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THREE-DIMENSIONAL PERIODIC STRUCTURES FOR ENHANCING LIGHT-MATTER INTERACTION AND ENERGY STORAGE
BY
HAILONG NING
DISSERTATION
Submitted in partial fulfillment of the requirements for the degree of Doctor of Philosophy in Materials Science and Engineering
in the Graduate College of the University of Illinois at Urbana-Champaign, 2014
Urbana, Illinois Doctoral Committee: Professor Paul V. Braun, Chair Professor John A. Rogers Professor Lane W. Martin Professor Kent D. Choquette
ii
ABSTRACT
Three-dimensional (3D) periodic architectures hold great promise for applications
ranging from manipulating the flow of light for integrated photonics to high power and
high energy batteries. Among the approaches to fabricate 3D meso-structured materials,
colloidal self-assembly and holographic lithography are particularly attractive owing to
their ability to create large, uniform templates. However, these 3D structures require
extrinsic functionalities (e.g. emitters, microcavities or energy materials) to fully utilize
their potentials. This thesis focused on additions of functional defects to the 3D networks
and studied the enhanced interactions between the embedded defects and the 3D host
materials.
A method based on epitaxial colloidal opal growth was developed to place
fluorescent nanoparticles at specific locations inside 3D silicon inverse opal photonic
crystals (PhCs), allowing the coupling between high dielectric contrast PhCs and
localized emitters to be investigated. Transfer-printing was next used to assemble a new
type of 3D PhC vertical microcavity consisting of a planar defect sandwiched between
two silicon inverse opals. This technique was similarly applied to embed pre-defined
high-quality defects into 3D holographic PhCs. Objects such as nanoparticle films,
spheres, and emitters served as defects and were introduced to well-defined positions.
Finally, interdigitated microbatteries were created from templates defined by both
3D holographic lithography and conventional UV lithography. The influence of electrode
width on liquid-phase ion diffusion was studied, which provided design parameters of
microbatteries for practical applications.
iii
ACKNOWLEDGMENTS
I have always felt extremely fortunate to have Prof. Paul Braun as my advisor. He
has taught me not only enormous amount of knowledge but also how to approach
research and convey scientific ideas at a top level. I am especially grateful for his
constant support and willingness that have allowed me to explore many ideas and to grow
quickly as a researcher in the past four years. I must admit that his mentoring and
research styles have had a great impact on me, from which I will continue to benefit in
my future career. I would also like to thank my committee, Prof. Rogers, Prof. Martin,
and Prof. Choquette. I truly appreciate their contributions and feedbacks to this work.
One of my best experiences at Illinois is that I have had the opportunity to work
with a few awesome colleagues and friends. The first person I would like to thank is
Kevin Arpin. He was not only a great colleague who was always ready to help and share
his working knowledge but also a supportive and honest friend. I have really enjoyed the
adventures (and exposures to the wild aspect of American culture) he provided in and out
of grad school. I was very fortunate to work (out) closely with Neil Krueger in the lab
(gym). Our common goal – to be either intelligent or athletic (or both) by the end of grad
school – has motivated us to be productive in both places. I immensely value our
collaboration and friendship. I would also like to thank Runyu Zhang for his ceaseless
help and support to my research and the fun we had together on the basketball court.
Dr. Masao Miyake and Dr. Agustin Mihi both deserve special thanks. I was very
fortunate to begin my research at Illinois with learning from Masao. He taught me
everything I know about the fabrication and simulation of 3D holographic photonic
iv
crystals, rare-earth emitters and also those important skills and techniques that I have
constantly benefited from throughout my entire PhD study. Agustin is one of my favorite
colleagues and friends. He was such a fun person to work with and also full of amazing
research ideas. I also own a special thanks to Dr. Joe Geddes for all the help and
insightful discussions he provided. I greatly admire his constant willingness to help and
sharp skills to solve problems.
I would like to thank Anthony Keum, Prof. Seok Kim, and Xing Sheng from
Rogers group. Without their help, I would not be able to finish this work at such pace.
My exposure and entrance to the advanced transfer-printing started with collaborating
with Anthony and Prof. Kim. Interactions with them have been a pleasant, simulating and
efficient learning process for me. Xing was one of the best collaborators I have worked
with. His knowledge on photonic and semiconductor devices has played an important
role in our successful achievement of hybrid III-V-porous silicon microcavities.
Acknowledgements are also due to Steven Zhang for guiding me into the field of
batteries and his informative tutorials on the principles of batteries whenever asked and
James Pikul for sharing his working experience on microbatteries with me. I would like
to specially thank my labmates Junjie Wang, Jiung Cho and Matt Goodman. Their help
and endless care to the lab has made my everyday research life so smooth and easy. A
number of facility staff Scott Robin, Dianwen Zhang, Tao Shang and Julio Soars deserve
special acknowledgements for offering equipment training and aiding the completion of
this work. I was also fortunate enough to reunion with my best friend since childhood
Qiyan Wang and also meet a number of new friends during grad school, Mark Losego,
v
Jian Yang, Simon Dunham, Henghua Jin and Andy Cloud, which has kept my non-
research life very interesting and enjoyable.
I must thank my parents. None of this could have been possible without their love,
support and sacrifice. My parents have given everything they can to provide the best
environment for me to grow up. While I cannot often be with them when they miss or
need me on the other side of Pacific Ocean, making them proud becomes my top priority
that has driven me to work hard and overcome all kinds of difficulties. Last but not least,
I would like to thank my wife. She gave up her whole career back home and followed me
to chase the American dream. I cannot express how grateful I am to have her and how
much I look forward to our future journey.
vi
TABLE OF CONTENTS
LIST OF ABBREVIATIONS viii CHAPTER ONE – INTRODUCTION TO 3D PHOTONIC CRYSTALS 1
1.1 Theory of photonic crystals 1 1.2 Experimental realization of 3D photonic crystals 6 1.3 Controlling spontaneous emission in 3D photonic crystals 17 1.4 3D photonic crystals for energy storage applications 25 1.5 References 27
CHAPTER TWO – CONTROL OF SPONTANEOUS EMISSION IN 3D SILICON PHOTONIC CRYSTALS 30 2.1 Introduction & motivation 30 2.2 Fabrication of silicon photonic crystal sandwich structures 31 2.3 Optical characterizations 34 2.4 Photonic density of states correlations 40 2.5 Conclusions 46
2.6 References 47 CHAPTER THREE – 3D SILICON PHOTONIC CRYSTAL MICROCAVITY 49
3.1 Introduction & motivation 49 3.2 Design of silicon inverse opal microcavities 50 3.3 Fabrication of silicon photonic crystal microcavities 52 3.4 Conclusions 60 3.5 References 61 CHAPTER FOUR – INCORPORATION OF FUNCTIONAL DEFECTS INTO 3D HOLOGRAPHIC PHOTONIC CRYSTALS 62 4.1 Introduction & motivation 62
4.2 Experimental procedures for embedding defects 63 4.3 Light-matter interaction between introduced defects & their hosts 67 4.4 Conclusions 72 4.5 References 73 CHAPTER FIVE – ASSEMBLY OF TUNABLE POROUS SILICON MICROCAVITY 75 5.1 Introduction & motivation 75
5.2 Printing hybrid porous silicon microcavity 77 5.3 Coarse tuning of microcavity resonance 81
vii
5.4 Fine tuning of microcavity resonance 83 5.5 Incorporation of solid state thin film emitters 84 5.6 Conclusions 87
5.7 References 87 CHAPTER SIX – HIGH POWER LITHIUM ION MICROBATTERY FROM 3D HOLOGRAPHIC LITHOGRAPHY 89 6.1 Introduction & motivation 89 6.2 Microbattery assembly 90 6.3 Electrochemical testing of microbatteries 94 6.4 Simulation and optimization of microbattery in COMSOL 103 6.5 Conclusions 114 6.6 References 115 CHAPTER SEVEN – CONCLUSIONS AND FUTURE WORK 116 7.1 Conclusions 116 7.2 Future work 118
viii
LIST OF ABBREVIATIONS
1D one-dimensional
2D two-dimensional
3D three-dimensional
ALD atomic layer deposition
Cu2O copper oxide
CVD chemical vapor deposition
DBR distributed Bragg reflector
DLW direct laser writing
DOS density of states
FCC face centered cubic
FDTD finite difference time domain
FWHM full width half maximum
HF hydrofluoric acid
ITO indium tin oxide
Li lithium
MnO2 manganese oxide
Ni nickel
PAG photo acid generator
PAG photo acid generators
PBG photonic band gap
PDMS Polydimethylsiloxane
pDOS photonic density of states
ix
PhC photonic crystal
PR photoresist
PSi porous silicon
QD quantum dots
Q-factor quality factor
RIE reactive ion etching
SE spontaneous emission
SEM scanning electron microscopy
Si silicon
SiO2 silicon dioxide
SOC state of charge
TPP two photon polymerization
1
CHAPTER ONE
INTRODUCTION TO 3D PHOTONIC CRYSTALS
1.1 Theory of photonic crystals
1.1.1 Origin of photonic band gap
Photonic crystals (PhCs) are materials with a periodically varying refractive index
on a length scale comparable to the wavelength of light. In such structures, photons with
a specific range of energies (the so called stop gap) cannot propagate along certain
directions, resembling the case that no electronic states are allowed inside the bandgap of
a semiconductor. The concept of PhCs was independently first advanced by Yablonovitch
and John in their efforts to control light for different goals, where Yalonovitch proposed
to suppress spontaneous emission using PhCs,[1] while John found strong localization of
light in those structures.[2] The physics that governs the stop gap for photons has its origin
in the coherence of scattered light from the periodic dielectric modulation. For a given
wavelength inside the band gap as shown in Figure 1.1a, the scattered light from each
layer is in phase with each other, producing a standing wave with the incident light that
does not travel in the PhC. When the wavelength is outside the band gap, the partial
scattered waves are out of phase and thus cancel each other (Figure 1.1b), allowing the
incident light to propagate through the structure.
