PDMS pre-polymer solution and
the curing agent – 1:10 w/w ratio for all samples – then vigorously mixed using a mechanical rotary
stirrer for 1 minute. Air bubbles and residual hexane were removed using vacuum desiccation and
the PDMS was poured into casts and cured at 60°C (more details in Section 4). We examined two
di erent concentrations of SiNCs/PDMS – 0.03 wt% and 0.12 wt%. Henceforth we will refer to the
lower and higher concentration samples with the preﬁx L- and H-. Speciﬁcally we characterized 4
sample sets and one control: L-NF, L-F, H-NF, H-F, and PDMS (no nanocrystal inclusions).

52

Figure 4.2: SiNCs in solution, (a) photographed in room light and b) under UV excitation Shown
are hexane, NF, and F samples from left to right.
(c) PL spectroscopy shows much brighter
luminescence for surface-functionalized (F) SiNCs as compared to nonfunctionalized (NF) SiNCs
in solution.

4.3.2 Nanocomposite characterization

4.3.2.1 Photoluminescence

Photographs of representative control PDMS and nanocomposite slabs in room light and under
ultraviolet (UV) excitation are shown in Figure 4.3a-b. We also measured PL quantum yields from
the nanocomposites, which are plotted in Figure 4.3c. These PL QYs underscore the high optical
quality of the nanocomposites, at 1.2% and 16% on average for H-NF and H-F samples.

4.3.2.2 Microscopy

Based on PL and visual inspection, it was not evident how the SiNCs were dispersed within the
PDMS. To investigate the dispersion of SiNCs in the nanocomposite, we ﬁrst assigned a speciﬁed
spatial Cartesian coordinate system (Figure 4.4a) to each slab, sliced the slabs to obtain thin
cross-sections, and then performed microstructural analysis using a Nikon AIRsi Confocal Laser
Scanning Microscope (CLSM), in which thin slices can be imaged with white light as well as

53

Figure 4.3: (a) Photographs of SiNC/PDMS nanocomposites in room light and (b) under UV
excitation (c) shows the average PL QY of SiNCs in hexane and in the nanocomposites, with error
bars reﬂecting the standard deviation of the measurements.

imaged at selected optical wavelengths using a laser excitation at 405 nm. The CLSM allows
visual detection of the SiNCs and clusters of SiNCs, which are otherwise di cult to discriminate
from the PDMS matrix using traditional optical or electron microscopy. The confocal nature of
the instrument allows depth resolution of the luminescent features as well, so that one can view a
composite 2-D image of 80 imaging planes (in X-Z, slice depth along the Y direction, Figure 4.4b,
or a 3-D depth-resolved map of the dispersion (Fig.4.4c). For this 3-D map, the position of each
SiNC cluster is represented in its physical location within the map as well as using color-coding
for depth. Based on this 3-D representation we found that, regardless of surface functionality, the
SiNCs are clustered together to some degree. This optical measurement technique is resolution-
limited, and so we were not able to visualize the degree of unclustered dispersion of the NCs within
the sample, but the results are important as they demonstrate the di erences in clustering behavior
of the samples.

Figure 4.5 shows representative composite images for each sample type, all acquired under
identical microscope settings. The depth of each image is ⇠60µm, and the representative images in

54

Figure 4.4:
(a) For reference during microscopy, we assigned each nanocomposite the same
coordinate system, (b) A single CLSM image of a slice is shown and (c) the depth-resolved
composite image (scale inset), demonstrates the distribution of small clusters within the sample
slab.

Figure 4.5 are a composite of 80 separate depth scans along that 60µm thickness. Visual inspection
of the CLSM images indicates an increase in number of SiNC clusters at higher concentrations
(HF and HNF) as compared to the lower concentration samples, which is not surprising. To
quantify any di erences between the samples, we analyzed the cluster size for each sample type
using a particle counting/volume algorithm via the ImageJ processing platform. We analyzed 3-4
individual images from di erent individual samples for each sample type using this algorithm to
build a comprehensive picture of the SiNC dispersion characteristics.

The results are plotted in Figure 4.6 using a box plot to display the ﬁve-number summary of
the data. First, we note that the smallest cluster size for our data was resolution-limited, and so the
lower bounds of each sample set are nearly identical. These lower bounds represent not the true
smallest cluster sizes but instead the smallest size that is resolvable using the CLSM instrument
and the processing algorithm. Aside from the increase in number of clusters, the data suggests that

55

Figure 4.5: Representative CLSM images of the SiNC/PDMS nanocomposites

Figure 4.6: SiNC cluster size statistics based on image processing from the CLSM images. 3
separate images were analyzed for each sample type. A key to the data representative is shown as
inset.

the average cluster size and median cluster size do not signiﬁcantly depend on the concentration
of SiNCs in the polymer, at least over the concentration range we investigated. This indicates that

56

for the SiNC concentrations that we studied, the mixing and curing procedure we used is su cient
to ensure relatively uniform dispersions. In other words, the distance between clusters within the
polymer host at our “high” concentration of SiNCs is still so large that the cluster formation and
dispersion is largely una ected. It is worth noting, however, that the larger particle concentrations
sometimes required slightly modiﬁed PDMS curing procedures, namely extended time on the hot
plate. At the 0.12 wt% concentration, which is modest by nanocomposite standards, the SiNC
clusters still seemed to interfere with the curing process to some degree, occasionally leading to
longer required curing times. Next we examine the e ect of surface functionality of the SiNCs
on the dispersions. One notable di erence between NF and F samples is the average cluster size.
Visual inspection of the CLSM images indicates that the clusters for the NF samples were larger than
for the F samples. The median cluster sizes as measureable using CLSM were nearly equivalent,
between F and NF samples, but the spread of data (both the interquartile ranges and the outlying
points) was larger for NF samples at both low and high concentrations. The mean cluster size of the
SiNCs in the PDMS, represented by the X symbols in the plot, is also larger for the NF samples as
compared to the F samples at both concentrations, although the discrepancy appears smaller at the
lower concentration. This ﬁnding is consistent with the fact that surface functionalization causes
a greater separation between individual SiNCs as well as improved integration with the monomer
chains before polymerization.