A periodic multilayer film, often called a 1D PhC or distributed Bragg reflector
(DBR), is the simplest PhC.[3] An important application for 1D PhCs is as dielectric
mirrors for semiconductor lasers.[4-8] 1D PhCs possess a highly angle-dependent stop gap
and thus can only control light in limited directions. 2D PhCs having periodicity in two
dimen
total
refrac
light
frequ
light
photo
Figurwave
nsions can e
internal refl
ctive index c
in a substa
uency band l
of these fre
onic band ga
re 1.1 Schelength is (a)
exhibit an o
ection to con
contrast betw
antial solid
lies within t
equencies is
ap (PBG).[16-
ematic illus) inside band
omnidirection
nfine the lig
ween the two
angle, or ev
the stop ban
s forbidden
-20]
stration of d gap and (b)
2
nal in-plane
ght in the thi
o materials i
ven in all s
nds for all di
to travel in
light propag) outside ban
bandgap, b
ird dimensio
s large enou
solid angles.
irections in
the crystal,
gating in pnd gap.[21]
but they can
on.[9-13] In 3D
ugh, stopgap
.[14, 15] When
three-dimen
, resulting in
photonic cry
only rely o
D PhCs, if th
s can exclud
n a specifie
nsional space
n a complet
ystals, whos
on
he
de
ed
e,
te
se
3
It is important to realize that PBGs only emerge for specific symmetries.
Researchers have made significant efforts to discover and fabricate PCs with large
PBGs.[22] To date, diamond-like structures have been found to generate some of the
largest gaps for a particular index contrast, and have received extensive attention for their
“champion” photonic quality. The rod-connected diamond PC, a structure formed by
dielectric rods which connect nearest-neighbor sites in the diamond lattice, exhibits an
exceptionally broad gap width: 30% of the mid-gap energy for a refractive-index contrast
of 3.6.[23] Besides the large bandwidth, diamond-based PBGs, which open between low
energy bands, tend to be more immune to lattice defects compared to high-energy-band
PBGs of other structures, such as those based on inverse opals.[17]
1.1.2 Photonic density of states for 3D photonic crystals
Photonic density of states (pDOS) describes the number of available
electromagnetic states in a unit volume at each photon energy. In a homogeneous
medium such as vacuum, pDOS can be calculated by treating light as optical standing
waves in an infinitely large cubic box,[24] in which the standing wave takes the form,
, , , sin sin sin sin . (1.1)
In Eqn. 1.1 L, c and λ stand for the length of the box, speed and wavelength of light,
respectively. Positive integer nx, ny and nz represent the “quantum numbers” of the optical
state in Cartesian coordinates. Substituting Eqn. 1.1 into wave equation gives,
. (1.2)
4
The quantum numbers can be related to the photon energy E by replacing λ with E=hc/λ
in Eqn. 1.2, suggesting that each state corresponds to one unique combination of the
quantum numbers. Thus, the total number of optical states within the photon energy E is,
2 . . (1.3)
The factors “2” and “1/8” in Eqn. 1.3 account for the polarization of light and the
positive nature of the quantum numbers, respectively. Finally, pDOS is the derivative of
the total number of states in unit volume with respect to the photon energy,
(1.4)
which can also be expressed in frequency ω as,
. (1.5)
The above derivation only applies to homogeneous media, which shows a quadratic
relationship between pDOS and frequency. For complex systems with spatial refractive
index variations such as 3D PhCs, the pDOS often takes a general form as,
∑ , , (1.6)
where the frequency , of the first n photonic bands is solved in Maxwell equations for
each k (or optical state) in the first Brillouin zone. The pDOS can be readily obtained by
taking the histogram of ω at an interval | |.[25]
Figurinver
high-
PhC.
its pD
mediu
relati
stopg
PBG
pDOS
energ
conse
re 1.2 Calcurse opals.[25]
Figure 1.
-symmetry k
Although it
DOS is only
um with an
ively low re
gaps in ГL an
also arises b
S is strongly
gy PBG. How
ervation. Hen
ulated photo
.2 shows the
k points in th
t possesses t
y moderately
equivalent
efractive ind
nd ГX direct
between the
y depleted a
wever, the p
nce, 3D PhC
onic band di
e photonic b
he first Brillo
three stopga
y modulated
refractive in
dex contrast
tions betwee
8th and 9th b
at the low-en
pDOS signif
Cs provide an
5
iagram and p
band diagram
ouin zone) an
aps in ГL dir
d at those sto
ndex. This is
t. When we
en the 2nd an
ands, as disp
nergy stopga
ficantly incre
n excellent m
photonic de
m (obtained
nd pDOS of
rection and
opgaps com
s because T
e replace th
nd 3rd bands
played in Fig
ap and comp
eases at the
means to mo
ensity of stat
by only cal
f a 3D titania
one gap in Г
mpared to a h
TiO2 inverse
he titania by
start to over
gure 1.3. In
pletely vani
band edge d
odify the pDO
tes for titani
lculating ω a
a inverse opa
ГX direction
homogeneou
opals have
y silicon, th
rlap and a fu
n this case, th
ished at high
due to energ
OS.
ia
at
al
n,
us
a
he
ull
he
h-
gy
Figurinver
1.2 E
techn
this
categ
assem
1.2.1
writin
writin
and r
re 1.3 Calcurse opals.[25]
Experimenta
The comb
niques has re
section, the
gorized as di
mbly and hol
Direct writi
Computer
ng have bee
ng utilizes s
rapidly solid
ulated photo
al realizatio
bination of n
ealized a var
e available
irect writing
lographic lith
ing
r-controlled
en applied t
ol-gel fluidi
dify in a co
onic band di
n of 3D pho
new theoret
riety of highl
methods fo
g, layer-by-la
hography.
writing tech
to make com
ic inks that c
agulation re
6
agram and p
otonic crysta
ical designs
ly functiona
or fabricatin
ayer assemb
hniques such
mplex 3D f
can readily f
eservoir (Fig
photonic den
als
s and moder
al PhCs in th
ng 3D PhCs
bly, top-dow
h as direct in
fine-scale st
flow through
gure 1.4a).[
nsity of stat
rn micro/nan
he optical wa
s will be r
wn etching, c
nk writing an
tructures. Th
h fine depos
26] Figure 1
tes for silico
no fabricatio
avelengths. I
reviewed an
colloidal sel
nd direct lase
he direct in
sition nozzle
1.4b shows
on
on
In
nd
f-
er
nk
es
a
polym
can s
silico
prope
dielec
inver
FigurSEM
In thi
negat
energ
volum
polym
featur
direct
meric wood-
ubsequently
on, germaniu
erties have a
ctric TiO2
rsion step an
re 1.4 DirecM image of 3D
Direct las
is process, a
tive photores
gy stays belo
me of the
merization.
re size as sm
t ink writing
-pile 3D PhC
y serve as sac
um) to open
also been wr
upon calcin
d thus simpl
ct ink writingD wood-pile
ser writing (D
a tightly focu
sist (Figure
ow the sing
focal po
Although r
mall as ~ 65
g, have been
C that was fa
crificial temp
n up a larg
ritten with s
nation. This
lified the fab
g of 3D PhCe structures.[2
DLW) relies
used laser be
1.5a).[30-34]
le-photon ab
int exceeds
estricted by
5 nm, a few
n achieved w
7
abricated in s
plates for hi
ge PBG.[27,
sol-gel inks
s technique
brication pro
s: (a) schem26]
s on two-pho
eam translate
The key of t
bsorption ed
s the exp
y Abbe’s di
orders of m
with careful s
such manner
gh dielectric
28] 3D PhC
that can dir
exempted
ocedures.[29]
matic of ink d
oton or multi
es in a desig
this techniqu
dge of the p
osure thres
iffraction li
magnitude le
selection of
r. The polym
c material in
Cs with exce
ectly transfo
the additio
deposition pr
i-photon pol
gned pathwa
ue is that the
photoresist, w
shold for
imit, 3D str
ess than thos
the excitatio
mer structure
nversions (e.g
ellent optica
orm into hig
onal materia
rocess and (b
ymerization
ay through th
e laser photo
while a sma
multi-photo
ructures wit
se realized b
on source an
es
g.
al
gh
al
b)
ns.
he
on
all
on
th
by
nd
the p
refine
work
made
struct
limite
exper
Figur(b) SE
photosensitiz
ed this techn
kers have als
e by DLW to
Although
tures with h
ed its usag
riments.
re 1.5 DirecEM image o
zer.[35] Besid
nique to real
so demonstr
o separate lef
direct wri
igh precision
ge only to
ct laser writinof 3D spiral s
des the woo
lize a spiral
rated for the
ft- and right-
iting holds
n, the slow n
low-volume
ng of 3D Phstructures.[31
8
od-pile geom
3D PhC, as
e first time
- circularly p
great prom
nature of its
e productio
hCs: (a) sche1]
metry, Misa
s shown in F
the ability o
polarized lig
mise to crea
s point-to-po
on of devic
ematic of two
awa et al. h
Figure 1.5b
of using 3D
ght.[36]
ate nearly
oint fabricati
ces for pro
o-photon lith
have recentl
b. Gu and co
D chiral PhC
arbitrary 3D
ion has so fa
of-of-concep
hography an
ly
o-
Cs
D
ar
pt
nd
9
1.2.2 Layer-by-layer stacking
Highly-developed conventional 2D lithography has also showed excellent
capability of fabricating 3D PhCs. In this case, each layer of the 3D lattice is
independently defined by e-beam lithography and subsequently stacked together with
precise registration. In particular, Noda et al. reported a wood-pile 3D PhC with a
complete PBG at near IR wavelengths by stacking III-V semiconductor stripes using
wafer fusion and laser-assisted alignment (Figure 1.6).[19] By intentionally removing two
stripes, they created a sharp 90◦ bend waveguide, presenting the first successful optical
circuit embedded in 3D PhCs. Arkawa et al. adopted different strategy when stacking the
3D structures, that each patterned planar component was lifted-off from the substrate and
assembled between three rectangular positioning pins via advanced micromanipulation
(Figure 1.7).[37] In addition, this method allowed them to introduce functional point
defects that contained quantum dots (QDs) in one of the 2D layers and for the first time
study the electron-photon coupling in 3D PhCs. Recently, Arkawa and co-works also
pushed forward this technique to accomplish a low-threshold 3D PhC laser with a record-
high cavity quality factor of ~43,000.[38]
Besides the obvious fact that layer-by-layer assembly requires complicated and
time-consuming processes, which hinders their potential for practical applications, the
major challenge remains for lattice registration, e.g., a spatial resolution of less than 10
nm may be necessary for creating PhCs that have PBGs in the visible spectrum.