4.3.2.3 Mechanical properties of the composite

To test the e ect of the concentration and surface functionality of the SiNC inclusions on the
nanocomposite mechanical behavior, we used uniaxial tensile testing and a material model to
estimate a characteristic mechanical parameter, and we also quantiﬁed rupture stretch. The resulting
stress-stretch data (Figure 4.7a) can be analyzed to extract the neo-Hookean coe cient (µ, with the
dimension of a stress) based on the following equation for Cauchy stress ( ):

  = µ ⇤(  2  ( 1/ ))

57

(4.1)

Figure 4.7: Mechanical properties of the nanocomposites. (a) shows a representative plot of Cauchy
stress vs. stretch data for each sample type. From this data, (b) we extracted the neo-Hookean
coe cient for each sample, showing a distinct reduction for the higher-concentration sample. (c)
the rupture stretch was una ected by the SiNC inclusions.

where   is the applied stretch. Figure 4.7b shows the estimated neo-Hookean coe cient of each
sample, with standard deviations plotted as error bars. The rupture stretch, shown in Figure 4.7c,
has a value of 2 for all samples including the control PDMS. However, the results for µ (Figure
4.7b) tell a di erent story: while the µ values for the L-NF and L-F samples are within the statistical
variation of the control PDMS, the H-NF and H-F samples both exhibit a reduced µ, correlating to
a reduced sti ness. This reduction is most prominent for the H-NF samples.

4.3.2.4 Discussion

These results have signiﬁcant meaning. From an optical standpoint, inclusion of surface-alkylated
SiNCs in PDMS via this procedure leads to brightly emissive nanocomposites with PL QYs around
16%, pointing to the promise of this materials system going forwards. The non-functionalized
SiNCs still result in emissive nanocomposites, but the PL QY of these is quite low (⇠1-2%) and is
di cult to measure at the lower concentration we studied. For optically emissive nanocomposites,
using the surface-functionalized SiNCs is critical. The CLSM images make it clear that clustering
of SiNCs within the PDMS is prevalent across sample type. Preventing this clustering e ect using
modiﬁed nanocomposite fabrication procedures or by tuning the surface of the SiNCs by using
other organic ligands with alternate terminal groups are promising routes to explore, for even

58

higher-e ciency photoluminescence from the nanocomposites. For example,Hill et al. revealed
that the ligand terminal group greatly inﬂuenced the dispersion of SiNCs in a PMMA host [65].
This calls for a comprehensive investigation of the interaction between the ligand termination and
the resulting dispersion of SiNCs in PDMS.

Mechanically, our ﬁrst conclusion is that inclusion of SiNCs at very low concentration (0.03
wt%) in the PDMS does not appear to inﬂuence either the rupture stretch or the modulus of the
PDMS nanocomposites. At low concentration, the SiNC cluster sizes within the nanocomposites
were widely dispersed, and despite a change in average cluster size between NF and F samples,
the low concentration did not inﬂuence the overall mechanical behavior of the nanocomposites
(although it did lead to bright PL in the case of the LF samples). Second, at higher concentrations
both the NF and F samples demonstrated a signiﬁcant reduction in the neo-Hookean coe cient µ and
the reduction is more prominent for the NF samples. When compared with the nonexistent change
in µ for the lower concentration samples, this indicates that the modulus of the nanocomposites
depends both on the surface functionality and on the SiNC concentration in the nanocomposite.

We were cognizant of the fact that mixing the SiNCs into the PDMS prepolymer solution might
be one cause for the large cluster sizes and the reduction in modulus. To test this hypothesis we
also used a shear mixer (FlackTek SpeedMixer at 3000 rpm, 1 minute) for mixing H-F samples into
PDMS. Our uniaxial tensile testing results and CLSM imaging from the samples prepared using
shear mixing do not indicate any signiﬁcant deviation from the samples prepared as described
above. This points to solubility as the major cause for cluster formation and modiﬁed µ for
the nanocomposites, rather than incomplete mixing. The reason behind the reduction in the
neo-Hookean coe cient for the high-concentration composites is cause for further consideration.
Previously, we estimated µ for non-functionalized SiNC layers deposited on PDMS and found a
value of 345 kPa, approximately twice the value for PDMS. While the parameters inﬂuencing this
value are undoubtedly complicated, we can at least use it as an approximation for the modulus of the
SiNC inclusions in the composites. Our CLSM imaging conﬁrmed that a large fraction of SiNCs
are not individually dispersed within the PDMS matrix but instead are grouped together in clusters

59

up to several hundreds of nm in diameter and thus have a high amount of self-same interaction
rather than exclusively interacting with the matrix, and so considering their modulus within the
nanocomposite as being similar to the modulus of SiNC layers is a reasonable approximation for
the purposes of our analysis. Here, we can base our discussion around that value of µ as estimated
previously. Our ﬁrst approach to estimating the theoretical modulus of the composite employed
the standard theory of mixtures, in which the µ values for the host and inclusion are treated as
unchanging. The volume fraction of the mixture is used to weight the two µ values, as

µCOMP = µPDMS ⇤  PDMS + µNC ⇤  NC

(4.2)

where, µ(COMP),µ(PDMS) and µ(NC) are the moduli of the composite, PDMS and SiNC respec-
tively, and  (PDMS)and  (NC)are the volume fractions of PDMS and SiNC (calculated for each
sample type using ImageJ). Based on the previous equation, the composite is expected to exhibit a
modulus of 190.59 kPa, which is nearly identical to the modulus of PDMS (190 kPa) on its own.
From the theory standpoint, this is logical due to the very low volume fraction of SiNCs (calculated
theoretically at 0.1% for the high concentration samples). However, our measured value of µ in
both the functionalized and non-functionalized samples was lower than the modulus of PDMS on
its own. Even using the more sophisticated Mori-Tanaka approach [13, 120], the theoretical mixture
modulus was near the value for PDMS.

Recognizing this discrepancy, we also considered the possibility that the µ of the PDMS
matrix itself may be altered by the presence of the SiNCs at higher concentration. The clusters
of SiNCs are large enough to conceivably interfere with cross-linking between polymer chains, in
e ect countering the concentration of cross-linking agent and reducing the modulus of the PDMS
[22, 126].