Figurunit c
Figurschemwood
1.2.3
altern
signif
PhCs
re 1.6 Layercell and (b) S
re 1.7 Laymatic of 3D d-pile PhCs.[
Top-down e
Advanced
native appro
ficantly simp
s via multi-a
r-by-layer asSEM image
er-by-layer structures e
[37]
etching
d top-down
oaches that
plify the fab
angle reactiv
ssemblies of of stacked w
stacking ofembedded w
semicondu
can offer
brication pro
ve ion etchin
10
f 3D PhCs bywood-pile Ph
f 3D PhCs ith quantum
uctor etchin
large-area a
ocesses. In 2
ng (RIE) on s
y wafer bondhCs.[19]
via high-pm dots and (b
ngs have p
and defect-f
2005, Noda
single-crysta
ding: (a) sch
precision alib) SEM ima
proved to b
free 3D Ph
et al. directl
alline Si sub
hematic of th
ignments: (aage of stacke
be promisin
hCs and als
ly created 3D
bstrates.[39] A
he
a) ed
ng
so
D
As
illustr
45º w
coolin
holes
relati
Figurdoub
micro
subse
from
a new
using
1.9).[
the ap
rated in Fig
with respect
ng was emp
s did not inte
ive to the con
re 1.8 3D sile-angled ion
Electroch
omachining,
equently rem
illuminated
w type of 3D
g a modified
[42] The key
pplied volta
gure 1.8a, tw
t to the surf
loyed to pre
ersect inside
nventional c
ingle-crystaln etching; (b
hemical etch
where Si w
moved by HF
d n-type Si o
D PhCs with
d procedure
of their suc
age to contro
wo RIE step
face. At the
event the def
the crystal,
configuration
lline Si PhC b) SEM imag
hing in HF s
was anodize
F.[40, 41] Depe
r nanometric
h simple cub
for macrop
cess was to
ol and stabili
11
ps were intro
e same time
formation of
producing a
n (Figure 1.8
fabricated bge of the 3D
solution has
d to SiO2 a
ending on th
c pores from
bic symmetry
porous Si ph
carefully mo
ize the micr
oduced to et
e, surface pa
f the side wa
a wood-pile
8b).
by deep-ion D PhC lattice
s been a we
at the wafer-
he Si doping
m p-type Si c
ry was demo
hotoelectroc
odulate the b
ro-sized pore
tch masked
assivation a
alls. The resu
PhC that wa
etching: (a) .[39]
ell-establishe
-electrolyte
g type, micro
can be deriv
onstrated by
chemical etc
backside illu
es formation
Si wafer at
and cryogeni
ulted deep a
as 90º rotate
schematic o
ed tool for S
interface an
ometric pore
ved. Recently
Gosele et a
ching (Figur
umination an
n. Verified b
±
ic
air
ed
of
Si
nd
es
y,
al.
re
nd
by
photo
PBG
electr
optoe
as po
locati
In ad
the an
Figurphoto
onic band str
with except
The fact t
rically cond
electronic de
oint defects,
ions on each
ddition, only
nisotropic fl
re 1.9 SEMoelectrochem
ructure calcu
tional uniform
that the 3D
ductive sing
evices. Howe
gain medium
h PhC lattice
limited crys
ow of etchan
M image of bmical etching
ulations, the
mity over a
PhC realized
gle-crystallin
ever, it is ve
m and wave
e are simulta
stal geometri
nt species du
bird’s eye vig.[42]
12
eir 3D PhCs
large area.
d in the abo
ne material
ry challengin
eguides into
aneously pat
ies can be cr
uring RIE.
iew of 3D m
exhibited a
ove fashions
ls makes i
ng to incorp
those 3D ar
tterned durin
reated by the
macroporous
a more than
are usually
it highly d
porate functio
rchitectures,
ng the top-d
ese approach
s Si structur
4% complet
composed o
desirable fo
onalities suc
, since all th
down etching
hes owning t
es created b
te
of
or
ch
he
g.
to
by
13
1.2.4 Colloidal self-assembly
Self-assembled 3D PhCs (so-called artificial opals), constructed by close-packed
spheres in a face-centered cubic lattice, have attracted enormous attention owning to their
ease of large-area and low-cost fabrication and because their optical properties can be
tuned simply by varying the sphere diameter.[43] Among all the proposed methods to
fabricate opals such as sedimentation,[44] cell confinement,[45] Lagmuir-Blodgett,[46] spin-
coating,[47] etc., the vertical deposition provides the highest optical quality and thus has
been widely adopted by researchers.[48] During the vertical deposition, the evaporation of
the solvent (generally water or ethanol) forces the spheres to align in the meniscus
formed between air, colloidal solution and the vertical substrate (Figure 1.10). Artificial
opals made of polystyrene or silica can be used as sacrificial templates for high refractive
index material inversions (e.g. TiO2, Si, GaAs, etc.) to increase their photonic strength.[49-
51] Recently, extensive efforts to introduce controlled functionalities into colloidal
crystals have also allowed people to thoroughly explore the light-matter interaction in 3D
opals, leading to the realization of emission modification, waveguides, optoelectronic
devices in self-assembled colloidal systems.[51-53]
Inherent to the nature of the self-assembly, colloidal PhCs usually possess high
density of undesired disorders and defects, which degrades their photonic strength.
Another major drawback is that only FCC opals can be easily obtained in a large scale
while methods for other symmetries often produce small or low-quality PhCs.[54]
Unfortunately, FCC PhCs are not ideal photonic structures, since even with the highest
available dielectric contrast (Si to air), they only exhibit a 5% full PBG in the defect-
sensitive high-energy band.[30]
Figurassem
1.2.5
poten
scale
interf
photo
rema
disso
confi
overl
Altho
requi
re 1.10 Colmbly and (b)
Holographi
Holograph
ntially the m
production
ference patte
oresist is cro
ins underex
lved. Up to
gurations. P
lapped non-
ough precise
ired in this c
loidal self-a SEM image
ic lithograph
hic lithogra
most promisin
n of 3D Ph
ern of mult
oss-linked a
xposed at d
o now, holog
ioneered by
coplanar las
e alignment
ase, the geo
assembled Pe of colloida
h
aphy, also c
ng way to re
hCs. At its
tiple collima
and thus ins
destructive
graphic litho
Turberfield
ser beams t
t of the ang
metry of the
14
hCs: (a) schal opal PhC m
called mult
ealize fast, n
heart, this
ated beams
soluble at c
interference
ography has
d et al., the fi
to create th
gle, polariza
e 3D structur
hematic of emade of silic
ti-beam inte
early arbitra
s technique
to photores
onstructive
e locations,
s been most
irst method d
he interferen
ation and po
res can be ea
evaporation-ca spheres.
erference lit
ary, defect-fr
translates
sist. Upon e
interference
which is
tly impleme
directly uses
nce (Figure
osition of e
asily adjuste
-induced sel
thography,
ree and large
the designe
exposure, th
e sites, but
subsequentl
ented in thre
s four or mor
e 1.11).[55, 5
each beam
ed by steerin
f-
is
e-
ed
he
it
ly
ee
re
56]
is
ng
an in
comm
Figurand (
incid
produ
simpl
betwe
the m
(PnP)
photo
1.13)
perio
the a
ndividual las
mon source v
re 1.11 4-beb) SEM ima
As shown
ent beam is
uces a woo
lifies the op
een each bea
most powerf
). In particu
oresist either
).[58-60] As a
dic intensity
bove two te
ser beam. In
via a single o
eam interferage of the 4-b
n in Figure
partially refl
od-pile like
tical setup a
am for achie
ful one amo
ular, inciden
r by imprin
a result, tho
y profile, wh
echniques, P
n another v
optical prism
ence lithogrbeam hologr
1.12, the pri
flected by its
symmetry.[
and also allo
eving more
ong the three
nt light is di
nting or by
se multiple
hich is subse
PnP can offe
15
variant, mult
m.
raphy: (a) scraphic PhCs
ism is placed
four facets.
[57] The sin
ows accurate
complex pat
e, is the ph
iffracted by
laminating
diffracted
quently reco
er a greater v
tiple beams
chematic of with FCC s
d on top of a
The resultin
ngle prism
e control of t
tterns. The
ase mask pr
a 2D gratin
an elastome
beams inter
orded in the
variety of co
s can be de
laser beam ymmetry.[55
a photoresist
ng five-beam
component
the relative
third approa
roximity na
ng created o
eric phase m
rfere and ge
photoresist.
omplex 3D
rived from
configuratio]
t film and th
m interferenc
significantl
optical phas
ach, probabl
ano-patternin
on the top o
mask (Figur
enerate a 3D
Compared t
structures b
a
on
he
ce
ly
se
ly
ng
of
re
D
to
by
addin
also
speci
Figurinside
Figur(PnP)(b) SE
ng more vert
be concurre
ially designe
re 1.12 Prise the prism a
re 1.13 Hol): (a) schemEM image o
tical levels o
ently embed
ed phase mas
sm holograpand (b) SEM
ographic lithmatic illustratof 3D PhCs r
of the diffra
dded into t
sk.[59]
phic lithograM image of 3
hography bating single inrealized by P
16
action grating
the 3D stru
aphy: (a) scD PhCs with
ased on masncident beamPnP approach
g. Additiona
uctures via
chematic of h wool-pile l
skless proximm diffractedh.[60]
ally, artificia
single expo
f laser beamlike symmet
mity field nad into multip
al defects ca
osure from
m propagatiotry.[57]
anopatterninple beams an
an
a
on
ng nd
17
1.3 Controlling spontaneous emission in 3D photonic crystals
1.3.1 Spontaneous emission
Light is typically created in two fashions: spontaneous emission and stimulated
emission. The former process refers to radiation of photons from high energy states to
low energy states with no regard to classical electromagnetic field, which is responsible
for most of incoherent light sources around us such as incandescent bulbs, fluorescent
lamps, LEDs and so on. The latter process entails the amplified radiation triggered by the
electromagnetic field with same wavelength, phase, polarization and direction, leading to
a coherent light source like lasers. In classical theory of light, spontaneous emission (SE)
is described as an irreversible emission of photon into free space with frequency (E2-
E1)/ħ, where E1 and E2 denote the energies of the ground and excited states and ħ is the
Planck’s constant. However, the presence of Planck’s constant clearly suggests that SE is
a quantum mechanical process. Indeed, a proper treatment of SE requires the quantization
of both the energy states and the electromagnetic field, because in reality SE is not an
intrinsic property of the emitters but the result from the interaction between electrons and
the reservoir field (vacuum state).[61] Hence, SE can be modified by tailoring the modes
of the surrounding vacuum field the emitters radiate into. It was first advanced by Purcell
in 1946 that the SE decay rate can be enhanced by placing an emitter in a cavity whose
resonance mode is near the emission frequency.[62] On the other hand, SE can also be
inhibited if the surrounding environments (such as 3D PhCs with full PBGs) do not have
available modes for the radiation to couple into.