Considering this, we can work backwards from our measured µ for the composites, assuming the
SiNC modulus as estimated in our prior work, and allow µPDMS to change. Using this approach,
we calculated that the PDMS modulus for the H-F samples was ⇠40% lower than the PDMS
without SiNCs, and for the H-NF samples it was even more dramatically altered, ⇠70% lower than
for PDMS alone. Based on established trends for the mechanical properties of PDMS as a function

60

of cross-linking agent concentration, this would mean that by including SiNCs (F and NF) in the
PDMS at a concentration of <1% causes an e ective change in the cross-linking ability of PDMS
equivalent to reducing the curing agent:prepolymer ratio to 1:12 and ⇠1:18, respectively, for the
H-F and H-NF samples [118].

4.4 Conclusions

We measured the luminescence and neo-Hookean coe cients of SiNC/PDMS nanocomposites
based on concentration and surface functionality. The PL QY from the nanocomposites is high at
16% for the samples made with surface-functionalized SiNCs, and still measurable at 1-2% for the
samples made with non-functionalized SiNCs at higher concentration. Our experiments indicate
that the mechanical behavior of luminescent SiNC/PDMS nanocomposites depends on both the
nanocrystal concentration within the PDMS as well as the surface functionality of the SiNCs. At
low concentration (0.03 wt%) the SiNC inclusions do not inﬂuence the rupture stretch nor the
neo-Hookean coe cient of the nanocomposites, regardless of surface functionality. However, at
higher concentration (0.1 wt%) there is a statistically signiﬁcant reduction in the neo-Hookean
coe cient. This is in contrast to the report by Chen et al. [23], who found similar PL results for
their SiNC/PDMS nanocomposites but reported a reduction in Young’s modulus (⇠three times µ at
low applied strain and assuming a Poisson’s ratio of 0.5) and the tensile strength of hydrogenated
(NF) samples but an increase in both for decyl-capped SiNCs in PDMS. We do not expect that the
di erence between decyl- and dodecyl-capped SiNCs should have such a profound impact on the
mechanical properties of the nanocomposites, neither at low nor at higher concentrations. Further,
the PL results from our study largely concur with their ﬁndings. Thus, these conﬂicting results
point to the importance of further characterization and an improved understanding of the exact
mechanisms of interaction between nanocrystals, their surface functionality, and manufacturing
procedures during fabrication of nanocomposites using elastomers such as PDMS.

61

CHAPTER 5

FUTURE

62

5.1 AuNP Application for Silicon NanoRods (SiNRs) Sythesis

Throughout our experiments on synthesis of AuNPs, our overarching objective was to fabricate
SiNRs with aspect ratio (AR) (Length/ Diameter) of 2-7. If successful, this work would introduce
a brand-new plasma-based ﬂow-through approach to continuous synthesis of SiNRs. Similar
anisotropic SiNR growth in a ﬂuid suspension was achieved by Heitsch et al, which used a ligand-
assisted solution-phase method, based on Au-seeded solution-liquid-solid (S-L-S) growth [63].
Our focus is on translating this approach from solution to the plasma, which can avoid solvent use
and could lead to wide tunability of the SiNR properties such as structure, AR, and surface in a
continuous nanomanufacturing method.

For plasma-based SiNR growth, we sought to build on existing methods and use seeded growth
based on AuNPs.
In Chapter 2, we reviewed di erent methods for fabrication of AuNPs and
introduced our plasma-based method which provides a new opportunity for fabricating SiNRs
based on multi-stage reactor design. Figure 5.1 shows a schematic diagram for the proposed multi-
stage reactor for producing AuNP-assisted SiNRs. The ultimate goal of this reactor is continuously
making gold nanoparticles and then using them as the seed catalysts for growing SiNRs and
continuously, using the plasma as energy source. For our preliminary trials, we constructed several
iterations of this reactor. The concept was based on multi-stage reactors, with the ﬁrst stage used
for AuNP synthesis, and a subsequent plasma stage for growth of SiNRs using a supersaturated
eutectic mixture method.

This reactor consists of gas cylinders connected to quarter inch tubing coming to MFCs and
tubes, non-thermal RF plasma, electrical circuit for matching the power and stainless steel tubing
for gas lines [107]. AuNPs are produced in the ﬁrst stage using the hot-wire or plasma-based
approach. After producing the AuNPs, in the second step these particles enter the second stage RF
plasma (operated at 13.56 MHz and with supplied powers ranging from 5 to 40 W), tuned with
a lab-constructed impedance matching networks including an adjustable capacitor, inductor coil,
and air gap. SiH4 was provided by a SiH4/Ar mixture cylinder (1% SiH4-99% Ar) as injected into
the reactor from a 0.5-inch O.D. sidearm at ﬂowrates between 5 and 50 sccm. The eutectic alloy

63

temperature for the Au/Si alloy is 363°C, and so we sought to explore the parameters necessary to
reach this temperature by adjusting reactor parameters such as power, gas ﬂowrate, overall pressure,
and reactant partial pressure.

Figure 5.1: (a) Schematic Diagram for SiNRs synthesis (b) early stage experimental set up with
hot wire method (Courtesy R.J.Anthony)

We modiﬁed the early stage design shown in Figure 5.1 by adding the Silane from an elevated
side arm between the two stages. The early design added silane in the middle of the second stage,
as presented in Figure 5.1, depicted in Figure 5.2. After using the reactor for ⇠ 20 mins, a thick
ﬁlm deposition of Silicon formed in the side arm and on the main reactor tube, preferentially on the
injection side. We hypothesized that e ective mixing of the AuNP-Ar ﬂow from the ﬁrst step and
Silane from the second stage did not take place, causing the SiH4 to remain spatially segregated,
decomposing into ﬁlm growth near the sidearm rather than mixing with the bulk ﬂow. Figure
5.3 shows the modiﬁed reactor, including the side arm for adding the Silane between two steps.
This modiﬁcation greatly reduced the gas mixture homogenization problem, and eliminated the
preferential ﬁlm growth near the side arm. We collected the reaction products at the bottom of
the second stage in the reactor, and performed TEM characterization to investigate possible SiNR

64

Figure 5.2: High ﬂowrate of AuNP-Ar from the ﬁrst step(600-1200sccm) and low ﬂowrate of SiH4
from the side arm did not mix properly in the second step of the reactor

Figure 5.3: Modiﬁed hotwire AuNP reactor for sythesis SiNRs

growth. Surprisingly the microscopy showed a widely size-dispersed set of silicon nanoparticles
with both crystalline and amorphous structures. We had expected to see some mix of AuNPs,
whether the SiNR growth had happened or not, as we were certainly injecting AuNPs into the