Control of SE is central to many processes involved with photon management
including light emitting sources[63, 64], solar energy[65] and so on. For example,
18
manipulation of propagation of SE may facilitate more efficient light extraction in optical
display devices.[66] Spectral redistribution of SE may permit more photons to couple into
useful optical modes in optical cavity devices (e.g. lasers).[38] In the following sections,
Fermi’s golden rule, the basic principle that governs SE, and its variation in the presence
of cavities will be firstly described, followed by a brief review of the progress achieved in
SE control using 3D PhCs.
1.3.2 Fermi’s golden rule
To illustrate the basic mechanism underlying SE and its decay rate, we consider a
two-level atom interacting with a continuum of quantized electromagnetic field modes in
free space.[67] The Hamiltonian of the two-level atom can be written as,
, (1.7)
where σ+ and σ- are the pseudo-spin operators and refer to the upward and downward
electronic transitions in the atom, respectively. The Hamiltonian of the field in free space
takes the form,
∑ , (1.8)
where and are the creation and annihilation operators of the field mode with index
k (stands for both wavevector and polarization). In Schrödinger picture, the interaction
Hamiltonian under dipole approximation and rotation wave approximation is,
∑ ∗ , (1.9)
19
The electron-photon coupling strength is defined as ∙ , in which
⟨1| |2⟩ is the dipole matrix element and r denotes the position of the atom. ,
and V are the permittivity of vacuum, permittivity of the medium and mode volume.
Therefore, the total Hamiltonian follows,
, (1.10)
∑ ∑ ∗ , (1.11)
Initially at time t = 0, the atom is in its excited state | and the field is in the vacuum
state|0 , so the evolved total state at time t is,
|Ψ | , 0 ∑ , | , 1 , (1.12)
Based on the perturbation theory, the equations of motion for the probability amplitudes
and , can be readily obtained by inserting Eqn. 1.12 into Schrödinger
equation,
∑ , , (1.13)
,∗ , (1.14)
Here | | corresponds to the probability of finding the atom in the excited state at
time t. To get an expression that only contains , we integrate Eqn. 1.14 and then
substitute it into Eqn. 1.13, yielding,
∑ | | , (1.15)
Since the reservoir field is continuum, we can replace the summation over k with an
integral as ∑ → , which gives,
20
, (1.16)
According to Weisskopf-Wigner approximation, varies sufficiently slowly
compared to the exponential factor in the time integral in Eqn. 1.16 such that it can be
evaluated at time t and removed from the integral. Also, it is convenient to introduce a
delta function that rewrites Eqn. 1.16 as,
, (1.17)
The last integral is the definition of photonic density of states, which counts the number
of modes per unit volume at a given frequency. The middle integral over can be
simplified as,
lim → , (1.18)
Inserting Eqn. 1.18 back into Eqn. 1.17 follows,
∆ , (1.19)
where
, (1.20)
is the analytical expression for Fermi’s golden rule, which states that the rate of SE
depends on the intrinsic emission dipole d and the density of available external modes the
emitter radiates into. The other term,
∆ , (1.21)
21
is called Lamb shift, representing a small shift in transition frequency. Finally, the time
evolution of transition probability can be derived from Eqn. 1.19,
| | . (1.22)
showing that electrons in the excited atomic state relax via an exponential decay.
1.3.3 Weak electron-photon coupling
When SE decay rate is much greater than electron-photon coupling strength ,
the light-matter interaction stays in the so called “weak coupling” regime, as manifested
in Fermi’s gold rule. In this case, the modes of electromagnetic field are treated as a
reservoir with infinitely short memory on the emitted photons, leading to an irreversible
process. If the emitter is placed in a homogeneous medium with refractive index n, the
SE decay rate can be obtained by inserting the pDOS of the free space Eqn. 1.5 into Eqn.
1.20,
, (1.23)
If the emitter is surrounded by a cavity, the free space pDOS is should be replaced by the
cavity pDOS,
/ , (1.24)
which leads to the SE rate in a cavity as,
. (1.25)
22
In Eqn. 1.25, / is defined as cavity quality factor (Q-factor). For a cavity
tuned near the atomic transition frequency ( 0), Eqn. 1.25 can be readily
simplified to,
, (1.26)
while for a cavity detuned from the emission frequency,
, (1.27)
Eqn. 1.26 and Eqn.1.27 clearly shows the enhancement and suppression of SE when the
atomic transition is on and off resonance with the cavity mode, respectively, which was
first discovered by Purcell and thus called Purcell effect. [62]
Fermi’s golden rule (weak coupling) predicts that for the time large enough so
that energy conservation is established but short enough for the first-order perturbation
theory to hold, the excited atom state relaxes via an exponential decay, whose rate is
determined by both the electronic property of the emitter and pDOS of the surrounding
medium. Therefore, for a given emitter, key to the manipulation of SE is full control of
the pDOS.
23
1.3.4 Control of light emission in 3D photonic crystals
The interaction between emitters and 3D PhCs has been intensively investigated
in the past decade in order to achieve full control of SE in all directions. The early
experiments on SE manipulation were carried out in 3D opal PhCs made of low dielectric
materials and infiltrated with organic dyes. Pioneered by Lawandy and coworkers, the
emission rate of a dye was found to be inhibited by a factor of 1.75 in polystyrene opal
PhCs, compared to that in a disordered structure.[68] Later, Gaponenko et al. reported a
two-fold change of SE rate from dyes dispersed in polymer-coated SiO2 opals.[69]
However, in principle colloidal opals with low dielectric contrast could only weakly
modify SE. As disputed later, the above observations were mainly attributed to the
variation of the chemical environment surrounding the emitters or the emission from the
PhC backbones. The first accurate measurement of SE modification in 3D PhCs was
performed by Vos et al., where they adopted highly efficient quantum dots (QDs) and
selectively infiltrated them inside a high dielectric contrast TiO2 inverse opal.[50] The
sharp emission of individual QD allowed them to precisely probe the modulated pDOS in
3D PhCs. Both pronounced enhancement and suppression of SE were observed and
quantitatively compared to a better-suited reference (Figure 1.14), a small-lattice PhC
whose stopgap was on the blue side of the QD emission. Besides dyes and quantum dots,
efforts have also been spent on incorporation of rare-earth emitters into 3D opal PhCs.[70-
72] In particular, van Veggel et al. successfully doped Er2+ into an inverse opal composed
of GaN/SiO2 composites via a simple solid state reaction. Their experiment revealed a
three-fold modification of Er2+ emission rate induced by the strong stopgap of the PhC
host (Figure 1.15).[73]
Figurembe
Figurrelati
epitax
These
excel
has y
re 1.14 (a) edded in titan
re 1.15 (a) ive measured
Control o
xial quantum
e experimen
llent agreem
yet been don
SEM of titania inverse o
Optical imad emission li
of SE by 3D
m wells and
nts were p
ment with the
ne with artifi
ania inverse opals with va
age of Eu2+
ifetime of Eu
PhC with a
d dots with
erformed in
e weak-coup
icial PhCs th
24
opal PhCs. arious lattice
+ doped comu2+ embedde
full PBG ha
the layer-by
n the near-
pling theory
hat operate i
(b) Emissioe constants.[5
mposite GaNed in GaN Ph
as also been
y-layer asse
-infrared wa
.[37, 74] Howe
in the visibl
on decay cu50]
N inverse ophCs.[73]
n realized by
embled wood
avelengths a
ever, such d
le spectrum,
urves of qdot
pals. (b) Th
y coupling th
d-piles PhC
and reveale
demonstratio
owing to th
ts
he
he
s.
ed
on
he
challe
Natur
archit
green
diam
the ex
that e
the m
photo
Figurwing 1.4 3D
are al
enges for fa
re, on the ot
tectures opti
n coloration
ond-based 3
xoskeleton c
exhibits 2.5%
measured rad
onic band ed
re 1.16 (a) s. (b) Compa
D photonic
Besides p
lso particula
abricating w
ther hand, lo
imized for v
of one of
D PC structu
can be used
% full PBG
diative lifetim
dge to 99±2 n
SEM imagearison of em
crystals for
photonic app
arly attractiv
wood-pile or
ong ago dev
visible wave
the weevil
ure of its exo
as a templat
in the visib
mes range o
ns within the
e of inversemission decay
r energy sto
plications, th
ve to electroc
25
r other diam
veloped an e
elengths. In 2
beetles com
oskeleton.[75
te for creatin
ble. Upon em
over a factor
e PBG (Figu
e diamond-liy of qdots in
rage applic
he periodic a
chemical sys
mond-like P
enormous pa
2009, Bartl
mes from li
5] Recently t
ng high refra
mbedding Q
r greater tha
ure 1.16).
ike structuren different ph
cations
and bicontinu
stems, wher
PhCs with s
alette of elab
et al. discov
ight reflectio
this group h
active-index-
QDs into tho
an 10, from
e obtained fhotonic envi
uous natures
re large surfa
small lattice
borate 3D P
vered that th
ons from th
as shown tha
-contrast Ph
ose structure
8±1 ns at th
from beetleronment.[76]
s of 3D PhC
face areas an
s.
C
he
he
at
C
s,
he
’s
Cs
nd
effect
these
place
densi
electr
densi
a 94%
confo
opal
ionic
active
nanop
2X g
repor
Figurobtain
tive pathway
merits, por
e for themse
ity batteries.
rodes can of
ity by simult
% porous m
ormally coat
structures ca
transport le
e materials
porous elect
greater energ
rted previous
re 1.17 (a) ned from sel
ys for mass
rous electrod
elves on the
.[77] For insta
ffer capacito
taneously mi
metallic curr
ted with batt
an provide d
engths to bo
for achievin
trodes to inte
gy density a
sly.[79]
Schematic lf-assembled
and charge
des obtained
e roadmap t
ance, Zhang
or-like powe
inimizing all
rent collecto
ery active m
deterministic
oost power
ng high ene
erdigited mic
and 2000X g
and (b) SEd colloidal op
26
transports a
d from 3D P
to realize h
g et al. have
er density w
l the resistan
or was deriv
materials. As
c porous cha
and also co
rgy. Later,
crobatteries.
greater pow
EM image opals.[78]
are often hig
PhC templa
high energy
e recently de
while mainta
nces inside b
ved from a
shown in F
annels that re
oncurrently a
Pikul et al.
. As a result,
wer density t
of High-pow
ghly favored
ates have rap
density and
emonstrated
aining battery
batteries.[78] I
3D opal P
igure 1.17,
educe both e
allow a larg
applied thi
, their batter
than other m
wer 3D LiM
d. Because o
pidly made
d high powe
that such 3D
y-like energ
In their work
PhC and the
such inverte
electronic an
ge loading o
s type of 3D
ries possesse
microbatterie
MnO2 cathod
of
a
er
D
gy
k,
en
ed
nd
of
D
ed
es
de
27
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29
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30
CHAPTER TWO
CONTROL OF SPONTANEOUS EMISSION IN
3D SILICON PHOTONIC CRYSTALS *1
2.1 Introduction & motivation
Compared to 1D and 2D PhCs, which inherently cannot control SE in all three
directions, 3D PhCs hold promise for complete control over SE. In recent years, SE
control by colloidal crystals and TiO2 inverted opals has been demonstrated.[2-7]
However, little work on SE has been conducted on silicon or other high refractive index
contrast inverse opal PhCs that exhibit strongly modified photonic DOS.[8, 9] The emitters
commonly used in the previous SE control studies were either organic dyes or colloidal
quantum dots. Although these materials exhibit high quantum emission efficiency, these
emitters are not ideal for several reasons: (i) the generally broad emission of dyes does
not allow investigation of sharp DOS features in an opal PhC;[5] (ii) both dyes and
colloidal quantum dots often suffer from photo oxidation and bleaching;[10, 11] (iii) and
perhaps most importantly, they are generally introduced into opal PhCs by infiltration
and thus randomly infill the PhCs, including the near-surface and surface of the PhC
where of course their optical properties are not influenced as much by the PhC. In this
letter, we demonstrate incorporation of LaF3: Nd nanoparticles at well-defined locations
in a silicon inverse opal. We probe the sharp DOS features of a silicon inverse opal PhC
by studying the stable narrow emission of the embedded Nd3+ at 1052 nm as a function of
* Content in this chapter was previously published by the author and reproduced with permission.[1]
31
silicon filling fractions and find that the radiative lifetime of the embedded particles is
strongly correlated to photonic DOS.
2.2 Fabrication of silicon photonic crystal sandwich structures
Accurate PhC SE studies require that the emitter be stable and have high quantum
efficiency as well as a narrow emission linewidth relative to the DOS features.
Neodymium-doped lanthanum fluoride is an excellent candidate in these respects. The
emission of Nd ion arises from parity forbidden transitions between energy levels of 4f
electrons and thus has very narrow (~20 nm) emission bands.[12] The LaF3 host provides
protection and also endues the Nd ions with high quantum yield.[13] For this study, the
LaF3: Nd nanoparticles were synthesized following a modified version of van Veggels’s
procedure.[14] The synthesis was carried out at 90°C rather than 75°C to increase the
particle size. After formation of the LaF3: Nd nanoparticles, an extra 7 nm of LaF3 was
grown on the shell to increase the luminescence yield.[15] After purification, the resulting
nanoparticles could be dispersed in isopropanol, and have an average diameter of 50 nm.
It is known that the Nd3+ quantum efficiency increases as the Nd doping concentration is
decreased.[16] We find that 0.5% Nd doped LaF3 is optimal taking into consideration both
quantum yield and brightness.
One of the outstanding difficulties in studying SE manipulation in PhCs is an
approach to place the emitters at well-defined locations within the three dimensional
structure. In our procedure, a thin layer of LaF3: Nd particles is introduced into the
colloidal crystal as a planar defect parallel to the <111> opal crystal plane. As shown in
Figure 2.1, the fabrication of the sandwich structure started with the growth of a silica
collo
assem
silica
solve
was f
crysta
was g
crysta
layer
able
length
Figurnanop
idal crystal
mbly.[13] The
a dispersion
ent was evap
formed by sp
al film follo
grown over
al. The seco
(< 100 nm)
to direct the
hwise into 4
re 2.1 Schparticles into
on a sapphi
e substrate w
of 442 nm d
porated leavi
pin-coating t
owed by heat
the nanopa
ond crystal g
) preserved
e growth of
4 pieces; thes
hematic illuso silicon inv
ire substrate
was placed i
diameter sili
ing behind a
the nanopart
ting at 150°
article layer
grew epitaxi
the surface
f the second
se pieces we
stration of erse opals.
32
e (1.5 cm ×
in a 15 mL
ica spheres
a colloidal cr
ticle solution
C for 30 mi
using the s
ially on the
morphology
crystal. The
ere then filled
steps (a th
2 cm) by e
scintillation
(3wt% in et
rystal of 5-8
n (3wt%) on
in.[17] A seco
same growth
first crystal
y of the firs
e resulting s
d with differ
hrough f) to
evaporation-
n vial which
thanol) at 32
layers. The
n top of the
ond colloida
h conditions
, as the thin
st opal film
sample was
rent amounts
o incorpora
-induced sel
h contained
2°C, until th
emitter laye
first colloida
al crystal film
s for the fir
n nanoparticl
and thus wa
carefully cu
s of silicon.
ate LaF3: N
f-
a
he
er
al
m
st
le
as
ut
Nd
33
To tune the filling fraction of the silicon inverse PhCs, each piece of colloidal
crystal was exposed to a different number of Al2O3 deposition cycles via atomic layer
deposition (ALD). This Al2O3 layer served to reduce the amount of silicon deposited in
the next step and thus provided control over the spectral position of the optical features.
In each ALD cycle (Savannah 100, Cambridge NanoTech) the water and
trimethylaluminium (TMAl) exposure times were 0.05 s and 0.10 s, respectively,
followed by 65 s pump time. The chamber was kept at 80°C during the deposition
process. Each ALD cycle conformally deposited 1.2±0.2 Å alumina over the template.
The 4 pieces were exposed to 10, 20, 30 and 40 ALD cycles, and subsequently
concurrently infiltrated with amorphous silicon by means of static chemical vapor
deposition (CVD) at 350°C for 5 hours using Si2H6.[18] Each opal presented an overlayer
of silicon, indicating that the pinch-off of the templates had been reached (all colloidal
crystals were maximally filled). The excess silicon was then removed by reactive ion
etching as previously described,[19] and the silica and alumina were concurrently removed
by immersing the samples in dilute HF (5% HF in a 50 : 50 ethanol : water mixture) for
30 min, followed by rinsing in isopropanol and drying at 60°C. The final result were
emitter-embedded silicon inverse PhCs of 20.8%, 19.7%, 17.9%, 16.6% silicon filling
fractions, starting from the assumption that 22.4% is the maximum filling fraction for an
inverse opal;[20] these values are extracted from calculations as will be discussed. A cross
section SEM image of the sandwich silicon inverse opal is presented in Figure 2.2a and
Figure 2.2b provides a closer look at the LaF3: Nd nanoparticle layer.
FiguremittNd na
2.3 O
2.3.1
with
show
lattic
Perot
value
crysta
cycle
to dip
incre
layer
re 2.2 Scannter-embeddeanoparticle l
Optical char
Reflection &
The refle
varying silic
w intense pe
e throughou
t interference
es, which is
al and went
es. The -L s
ps in spectr
asing light s
s).
ning electrod silicon invlayer.
racterization
& transmissi
ctance and t
con filling fr
aks arising
ut the first an
e. The reflec
consistent w
through ide
stopgap is al
ra. On the b
scattering an
on micrograpverse opal a
ns
ion measure
transmittanc
ractions (SF
from the
nd second cr
ctance peaks
with the fact
entical fabric
lso evident i
blue side of
nd absorption
34
ph (SEM) imand (b) an en
ments
ce spectra of
F) are given
-L stopgap,
rystals, as w
s of the four
that they we
cation proced
n the transm
f the stopga
n by silicon
mages of (a)nlarged view
f the sandw
n in Figure 2
evidence o
well as secon
r samples m
ere originall
dures except
mittance mea
aps, the low
through the
) the cross sw of the emb
wich silicon
2.3. The refl
of a highly
ndary lobes d
manifest simi
ly cut from t
t for the num
asurements c
w transmittan
e thick opal
section of thbedded LaF
inverse PhC
flectance plot
ordered opa
due to Fabry
lar maximum
the same opa
mber of ALD
correspondin
nce is due t
crystals (~1
he F3:
Cs
ts
al
y-
m
al
D
ng
to
14
Figurembeverticnanopshowfrequ
re 2.3 Refledded siliconcal line reprparticles; in
wn in arbitrauency, where
lectance (ren inverse opaesents the 1gray is the
ary units. Ae lattice cons
ed line) andals with decr052 nm cennanoparticl
All spectra stant a = 625
35
d transmittanreasing SFF nter wavelenle emission
are plotted5 nm.
nce (blue lfrom bottom
ngth of the espectrum. T
d in both w
line) spectram to top. Theemission of The emissionwavelength
a of emittere dashed grathe LaF3: Nn spectrum and reduce
r-ay
Nd is
ed
36
As the alumina coating is increased from 10 cycles to 40 cycles, the filling
fraction of silicon inversion is reduced from 20.8% to 16.6%, causing the stopgap of the
silicon inverse PhCs to blue-shift about 0.05 in reduced frequency units (a/, where a =
625 nm, is the lattice constant of the opal crystal), corresponding to a 100 nm wavelength
shift, which allows finely tuning the relative spectral position between the stopgap and
the 1052 nm Nd3+ emission line (gray dashed line).