65

second stage. EDS provided some proof for AuNP from the ﬁrst step collected after the second
step of adding Silane without turning on the RF plasma in second stage. We focused on the form
of the plasma, which included a ﬁlamentary structure (constricted region of high plasma density)
during the second reaction. Other reports have indicated that, in argon/silane plasmas, the plasma
density and temperature spikes reached by the nanoparticles in the plasma are related [3, 84, 109].
We hypothesized that the constricted plasma ﬁlament caused su cient power density to in fact
coalesce the smaller SiNCs typically produced under the same pressure/ﬂow/power conditions
but with a di use plasma. While Si nanoparticles exhibit size-related melting point depression
[67, 108], the larger crystalline Si nanoparticles we observed had diameters >10 nm, indicating
that they reached temperatures in excess of 1300 K. Not only this is higher than the eutectic
alloy temperature we were seeking to reach, it also exceeds the melting temperature of gold. Our
hypothesis was that by passing through the ﬁlamentary region, the AuNPs were re-evaporated
(explaining their disappearance before collection on the TEM grid) and the SiNCs were coagulated
and recrystallized. The solution we proposed was to eliminate the ﬁlament using gas ﬂowrates.

We were able to stop the ﬁlament in the reactor by controlling the gas species ﬂowrate inputs
to the second stage (in particular, by keeping the SiH4 ﬂowrate <5 sccm) and by controlling the
power to the plasma (Rf power limited to the range of 5-15 Watts). Photographs of this second-
stage plasma with ﬁlament present (a) and the di use plasma (c) achieved using limited power
and reduced SiH4 ﬂowrate are presented in Figure 5.4. Unfortunately, though, we probed a wide
range of parameters within the conﬁnes of avoiding the ﬁlamentary plasma, we never observed
anisotropic growth from this reactor conﬁguration.

66

Figure 5.4: (a) Filament formation associated with high plasma density (b) High plasma density
resulted in wide size range of Silicon nanoparticles with both crystalline and amorphous structures.
(c) Controlling the SiH4 ﬂowrate and RF power stoped ﬁllament and resulted in narrow size
distribution for Silicon nanoparticles

As presented previously, the plasma-based method works with wider ranges of experimental
parameters, compared to the hotwire method. Thus, we next attempted to combine plasma-
synthesized AuNPs and SiNR growth. Figure 5.5 shows a schematic picture of the projected
dual-plasma reactor. Future work will involve exploring the full range of capabilities of this
dual-stage reactor.

67

Figure 5.5: Schematic of dual-plasma reactor for AuNPs-SINRs simultaneous synthesis

5.2 Modulus of SiNC Layers on PDMS: Parametric Investigation

5.2.1

Introduction

In-situ investigation of mechanical properties of nanomaterials in form of thin ﬁlms is one of the
challenging obstacles in advancement of nanotechnology and nanomaterial ﬁelds. There are many
studies dedicated to plastic properties of individual nanoparticles in crystalline and amorphous
forms [100]. Nanoindentation experiments and other physical probe-based techniques can be
used to evaluate single-nanocrystal deformation in elastic and plastic regime. These techniques
generally require especial microscopy techniques like tunneling electron microscopy. However, as
these approaches are mostly irreversible and/or destructive, there is a signiﬁcant interest in non-
destructive methods for evaluating properties of thin ﬁlms, especially for layers of nanocrystals
on elastomeric substrates. The future of electronics strongly relies on advancement in mechanical
ﬂexibility, sustainability and biocompatibility of electro-optical sensors, health monitoring systems,

68

photovoltaic devices, displays, and functional coatings. We recently introduced a novel in-situ
approach for estimating mechanical properties of luminescent SiNCs deposited on ﬂexible substrates
by measuring the onset of instabilities in the SiNC layers on PDMS under ﬁnite bending deformation
as reported in Chapter 3. This study revealed a neo-Hookean coe cient (µ, analogous to shear
modulus at low stress/strain) of a 4.5 µm thick SiNC ﬁlm to be 345 ± 23 kPa, which is signiﬁcantly
smaller than the young modulus of a single SiNC at the value of 170 GPa. There are many
parameters that could inﬂuence the mechanical behavior of the SiNC layers on PDMS, including
layer thickness, surface functionality, and mechanical properties of the PDMS substrate. In order to
reveal some of these interrelationships, we investigated these parameters regarding their inﬂuence
over the mechanical properties of SiNC layers on PDMS. Our results indicate that the layer of
SiNCs and PDMS parameters exert competing inﬂuences on the mechanical properties of the SiNC
layers. For example, we found that reducing the sti ness (reducing µ) for PDMS leads to a decrease
in µ for the SiNC layer. However, the relationship is not monotonic. Unraveling the connection
of these parameters on the physical, mechanical, and optical properties of the SiNCs will allow
advancement for predictive engineering of nanomaterial thin ﬁlms, for stretchable and ﬂexible
devices.

5.2.2 Experimental work

We prepared the PDMS in lab using a Sylgard prepolymer PDMS 184-kit and curing agent catalyst
produced by Dow Corning. The PDMS mechanical properties depend on the prepolymer PDMS
kit-184/curing agent catalyst ratio and curing conditions during PDMS preparation. We produced
a range of substrate with same geometrical shape and dimensions (deﬁned earlier in Chapter 3
as slenderness ratio). For experiments designed around modifying the SiNC layer, the prepoly-
mer:curing agent ratio was 10:1, as before. For experiments to test the dependence of SiNC layer
modulus on PDMS modulus, we altered this ratio to produce PDMS slabs with lower moduli. With
a higher fraction of curing agent, the modulus of the resulting cured PDMS increases. [78]. We
attempted to increase the modulus for this study by using a 7:1 ratio of prepolymer to curing agent,

69

but our experiments revealed that the critical angle required for bifurcation onset was too large to
measure practically. Instead we reduced the amount of curing agent, using ratios of 10:1, 12:1,
and 20:1 for this parametric study. We used the same method for evaluating µ for the substrates as
discussed earlier in Chapter 3 and Chapter 4. For performing this parametric study, we chose a 3/1
slenderness ratio for our substrate.