2.3.2 Spectral & time-resolved photoluminescence measurement
Both spectral and time-resolved experiments were performed to study the
influence of silicon inverse PhCs on the spontaneous emission of the sandwiched LaF3:
Nd nanoparticles. The optical setup is shown in Figure 2.4. The laser used was a tunable
Ti: Sapphire CW (continuous wave) laser (Spectra-Physics 3900S) pumped by a 2 W Ar
ion CW laser operating at the all line mode. The output wavelength of the Ti: Sapphire
laser was set at 796 nm with an average power of 100 mW. The excitation laser was then
modulated by an optical chopper at 100 Hz and focused onto the sample by a 60×
objective with a 0.85 numerical aperture. The emission from the sample was collected by
the same objective and directed into a monochromator which was set to 1052 nm with a
spectral resolution of 50 nm. To obtain the emission decay curve, an oscilloscope was
used to record the signal from the NIR PMT detector (Hamamatsu H10330A-75) after the
excitation was blocked by the chopper. The temporal resolution for such lifetime
measurement is estimated to be about 15 µs taking into account the system responses and
the modulation of the excitation beam of a finite size. The PL spectra were obtained using
the same setup but with a spectral resolution of 6nm. To ensure that PL and lifetime
meas
meas
light
can b
Figur
(PL);
emiss
differ
to ex
concl
it has
urements w
urements by
from the sam
be imaged an
re 2.4 Exper
Silicon a
; however,
sion from th
rent photoni
xtract quan
lusions since
s been repor
were perform
y Fourier tran
mple surface
nd compared
rimental con
also absorbs
as displayed
he Nd3+. The
c environme
ntitative con
e the PL me
rted that the
med at the s
nsform infra
e was directe
d to those obt
nfigurations o
at this exc
d in the Fig
e PL spectra
ents; to conf
nclusions. P
asurements w
emission ou
37
ame spots a
ared spectros
ed into a CC
tained from
of PL and lif
citation and
gure 2.5a,
a suggest th
firm this, tim
PL intensity
were perform
ut of the pho
as the reflec
scopy (FTIR
D camera so
the microsc
fetime meas
exhibits a w
the PL spe
hat the PL in
me-resolved
y alone is
med over fin
otonic structu
ctance and
R), the scatter
o that spots o
ope on FTIR
urements.
weak photol
ectra were d
ntensity was
measuremen
not suffici
nite collectio
tures has a s
transmittanc
red excitatio
on the sampl
R.
luminescenc
dominated b
s modified b
nts were use
ient to draw
on angles an
trong angula
ce
on
le
ce
by
by
ed
w
nd
ar
38
dependence.[21] The time-resolved data (Figure 2.5b), measured at a center wavelength of
λ = 1052 nm with a 50nm spectral resolution, provide direct evidence of emission
manipulation by PhCs through the differences in the radiative lifetime. All decay curves
were fitted by a double exponential function.[22] The luminescence decay times for the
silicon inverse opals with a = 625 nm are 352±13 µs (20.8% SFF), 418±15 µs (19.7%
SFF), 565±20 µs (17.9% SFF) and 501±15 µs (16.6% SFF). A possible explanation on
the multiple decay channel behavior is that the probabilities of the radiative emission of
Nd ions are different near the particle surface and in the core. It would be rather
challenging to investigate the influence of photonic environments on each decay channel
individually. In this work, we used the averaged emission lifetime that was previously
employed to study the quantum yield of Nd3+ doped nanoparticles. The averaged
radiative lifetime can be calculated as follows:
τ
, (2.1)
where A1, A2 , t1, and t2 are fitting constants. For example, the fitting result of 17.9% SFF
data is shown in Figure 2.6. The reduced chi-square and averaged lifetime are 7.43 ×10-5
and 565 µs, respectively.
Figurwavefabric
FigurChi-s
re 2.5 (a) Eelength of 10cated with d
re 2.6 Emissquare is 7.4
Emission spe052 nm for ifferent SFF
sion decay c3 ×10-5.
ectra and (bLaF3: Nd n
F.
curve fitting
39
b) emission nanoparticle
g of 17.9% S
decay profis embedded
SFF data (a
les measured in silicon
= 625 nm).
ed at a centeinverse opa
The reduce
er ls
ed
40
To qualitatively determine the effect of photonic crystals on SE control, a
reference system, e.g., a photonic crystal with a smaller lattice parameter in which the
emission is on the red side of the stopgap, has been previously utilized.[23] We
constructed a similar reference structure with embedded emitters from 240 nm silica
spheres (a = 340 nm) which was subsequently coated using 10 cycles of alumina ALD
and filled with Si via CVD. The ALD and Si infilling procedure was applied to ensure the
nanoparticles in the reference sample were in a similar chemical environment as the other
samples and to ensure all samples had been similarly processed. The SFF value and
lifetime of the reference were found to be 20.0% and 390±16 µs. Compared with this
reference sample, the emission is enhanced for the inverse opal with a = 625 nm and
20.8% SFF and is suppressed for the samples with 17.9% SFF and 16.6% SFF, and about
the same as the 19.7% SFF sample .
2.4 Photonic density of states correlations
In the opal sandwich structures, the thin LaF3: Nd particle layer that has a
refractive index different from the surrounding materials, which introduces a planar
defect within the PhC. It has been shown that such a structural defect can create a spectral
defect state in PBGs, manifesting itself as a dip (spike) within Bragg peaks in reflectance
(transmittance) spectra.[24, 25] The spectral position of the state is determined by the
refractive index and thickness of the defect material and the optical properties of the
surrounding PhC. Regardless of the thickness of the planar defect, a defect state will
always be produced in any PBGs and periodically sweeps the entire PBG as the thickness
of the defect layer increases.[26] For our samples, the defect state is not apparent in the
41
reflectance and transmittance spectra. It might overlap with the tail of the Bragg peak on
the high energy side given the fact that the ratio of the emitter layer thickness to the
lattice constant is small (~ 0.1). Most important for this work, the defect mode does not
overlap with the Nd3+ emission spectrum. Thus, we can neglect the influence of the
defect state on the SE of the emitters and approximate the sandwich structure as a perfect
infinite silicon inverse opal PhC in the band structure and DOS calculations.
The photonic band structures and density of states were calculated using the MIT
Photonic Bands computer program to interpret the dynamics of the emission measured in
the silicon inverse PhCs (a = 625 nm).[27] The -L stopgap for each PhC was fitted to
their respective reflectance peak (Figure 2.3) by adjusting within the experimental
uncertainty the thickness of silicon coating and the size of interconnecting windows
between two adjacent opal lattices. The effective refractive index was derived from the
dispersion of the first photonic band close to the zone center. The fitted parameters
served as inputs to simulate the corresponding photonic DOS spectrum, which was
carried out using 23508 k-points that were evenly distributed in the irreducible Brillouin
zone and carefully weighted.[28, 29] The silicon filling fraction was calculated by modeling
the silicon inverse structures using Mathematica with the fitted parameters and
numerically integrating over the silicon volume. Figure 2.7 shows the evolution of the
DOS for structures with varying SFF and the PL spectrum of Nd3+. As the SFF decreases,
while the shape of the DOS curves remains the same, the sharp DOS features gradually
migrate to the high energy side, enabling the Nd emission to probe different DOS
regions. For example, the PL (centered at a/λ = 0.59) lines up with the high DOS peak
near the -L band edge for the case of 20.8% SFF, which is consistent with the enhanced
42
emission (short lifetime) observed experimentally for this sample. For the samples of
17.9% SFF and 16.6% SFF, in which inhibited emission (long lifetimes) were observed,
the PL falls within the strongly depleted DOS region of the stopgap. The difference in the
measured lifetime for those two cases suggests a variation in the photonic strength within
the stopgap, where the former presents a lower DOS than the latter according to the
simulation.
Since structural fluctuations among inverse opals of different batches during the
fabrication (e.g. crystal thickness, location of emitters and density of crystal defects)
might affect the optical measurements, a quantitative analysis is directly performed
among these samples with a = 625 nm that were cut from the same crystal and thus have
identical structural conditions; the only difference between the samples is the amount of
Al2O3 in the intermediate step and the resulting SFF. To examine the influence of the
Al2O3 ALD on Nd3+ emission (although Al2O3 was completely removed in the oxide
etching step for the sample studies), time-resolved emission measurements were
performed on two control samples with similar SFF but different amount of ALD. We
fabricated another small lattice parameter reference sample (which also presents no DOS
modification at 1052 nm Nd3+ emission peak) from 320 nm silica spheres (a = 453 nm)
and 20 cycles of Al2O3 ALD. The corresponding SFF and emission lifetime of this
sample were found to be 19.3% and 403±17 µs. In Figure 2.8, the measured emission
decay profile of a = 453 nm sample is plotted together with that of a=340nm reference
sample which was made with 10 cycles of ALD and previously mentioned. Although
these two samples were constructed with different amount of ALD, they showed similar
radiative lifetime, suggesting that ALD processing does not alter Nd3+ emission
43
properties. The sample with 19.7% SFF is used as a reference because its emission is the
least modified according to the small lattice reference (a = 340 nm). Table 2.1 displays
the measured emission rates (inverse lifetime) and the corresponding DOS values
(integrated over the emission spectrum of Nd3+) for various SFF normalized by 19.7%
SFF data. The emission rate and DOS for each sample are highly correlated: the change
of the relative emission rate ranges from 19% SE enhancement (20.8% SFF) to 26% SE
inhibition (17.9% SFF); the relative DOS changes for these two cases are a 77% increase
and 68% reduction, respectively. To examine if the photonic crystal effect is responsible
for the SE change rather than differences in the effective refractive index of each sample
due to the variation in SFF, we modeled the PhC as a homogeneous medium (HM) with
the effective refractive index obtained from the band structure calculations, and the SE
rate in such a homogeneous medium is proportional to the refractive index.[30, 31] We
estimated the change of the emission rate due to the variation in the SFF, which is found
to decrease monotonically in a very small magnitude (≈ 2% per 1% SFF), e.g. the
emission rate in a HM increases 2.0% when SFF increases from 19.7% to 20.8% and
decreases 4.3% and 7.0% as SFF is reduced from 19.7% to 17.9% and 16.6%. A much
larger change of emission rate was experimentally observed among these samples,
indicating the decisive factor for controlling SE was the photonic crystal.
Figurdecresimul1052 nanop
re 2.7 Simeasing SFF lated -L ganm center w
particle emis
mulated DOSfrom bottom
ap positions wavelength ossion spectru
S of siliconm to top (bfor each ca
of the emissum (shown i
44
n inverse opblack lines)ase. The dassion of the Lin arbitrary u
pals (a = 6). The shadshed gray veLaF3: Nd nanunits).
625 nm) fabowed areas ertical line rnoparticles;
bricated witrefer to th
represents thin gray is th
th he he he
FigurLaF3
nm (b
Tablemissinver
re 2.8 Emis: Nd nanopablack) and 3
e 2.1 Measusion rates inrse opals with
ssion decay articles embe40 nm (red)
SFF Γ
20.8%
19.7%
17.9%
16.6%
ured emission homogeneh a = 625nm
profiles meedded in sili, respectivel
Γ/ Γ19.7%SFF D
1.19
1.00
0.740
0.834
on rates (Γ)eous media
m, normalize
45
easured at aicon inverse y.
OS/DOS19.7%SFF
1.77
1.00
0.324
0.523
), calculatedwith effect
d by 19.7%
a center wav opals with
F ΓHM/ΓHM_19.7
1.02
1.00
0.957
0.930
d DOS valuetive refractiSFF data.
velength of lattice param
7%SFF
es (DOS) anive indices
1052 nm fometers of 45
nd calculate(ΓHM) of th
or 53
ed he
46
Although the general trends in the experimental SE data and the DOS calculations
agree, the magnitude of the SE rate change is less than the variation of the DOS. This is
due to various non-idealities in the experiment: the self-assembled PhCs contain defects
which degrade their photonic strength; the double exponential decay of the emission
suggests the presence of a non-radiative decay process, and the experimental samples
have finite thickness.
2.5 Conclusions
In conclusion, we have demonstrated that the photonic band structure of silicon
inverse PhCs can be specifically and finely tailored using ALD at an intermediate step
between template fabrication and silicon inversion. This is coupled with the incorporation
of rare earth nanoparticle emitters into the silicon inverse PhCs at a well-defined location
provided by a simple experimental procedure. We utilize the narrow emission linewidth
of rare earth nanoparticles and the filling fraction tuning enables to study the effect of the
photonic DOS on spontaneous emission. Time-resolved experiments reveal that the
emission rate of embedded emitters can be strongly manipulated by the stopgap of silicon
inverse opals; up to a 61% change of emission decay rate is observed between the
enhanced and the inhibited SE.
47
2.6 References
[1] H. L. Ning, A. Mihi, J. B. Geddes, M. Miyake, P. V. Braun, Adv Mater 2012, 24, Op153. [2] E. P. Petrov, V. N. Bogomolov, I. I. Kalosha, S. V. Gaponenko, Phys. Rev. Lett. 1998, 81, 77. [3] S. V. Gaponenko, V. N. Bogomolov, E. P. Petrov, A. M. Kapitonov, D. A. Yarotsky, I. I. Kalosha, A. A. Eychmueller, A. L. Rogach, J. McGilp, U. Woggon, F. Gindele, J. Lightwave Technol. 1999, 17, 2128. [4] S. Kubo, A. Fujishima, O. Sato, H. Segawa, J PhysChem C 2009, 113, 11704. [5] I. S. Nikolaev, P. Lodahl, W. L. Vos, J Phys Chem C 2008, 112, 7250. [6] I. S. Nikolaev, P. Lodahl, A. F. van Driel, A. F. Koenderink, W. L. Vos, Phys Rev B 2007, 75, 115302. [7] P. Lodahl, A. Floris van Driel, I. S. Nikolaev, A. Irman, K. Overgaag, D. Vanmaekelbergh, W. L. Vos, Nature 2004, 430, 654. [8] A. Blanco, E. Chomski, S. Grabtchak, M. Ibisate, S. John, S. W. Leonard, C. Lopez, F. Meseguer, H. Miguez, J. P. Mondia, G. A. Ozin, O. Toader, H. M. van Driel, Nature 2000, 405, 437. [9] E. Palacios-Lidón, A. Blanco, M. Ibisate, F. Meseguer, C. López, J. Sánchez-Dehesa, App Phys Lett 2002, 81, 4925. [10] G. van den Engh, C. Farmer, Cytometry 1992, 13, 669. [11] W. G. J. H. M. van Sark, P. L. T. M. Frederix, D. J. Van den Heuvel, H. C. Gerritsen, A. A. Bol, J. N. J. van Lingen, C. D. Donega, A. Meijerink, J Phys Chem B 2001, 105, 8281. [12] G. A. Kumar, C. W. Chen, J. Ballato, R. E. Riman, Chem Mater 2007, 19, 1523. [13] K. S. Upadhyaya, R. K. Singh, Journal of Physics and Chemistry of Solids 1975, 36, 293 [14] J. W. Stouwdam, F. C. J. M. van Veggel, Nano Lett 2002, 2, 733. [15] X. F. Yu, L. D. Chen, M. Li, M. Y. Xie, L. Zhou, Y. Li, Q. Q. Wang, Adv Mater 2008, 20, 4118. [16] M. C. Tan, G. A. Kumar, R. E. Riman, M. G. Brik, E. Brown, U. Hommerich, Journal of Applied Physics 2009, 106, 063118. [17] R. Pozas, A. Mihi, M. Ocana, H. Miguez, Adv Mater 2006, 18, 1183. [18] S. A. Rinne, F. Garcia-Santamaria, P. V. Braun, Nat Photon 2008, 2, 52. [19] F. Garcia-Santamaria, E. C. Nelson, P. V. Braun, Phys Rev B 2007, 76, 075132. [20] Y. A. Vlasov, X.-Z. Bo, J. C. Sturm, D. J. Norris, Nature 2001, 414, 289. [21] A. F. Koenderink, L. Bechger, H. P. Schriemer, A. Lagendijk, W. L. Vos, Phys. Rev. Lett. 2002, 88, 143903. [22] G. A. Kumar, E. D. la Rosa-Cruz, K. Ueda, A. Martínez, O. Barbosa-García, Optical Materials 2003, 22, 201 [23] E. Bovero, F. C. Van Veggel, J Am Chem Soc 2008, 130, 15374. [24] N. Tetreault, A. Mihi, H. Miguez, I. Rodriguez, G. A. Ozin, F. Meseguer, V. Kitaev, Adv Mater 2004, 16, 346. [25] E. Palacios-Lidon, J. F. Galisteo-Lopez, B. H. Juarez, C. Lopez, Adv Mater 2004, 16, 341.
48
[26] E. Yablonovitch, T. J. Gmitter, R. D. Meade, A. M. Rappe, K. D. Brommer, J. D. Joannopoulos, Phys Rev Lett 1991, 67, 3380. [27] S. Johnson, J. Joannopoulos, Opt Express 2001, 8, 173. [28] J. Hama, M. Watanabe, T. Kato, J Phys-Condens Mat 1990, 2, 7445. [29] M. R. Jorgensen, J. W. Galusha, M. H. Bartl, Phys Rev Lett 2011, 107, 143902. [30] A. M. Fox, Quantum optics: an introduction, Oxford University Press, 2006. [31] L. A. Stewart, Y. Zhai, J. M. Dawes, M. J. Steel, J. R. Rabeau, M. J. Withford, Opt Exp 2009, 17, 18044.
49
CHAPTER THREE
3D SILICON PHOTONIC CRYSTAL MICROCAVITY
3.1 Introduction & motivation
3D Photonic crystals (PhC) are considered one of the most promising means to
realize miniaturized low threshold lasers. This is because 3D PhCs microcavities can
possess both a high cavity quality factor and a small mode volume. If the 3D PhC has a
full photonic band,[1-5] spontaneous emission at wavelengths other than cavity modes is
suppressed and consequentially the lasing threshold is reduced via the Purcell effect.[6]
However, limited work has been done on 3D PhC lasers due to difficulties in fabricating
defect cavities and incorporating gain medium into the 3D PhC network. Reference[7]
demonstrated a 3D PhC laser with a record-high cavity Q factor and ultralow lasing
threshold using a layer-by-layer assembly method. Such device is expensive and difficult
to scale due to the complex fabrication procedures. Lasing was also realized in polymer-
based colloidal PhC systems,[8, 9] in which two polymeric 3D PhCs sandwiched a thin
film of dye-doped polymer that served as the lasing cavity and gain medium. However,
polymer PhCs can only weakly interact with light and manifest a weak photonic band gap
with strong angular dependence due to their low refractive index modulation. Thus, these
3D structures actually claim few advantages over their 1D and 2D counterparts.[6, 10, 11] In
this work, our goal is to develop a new route to fabricate 3D PhC microcavities using
high dielectric contrast silicon inverse opal PhCs. Our method paves the way for
achieving 3D PhC lasers that utilize Si photonics to provide cavity resonances and III-V
compound semiconductors to serve as the gain media.
3.2 D
cavity
conve
confi
13] Ho
high
light
FigurPhC,
Design of sili
As shown
y sandwiche
entional ver
nement com
owever, we
dielectric co
over a large
re 3.1 Schema ½ lambda
icon inverse
n in Figure
ed between
rtical cavity
mes from 1D
expect that
ontrast 3D P
range of an
matic of a 3a defect and a
e opal micro
3.1, the 3D
two Si inve
surface emi
PhC (DBR)
3D PhC ver
PhCs can h
gles.[14, 15]
3D photonic a bottom 3D
50
ocavities
D PhC micro
erse opal Ph
itting laser (
) and light c
rtical cavities
ave 3D pho
crystal micD Si PhC.
ocavity is c
hCs. This st
(VCSEL), w
an be trappe
s outperform
otonic bandg
rocavity tha
omposed of
tructure is s
where the top
ed in vertica
m those with
gap and stro
at consists o
f a ½ lambd
similar to th
p and bottom
al direction.[1
h 1D PhCs, a
ongly confin
f a top 3D S
da
he
m
12,
as
ne
Si
51
The above 3D Si PhC vertical cavity is simulated in commercial finite difference
time domain (FDTD) software from Lumerical Solutions. Figure 3.2a shows the
simulated cavity mode as a function of the cavity thickness, where the top and the bottom
PhCs both have 9 layers of lattices and the refractive index of the cavity material is 3.5.