As we mentioned earlier there are many advantages for gas phase processing. Similar to
Chapter 3 and Chapter 4, we synthesized SiNCs using a low-pressure non-thermal plasma. The
deposition of nanocrystals occurs directly out of the described reactor into thin-ﬁlm layers using
inertial impaction through a slit-shaped oriﬁce onto a substrate placed beneath[68, 124]. This
reactor enables us to control thickness of the ﬁlm, pressure, sample-oriﬁce distance and forwarded
RF power to the glass tube.

As we discussed in Chapter 3 there is a stand o  distance between oriﬁce outlet and the top
surface of PDMS substrate which is around 4-5 mm. This distance plays a signiﬁcant role on the
thickness and porosity of the thin ﬁlm during deposition via back and forth rastering motion of
the substrate. For instance, decreasing the distance would result in a thin ﬁlm with lower porosity
(higher density) and increasing the distance on the other hand, would result in higher porosity
(lower density) [5, 68]. This is a consequence of the dynamics of SiNCs exiting the reactor from
oriﬁce outlet.

Since we ﬁxed the value for stand o  distance and the frequency of rastering motion for substrate
(3 sec for each back and forth motion), the only parameter which impacts the thickness of the thin
ﬁlm is the number of rastering cyclces. In other words, changing the time of rastering from 20
mins as discussed in Chapter 3 to 10 mins and 5 mins results in a thinner and thinnest ﬁlm in our
results, respectively. We also produced samples with higher rastering time, 30 mins, as well. We
used SEM analysis to evaluate the thickness of the layers depending on rastering time.

As we discussed the surface functionalization of the SiNCs will improve their PL properties,
but changing the surfaces may also inﬂuence the mechanical behavior of the layer as a whole. We
used a facile method to surface functionalize the SiNCs in the gas phase during the plasma process

70

as previously studied by Anthony et al. [4]. Here, we used 1-dodecene for surface functionalization.
We also measured the PL and QY for the produced ﬁlm and results presented in Fig 5.10. We
also used SEM to compare the surface morphology of non-surface functionalized and surface
functionalized thin ﬁlms under similar conditions.

In a summary of our experiments, Figure 5.6 presents a work ﬂow of this parametric study.

Figure 5.6: Schematic diagram for investigating the impact of ﬁlm thickness, surface functionality
of SiNCs via gas phase plasma and shear modulus of the PDMS substrate on mechanical properties
of SiNC thin ﬁlm

5.2.3 Results

Film Thickness: We used a mechanical diamond glass cutter pen to mark our samples for thickness
measurements. Figure 5.7 depicts the cross sectioning of samples produced via non-thermal plasma
with various rastering deposition times. Regarding the the SEM image, which consisted from
samples with 5 mins , 10 mins and 30 mins rastering time, the thickness of the ﬁlm increased
linearly with respect to the time of rastering deposition in plasma reactor. We performed a linear

71

regression and the R2 value was 0.9918. We also attempted a exponential ﬁt, for which the R2 value
was 0.9602. We expect that the deposition rate is linear, but it is possible that at low thicknesses
there is a deviation from this behavior. Future work will investigate the sub-micron ﬁlm thicknesses
to evaluate if the rate remains constant even very close to the substrate, where packing e ects may
inﬂuence the porosity of the SiNC layer. In Figure 5.7 sample in red represented the results from
Chapter 3 for comparison.

Figure 5.7: SEM images for PDMS with 3:1 slenderness ratio as deﬁned in chapter 3 with rastering
deposition time of 5 mins (Sample 1), 10 mins (Sample 2) and 30 mins (Sample 3) and averaged
measured thickness of the ﬁlm with corresponding standard deviation in comparison with the
sample discussed with same slenderness ratio and 20 mins (in red)of rastering deposition time from
previous chapter

Modulus of SiNC Layers (Estimated Values): We used the same methodology employing
the onset of bifurcation to estimate the neo-Hookean coe cient for the SiNC layers subject to the
parametric variation as noted above. The results are plotted in Figure 5.8. The asterisks on the bars
indicate samples which had a statistically signiﬁcant deviation from the value we obtained for the
samples originally examined and reported in Chapter 3.

72

Figure 5.8: The methodology introduced in Chapter 3, employed to estimate neo-Hookean coef-
ﬁcient for SiNC layars corresponds to parametric variation. The asterisks indicate samples with
signiﬁcant deviation from value we abrtained in Chapter 3

The ﬁrst observation is that the thickness of the SiNC layers does inﬂuence the neo-Hookean
coe cient, with thinner layers exhibiting increased µ as compared to the thicker layers. However,
increasing the thickness beyond the 4-5 µm, which was our standard, does not appear to alter the µ
of the SiNC layers. From this, our hypothesis is that the soft PDMS substrate creates microstructural
changes or densiﬁcations of the SiNCs, altering their mechanical behavior close to the substrate -
yet once the layer reaches a certain critical thickness, the e ect diminishes as the subsequent layers
are deposited and they are less able to respond to the substrate’s mechanical properties during
deposition. Figure 5.9 presents the calculated mechanical property (neo-Hookean coe cient )of
the SiNC thin ﬁlm deposited on PDMS. The asterisks show the samples with neo-Hookian value
signiﬁcantly deviate from samples reported in Chapter 3.

Next, we can observe that the modulus of the PDMS also changes the resulting SiNC modulus,

73

Figure 5.9: Neo-Hookean coe cient of all sample introduced in this section.The samples with
asterisks, present the data with siginﬁcant deviation from samples reported in Chapter 3

with decreasing PDMS modulus creating a decreased µ for the SiNCs. Our hypothesis to this
regard is still very nascent and this research is ongoing, but we can draw some ideas from what we
observed during the thickness measurements. The thinnest layers of SiNCs on substrates made with
the standard PDMS recipe (10:1 prepolymer:curing agent) have a higher modulus (are sti er) as
compared to the thicker layers. From this it is reasonable to expect that near the substrate there are
e ects from the impact of the NCs onto the PDMS, but as the layer becomes thicker these substrate
e ects are diminished. When we decrease the modulus of the PDMS by altering the recipe, we
possibly alter the damping e ect of the SiNCs impacting upon the substrate. We know that the
thickness of the SiNC layer does not vary with changing PDMS modulus, but it’s possible that the
propagation of the substrate e ect is reduced for the substrates with lower modulus. Our ongoing
experiments will check this possibility by studying very thin layers of SiNCs (rastered for  5 mins)
to test the limits of the substrate e ect in terms of how far into the SiNC ﬁlm is plays a role in the
mechanical behavior.

74

Notable is that the surface functionalization plays no statistically signiﬁcant role in the modulus
of the SiNC layers (see Figure5.9). This is especially important regarding the fact that the SiNCs
with in-ﬂight surface functionalization display much higher PL quantum yields as compared to
those made without surface functionalization (See Figure. 5.10). We can conclude that, for
brightly emissive layers of SiNCs on PDMS, one can perform surface functionalization without
disrupting the layer’s mechanical behavior.

Figure 5.10: The surface functionalization applied in gas phase during plasma processing and
sythesis of SiNCs. The PL and QY measured. This results coupled with mechanical property
estimation of corresponding samples presented in this Figure. The results here, conﬁrm that
improving the oprical properties and QY do not interfere with mechanical property of the SiNC
thin ﬁlm.

5.2.4 Conclusions

This study is nearly complete. In summary, we have investigated some of the parameters of SiNC
layer properties and PDMS properties. We discovered that reducing the thickness of the SiNC
layer increases the modulus of the layer, and that in-ﬂight surface modiﬁcation greatly improves

75

the PLQY but does not appear to signiﬁcantly alter the modulus of the SiNCs. However, reducing
the PDMS modulus by altering the recipe for fabrication causes a signiﬁcant reduction in the SiNC
layer modulus. Our ongoing experiments and modeling will seek to explain these parametric e ects
and recommend future experiments.

76

5.3 Unbending- Another Finite Bending Condition for Mechanical Property

Estimation

5.3.1

Introduction

Prior experiments involved the deformation of SiNC layers deposited on PDMS in a relaxed
condition, and the bifurcations we observed upon bending the system as a whole. Here, instead we
focus on experimental observations and physical circumstances which result in surface instabilities
(wrinkling) of SiNCs layers deposited onto PDMS substrates in a prestretched condition, with
future applications in tunable gratings and new manifestations of LEDs and other devices. Prior
work has shown the emergence of wrinkling patterns in these systems [104], but here we have
performed experiments to understand the conditions of the SiNC layers and the PDMS under which
these wrinkling patterns are formed.

5.3.1.1 SiNC deposition

In this section, we discuss the various deposition schemes for placing the SiNCs layers onto the
PDMS. The goal is to deposit a uniform layer of SiNC onto the PDMS as thin ﬁlm, under varying
pre-stretched conditions. We are interested in studying how the optical and physical properties of
the SiNCs are altered when various degrees of pre-strain are applied to the PDMS.

We prepared a stage (Figure 5.11) which can be used to obtain any stretching ratio between 0-
100%, and which is compatible with holes in both ends to be mounted onto a pushrod in the plasma
reactor. This apparatus helps us to use PDMS as a substrate since we apply coaxial stress-strain
to substrate. One end is ﬁxed, and the other end is moving with a corresponding twisting screw.
We achieve uniform deposition by rastering the substrate as mounted on the stage underneath the
oriﬁce. The rastering is a back and forward movement which helps to keep the thickness uniform
throughout the substrate. While this allows for easy control over the applied pre-strain in the long
direction and uniform deposition, it also leads to a pronounced bi-axial deformation due to the
Poisson’s e ect of the PDMS, meaning that during the applied linear pre-strain there is also a

77

compression across the perpendicular edge of the PDMS. This compression is measured to be ⇠
10 percent at the waist of compression in the middle section (Figure 5.12).

Figure 5.11: Sample holder for various stretching ratios (left), sample holder for ﬁnite bending
around pushing rod (right)

Figure 5.12: Above view of sample, relaxed after deposition

5.3.2 Results

As we discussed in previous section about loading the substrate in the reactor, here, we present the
results for both methods. In the ﬁrst method, the PDMS sample was loaded on the sample holder
and pre-stretched between 5% to 40% beyond the initial length (relaxed). After the SiNCs were
deposited uniformly onto the PDMS in the reactor via inertial impaction, we used the screw on
the stage to return the substrate to its original length steadily. During this process, SiNCs (which
were experiencing no tensile force or compression on the pre-stretched PDMS) start to experience
compression and this makes surface instabilities on SiNC layers. These wrinkles are observed by
mounting the stage-sample setup in a vacuum chamber of SEM. The corresponding wrinkles for
20% applied pre-strain are depicted in Figure 5.13. Their axis can be categorized based on the area
in which they formed, depending on the tensile or compressive forces that they experience during
relaxation of the PDMS. Again, we see this combination of forces due to the Poisson’s ratio of

78

the PDMS, leading to the SiNC layer experiencing compression near the clamps and tension near
the waist of the pre-stretched PDMS. We categorized these responses to pre-stretched forces into
3 regimes based on position on the pre-stretched substrate: pure compression, pure tension, and a
combination. In order to study the whole deposited area, we divided each deposited area of sample
to 15 cells, including 3 rows (A-B-C) and 5 columns (1to5). The corresponding deposited area
with cells shown in Figure 5.14 and Figure 5.12. Since the PDMS applied pre-strain was 20%, the

Figure 5.13: SEM image of deposited area with 20 percent pre-stretching condition, Column A3-
B3-C3 (left), Columns A2-B2-C2 and A4-B4-C4 (middle), Columns A1-B1-C1 and A5-B5-C5
(right)

Figure 5.14: SEM image of deposited area with 20 percent pre-stretching condition, Column A3-
B3-C3 (left), Columns A2-B2-C2 and A4-B4-C4 (middle), Columns A1-B1-C1 and A5-B5-C5
(right)

shape of the substrate during the SiNC deposition is not rectangular. Moreover, this new geometry
will result in an interesting force response that include wrinkles on the surface of the SiNCs. The
SEM images from di erent regions are shown in Figure 5.13. We can categorize the corresponding
surface instabilities into 3 distinctive types: each occurs in speciﬁc cell areas depicted in Figure
5.14. As we can see, in the area near clamps of both side, where the deposited area changes

79

its thickness to some su cient amount we observe vertical Lines. These areas include columns
A1-B1-C1 and A5-B5-C5. However, in the column A2-B2-C2 and A4-B4-C4 we observe transient
or inclined directions with respect to the X-Y Cartesian coordinate introduced in Figure 5.12. In
the center, including middle column (A3-B3-C3) we observe horizontal wrinkles. Since these
mechanical instabilities are di erent in terms of direction they appear throughout the deposited
area, we justify the observation from mechanical point of view.

Figure 5.15 shows a diagram of the forces on the PDMS during stretching and corresponding
tension/compression experienced by the PDMS during stretching. In order to understand the force
distribution (tensile vs. compressive), we represented the PDMS in a top view schematic before
deposition (in pre-stretch condition) and after deposition (in relaxed condition). The stage is set to
move in one direction and it designed to stretch the PDMS from 100% (Zero sterteched) to 200%
(double stretched), depending on the experiment design and initial PDMS dimensions. The Poisson
Ratio for PDMS is ⇠ 0.5 [138]. This means that the strain in longitudinal direction is twice the
strain in compression along the middle section of the sample. So, the PDMS area near the clamps

Figure 5.15: pre-stretched before deposition or during deposition

experience tension, however, one clamp is ﬁxed in position. During deposition SiNCs these areas
are deposited on a stretched surface but experience zero force while the SiNC layers are formed.
After deposition, when the PDMS is relaxed back to its original length, SiNCs will experience
compression near the edges (columns 1 and 5 in our areal mapping scheme) and they start to form
surface instabilities as shown in the Figure 5.13 (right). However, the middle section of PDMS
(column 3) is under compression during the deposition. By returning to relaxed condition, the
SiNC layer in the middle section would experience tension. Ultimately, in this middle section the

80

SiNC layers form instabilities perpendicular to the ones near edges. Between the edge regions and
the middle region (columns 2 and 4), the SiNCs experience a hybrid of these responses.

In order to limit the scope of this study, we focused on to the formation of instabilities near
edges, which result from the SiNC layers experiencing pure compression after the PDMS is relaxed.
As a part of manufacturing design, we had to provide a design to able reproduce this (near edges)
kind of surface instabilities which in order to categorize them in a similar formation direction since
we apply one-dimension force to the substrate. Due to the nature of the PDMS, it is impossible
to eliminate the tensile force experienced by the SiNCs when the PDMS is pre-stretched using the
stretching stage we described hre so far. Thus, we have to provide a new experimental set up for
loading the PDMS: our solution is to use ﬁnite bending of the PDMS. This approach will mostly
eliminate the Poisson’s ratio e ects on the region in the middle of the PDMS. Since the PDMS
wraps around the pushrod, the upper surface (SiNC-coated surface) experiences a strain in the
circumferential direction and there would be no dimensional change corresponding Poisson’s e ect
for substrate.

We explored ﬁnite-bending as a way to minimize the bi-axial e ect as observed for the stretching
stage described above. To achieve this, we cast the PDMS into aluminum molds of calculated size,
attached handles to the PDMS during the process with a commercial painting foam (Foam Board),
and wrapped it around the pushrod, gripping the handles together at the base using a staple. Then,
SiNCs are deposited onto the PDMS (which is deformed via this ﬁnite bending), and when the
handles are released, the PDMS returns to its relaxed state. Figure 5.11 shows both setups. Figure
5.16 shows a picture of a prospective sample going under ﬁnite bending elastic deformation. We
wrapped PDMS around the pushing rod and deposited SiNCs on the wrapped substrate. In order
to guarantee an evenly deposited ﬁlm on substrate we have to clamp the PDMS (Figure 5.16).
The microscope picture shown in right side of the Figure 5.16 , is captured by Optical-Electrical
Microscope (Keyennce) under aforementioned conditions. This picture conﬁrms the formation of
wrinkles with the new design for substrate. However, since we applied some force to clamp the
PDMS all around the push-rod, there would be extra force we applied for wrapping around the rod.

81

As the next step, we had to ﬁnd exact substrate dimensions that allow the substrate to wrap around
the rod without the need of additional clamping or stretching.

Figure 5.16: Early design for ﬁnite bending surface instabilities by wrapping PDMS around the
rod, the picture on right is the microscope image of the sample after unwrapped with the scale bar
of 10 micrometer

To measure the exact size, we need the sample to be wrapped around the pushrod with no tension.
We used a mathematical calculation to ﬁnd the width to thickness (w/t) ratio for a corresponding
pre-strain (here it is 20%). So then with that ratio, we are able to wrap the substrate around the
pushrod with no extra force which makes it feasible to have surface instabilities based on pure
ﬁnite bending.
(Figure 5.17). Based on the Constitutive Laws and Neo-Hookean assumptions
given in Figure 5.17, we are able to apply the corresponding w/t ratio to make an aluminum cast
for fabricating PDMS with handles. Figure 5.17 shows the corresponding result for this set up,
beside a picture of the setup, including the aluminum cast and ﬂexible handles for sample to be
wrapped around the push-rod and securely ﬁxed in place during deposition. This helps us to make
sure that during the deposition the rotation of substrate with the pushrod for uniform deposition
will not change the position and distance of substrate from the oriﬁce which is an important factor
for formation of uniform and continuous SiNCs ﬁlm on the substrate. The thickness and width of
the PDMS sample, in the reference conﬁguration, is taken in such a way that there are no tractions
on the surface when it is rolled on a cylindrical rod of a given outer radius. When the system
undergoes the unbending deformation the thin SiNC layer deposited on the outside surface of the
PDMS experiences a compressive force. This is similar to the compressive force faced by the top
surface of the SiNC in the method involving ﬁnite bending. Therefore, like in the ﬁnite bending
experiments, we expect the formation of instabilities on the surface of thin ﬁlm above a critical

82

compressive stress. This value of compressive stress depends upon the geometric conﬁguration and
mechanical properties of the system. We now able to produce surface instabilities based on pure

Figure 5.17: A new method for pure bending surface instabilities, the mathematical correlation and
assumptions for incompressible, Neo-Hookean material (left)

bending, and the microscope image also conﬁrms this result. Figure 5.18 shows a set of wrinkles
produced by pure bending e ect. These wrinkles spanned over di erent areas in the sample
including rows A1-A2-A3-A4-A5 and C1-C2-C3-C4-C5. The surface instabilities will be studied

Figure 5.18: A new method for pure bending surface instabilities, the mathematical correlation and
assumptions for incompressible, Neo-Hookean material (left)

by their amplitude and the size. Since the mechanical properties leading to instability formation
are dependent on the bilayer system of PDMS and SiNCs, this study provides an opportunity
to correlate the luminescence and mechanical properties of the SiNC layer based on the surface
wrinkles. Our early calculation is based on images from an optical microscope which helps us
measure the size of the wrinkles based on where the wrinkles are formed. Table 5.1 provides a list

83

of operation conditions used for making surface instabilities in SiNC plasma reactor. The time of

Table 5.1: List of operation used for making surface instabilities in SiNC plasma reactor at PNL

deposition can turn into number of rotations of sample around the rod underneath the oriﬁce. To
calculate this number, we measured each rotation time under the oriﬁce in the plasma deposition
reactor which is ⇠ 3 seconds. This will help us to compare the thickness of the ﬁlm quantitatively.
The forward power is the power sent out to the plasma reactor and can be optimized with matching
box. It is the operator’s duty to check the appearance (shape, intensity and length) of the plasma
for set of recorded parameters. The power range stated in the table is used to form SiNCs with high
crystallinity. The crystallinity can be checked by XRD. To verify the crystallinity, we performed
a brief power study using the same ﬂowrates and pressure, however, the power received was 30
and 25 respectively. Figure 5.19 shows the XRD patterns from these samples, indicating a higher
degree of crystallinity for 30 Watt power than 25 Watt. The clear peaks from the [111], [220] and
[311] crystallographic planes at 28.460°, 47.510°and 55.910°observed respectively [130]. This can
be further veriﬁed by Raman Spectral Analysis [74, 107, 128]. This has been proven with other
studies made for SiNCs [3].

84

Figure 5.19: XRD for SiNCs deposited on PDMS, 30 Watt(Left), 25 Watt (right)

Figure 5.20 shows a peak-to-peak distance measurement for these surface instabilities as based
on the optical microscopy. The microscope image shown in Figure 30 is one of the corresponding
samples used for this measurement. Lamda ( ) is the length of peak to peak separation of two
wrinkles next to each other. This   changes from⇠ 1 micron to more than 20 microns, corresponding
to the di erent cell areas we deﬁned for each PDMS sample (see Fig. 22). The aforementioned
system includes 15 equal size cells and these wrinkles repeated with di erent size in various cells.
Moreover, since each rotation under oriﬁce takes ⇠ 3 seconds, then 10 rounds equal to 30 seconds
(⇠ 100nm), 20 rounds equal to 60 seconds (⇠ 200nm) and 30 rounds equal to 90 seconds (⇠
300nm). These thickness calculations are based on experimental image samples captured by SEM.
Increasing the time or number of rotation (while we keep same speed for rotation) will result in
thicker ﬁlms on the substrates. Regarding the previous SEM thickness measurements Figure 5.20
includes average size with corresponding standard deviation. Since we have 3 samples with varying
thicknesses we can measure the   of the wrinkles for each sample. All calculations are based on
optical microscope images.

85

Figure 5.20: Lambda ( ) measurement for surface instabilities, 10 round rotations in the reactor(up-
left), 20 round rotations(up-right) and 30 round rotations (bottom), numbers are corresponding to
the cells that measurement done with speciﬁed standard deviation for each averaged size letter
corresponding to the cell which wrinkle instabilities observed

Figure 5.21: A sample of PDMS deposited with SiNCs under pre-stretched condition for 10 round
rotation(left), deﬁnition of Lambda ( ) and thickness of the SiNCs ﬁlm(right)

The   measured based on the schematic diagram shown in Figure 5.21. As demonstrated in
Figure 5.20,   is not changing much for most of the samples, however, near the edges, including(A1-
A5 and C1-C5) the standard deviation is bigger, which means the size of   changing from few
microns to 10 microns.

To able to track the surface instability regimes and their amplitude changes on the surface we
have to observe the cross-section image of formed wrinkle patterns by microscopy. Cross-sectional

86

imaging for these soft substrates has been an issue for many instruments and techniques including
surface and thickness measurement by spectroscopic ellipsometry and Focused Ion Beam (FIB)
cross-sectioning with SEM. Atomic Force microscopy (AFM), however, is a physical and optical
method which provides new information about surface morphology in micrometer scale. This
method has its own limitation including the maximum scanned size per each cycle. This method
has been used for the sample shown in the result section for 20 rounds and the image shown in
Figure 5.22. In this ﬁgure, a 3D model is constructed based on surface topology data from the AFM
tip over a 20x20 micrometer area. This 3D model is conjugated with a cross section estimation to
help us evaluate the amplitude of these surface instabilities.

In future experiments, we will study the inﬂuence of SiNC size and the regime of wrinkles.
Ultimately, we intend to study multilayer ﬁlm composed from SiNCs with di erent sizes to study
varied emission wavelengths.

Figure 5.22: Atomic Forced Microscopy (AFM), AFM scanned a line across a 20 by 20 micrometer
area and provided a 3D model and a cross-sectional view for investigating the amplitude of these
surface instabilities

5.3.3 Conclusion

In this section, we introduced another class of ﬁnite bending problem. Surface instabilities formed
on thin ﬁlm of SiNCs deposited on PDMS and their corresponding patterns categorized based
on tenssion and compression deformation of the substrate. In order to control the formation of
surface instabilities, we intorduced ﬁnite bending (unbending) for prestretched substrate (PDMS)
undergoing plasma deposition process. In this conﬁguration, ﬁnal bilayer system unwrapped and
compressed the thin ﬁlm of SiNCs on top. Surface instabilities formed and their size and presence

87

redorded for di erent thicknesses of the thin ﬁlm. Our study showed these surface instabilities
change to cracks and fractures on the thin ﬁlm by increasing thickness of the thin ﬁlm. The size of
these instabilities will be correlated to the mechanical property ratio of the thin ﬁlm to the one of the
substrate. These results will be coupled by a theoretical moding to estimate mehanical properties
of the SiNCs deposited on PDMS as thin ﬁlm.

5.3.4 Future work

A theoretical model will be applied to the system introduced in this section for mechanical property
estimation. This evaluation will be compared with results presented in Chapter 3 and parametric
study introduced earlier in this chapter. This e eort will provide a new technique for estimating
the mechanical property measurement of the thin ﬁlms of SiNCs in an in-situ and non destructive
fashion.

88

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