Similar as the conventional DBR cavity, the cavity mode in our design periodically
sweeps the photonic bandgap with increasing cavity thickness. The order of cavity mode
m also scales with the cavity thickness L as mλ = L, e.g. the defect forms a ½ lambda
cavity when its thickness ranges from 150 nm to 350 nm. The simulated quality factor
(Q-factor) of the microcavity is shown Figure 3.2b. The decrease of the Q-factor at the
band edge is due to the reduced reflectivity at those wavelengths. Here the calculated Q-
factor can only predict its experimental upper limit, since the real value is purely
determined by the quality of the fabricated silicon photonic crystal in experiments.
Figure 3.2 Simulated (a) cavity mode and (b) Q-factor of a 3D Si opal PhC microcavity with 9 periods of lattice for both top and bottom PhCs as a function of cavity thickness. The gray area in (a) stands for the photonic band gap of the Si inverse opal.
100 200 300 400 500 600 700 800
1450
1500
1550
1600
1650
1700
1750
Wav
elen
gth
(nm
)
Defect thickness (nm)100 200 300 400 500 600 700 800
0
2000
4000
6000
8000
10000
Q-f
acto
r
Defect thickness (nm)
(a) (b)
52
3.3 Fabrication of silicon photonic crystal microcavities
The 3D PhC microcavity is assembled by transfer-printing, as illustrated in
Figure 3.3. A bottom Si inverse opal PhC with 15 layers (7 μm) is prepared as reported
previously.[14, 16] A thin Si film (~ 300 nm) is then transfer-printed onto the 3D PhC
surface, followed by the printing of a top 3D silicon PhC that has 9 layers (4 μm) and is
lifted off from a sacrificial substrate. This method offers great flexibility in material
selections and structural designs. For instance, the Si thin film can be replaced by a III-V
compound semiconductor quantum-well layer or a polymer film embedded with colloidal
quantum dots, as to potentially achieve a hybrid lasing system.[17] The cavity layer can
also be patterned in advance of printing, in order to alter the modal field or enhance Q-
factor.[18, 19]
In this work, the transfer printing is performed using micro-structured PDMS
stamps as first developed by the Rogers group at Illinois.[20] Due to the viscoelastic nature
of the PDMS surface, objects (ink) can be detached from the donor substrate by rapidly
retrieving the stamp and transferred onto a receiver substrate by slow retraction.[21] The
micro-tips on the stamp can dynamically switch the adhesion between the ink and the
stamp. During the inking step, the applied force collapses the micro-tips, leading to a
conformal contact and thus maximal interfacial adhesion. After retrieval, the micro-tips
undergo elastic relaxation and gradually separate the ink from the backing layer, which
reduces the adhesion and facilitates the release of the ink.
Figur
conce
requi
we in
in co
schem
therm
by C
(NOA
50/50
are re
subst
re 3.3 Schem
Synthetic
entration ch
ired between
ntroduce a p
ontact with t
matically illu
mal oxide wa
VD and sub
A). The back
0 ethanol/ w
emoved simu
trate, and the
matic illustra
opal PhCs
hange during
n the cavity l
rocedure to
the wafer su
ustrates the p
afer with 1 μ
bsequently g
k surface of
ater), during
ultaneously.
e overview o
ations of fabr
often exhibi
g self-assem
layer and th
utilize the b
ubstrate and
process, star
μm thick SiO
glued onto a
the Si inver
g which both
Figure 3.5a
of its back su
53
rication proc
it terrace-lik
mbly. Howev
he PhC mirro
back surface
d thus posses
rting with th
O2 layer. Th
a sapphire w
rse opal is ex
h SiO2 spher
a shows the
urface is disp
cedures for 3
ke surfaces,
ver, opticall
ors for achie
e of Si invers
sses atomic
he colloidal g
he opal is co
wafer using N
xposed after
res and therm
Si inverse st
played in Fig
3D Si PhC m
as a result o
ly smooth i
eving high Q
se opals that
smoothness
growth of S
onformally c
Norland opt
r HF etching
mal oxide sa
tructure rele
gure 3.5b.
microcavity.
of the colloi
interfaces ar
Q-factor. Her
t is originall
s. Figure 3.
iO2 opal on
oated with S
tical adhesiv
g (10% HF i
acrificial laye
ased from th
id
re
re
ly
.4
a
Si
ve
in
er
he
Figur
Figursubst
X 10
help
receiv
highl
need
re 3.4 Schem
re 3.5 SEM trate (b) over
After flipp
0 μm X 0.3
of micro-st
ver substrat
ly porous (78
Up to her
to stack an
matic of proc
images of (arview of the
ping up the
μm) onto th
tructured sta
e, enabling
8%) Si inver
re, we have m
nother Si Ph
cedures to ex
a) cross-sectback surfac
back surface
he smooth 3
amps, trans
printing the
rse opal.
managed to
hC onto the
54
xpose the ba
tion view of e of the Si in
e of Si inver
3D PhC stru
fer-printing
e delicate S
place a Si c
e top of the
ack surface o
f the Si invernverse opal.
rse opals, w
ucture shown
virtually ex
i thin film
cavity layer o
e cavity lay
of Si inverse
rse opal relea
e print a Si
n in Figure
xerts little
without dam
on a 3D Si P
yer to finish
opals.
ased from th
film (100 μm
3.6. With th
force on th
mages onto
PhC and onl
h the device
he
m
he
he
a
ly
e.
Howe
appli
the fr
transf
Figurinver
prote
listed
ever, if the P
ed force oft
fragile 3D Ph
fer-printed w
re 3.6 SEMrse opal PhC
Figure 3.
cted and su
d below:
1. The S
thick
structu
PhC ink is p
en causes th
hC network
with minima
M image of a by transfer-
.7 shows th
spended by
Si-coated SiO
SiO2 layer. T
ure is pinche
prepared as
he structure
. Hence it r
l mechanica
a 320 nm thi-printing.
e process fl
dangling ph
O2 opal is de
The CVD co
ed off, to ens
55
those thin fi
to collapse
requires that
al impact.
ick Si film p
low for fabr
hotoresist an
eposited on
oated Si shou
sure the succ
film inks, e.g
during the p
t the 3D PhC
placed onto
ricating Si in
nchors. The
a thermal ox
uld be thick
cess of the w
g. the cavity
pick-up proc
C must be p
the back su
nverse opal
details for
xide substra
k enough suc
wet etching l
y layer,[22] th
cess owing t
protected an
urface of a S
inks that ar
each step ar
ate with 1 μm
ch that the 3D
ater.
he
to
nd
Si
re
re
m
D
56
2. AZ4620 photoresist (PR) is spin-casted on the colloidal film and patterned
into 160 μm X 160 μm squares. After development, the PR is hard-baked in
oven at 150 C for 30 min to increase its resistance to acids.
3. To etch Si and SiO2 at the exposed regions, Si RIE (10 SCCM O2, 10 SCCM
SF6, 50 mTorr, 70 W, 80 s) is first applied to remove the ~ 100 nm Si
overlayer, which reveals both SiO2 colloids and the underneath thermal oxide
sacrificial layer. SiO2 is then etched in 10% HF and 0.5% surfactant (3M
Novec 4200) solution, with the etching rates of 500 nm/min for SiO2 spheres
and 300 nm/min for thermal oxide. The left Si scaffold is subsequently etched
in 12 : 6 : 1 HNO3 (70%) : H2O : HF (49%) solution for 2 min.
4. After rinsing off the first PR pattern with acetone, another step of
photolithography with AZ4620 PR is performed to only cover the central
region of each square island with a 120 μm X 120 μm pattern. As a result, the
Si coating is partially exposed and then etched by RIE (20 SCCM O2, 20
SCCM CF4, 150 mTorr, 70 W, 80 s), followed by stripping off the PR pattern.
5. The third photolithography is carried out with AZ5214 PR on the square opal
islands to define the PR anchor that consist an 80 μm X 80 μm square and
four 5 μm X 40 μm stripes connecting the edge of the squares to the substrate.
Upon removal of all the remaining SiO2 (10% HF & 0.5% surfactant for 90
min), each PhC ink is released from the substrate but secured by the four PR
anchors.
Figurre 3.7 Schem
matic of fabrrication proc
57
cess for 3D SSi inverse oppal inks.
mech
inkin
adhes
objec
3.8a.
subst
Figur
unifo
Perot
witho
Figursubstrepre
The role o
hanical impa
ng step; (2) th
sion between
ct from the s
The PR res
trate at the P
re 3.8b disp
orm color an
t fringes in
out damages
re 3.8 Optictrate and (b)sents 100 μm
of PR ancho
act via defor
he flat squar
n the PDMS
substrate. Th
sidue of a d
PR anchors.
plays the bac
nd measured
Figure 3.9
and forms a
cal images o) the back sum.
ors is two-fol
rmation and
re PR other t
S stamp and
he optical im
elaminated i
Therefore th
ck surface o
d high reflec
also suppo
a tight contac
of (a) 3D Siurface of a
58
ld: (1) the da
thus protec
than the non
d the PhC in
mage of Si in
ink suggests
his procedur
of 3D Si PhC
ctance peak
ort the fact
ct with the r
i inverse opPhC ink pri
angling PR b
ct the fragile
n-planar opal
nk for determ
nverse opal
s that the in
re can keep
C printed on
k as well as
that the 3D
receiver subs
pal inks thatinted onto g
bonds absor
e 3D structu
l surface can
ministically d
inks is show
nk is dissocia
the PhC str
nto a glass s
the well-de
D PhC is tra
strate.
t are suspenglass substra
rb the most o
ure during th
n improve th
detaching th
wn in Figur
ated from th
ructure intac
substrate. Th
efined Fabry
ansfer-printe
nded from thate. Scale ba
of
he
he
he
re
he
ct.
he
y-
ed
he ar
59
Figure 3.9 Reflectance measurement of the 3D Si PhC ink printed on glass substrate.
We finally assembled a 3D PhC vertical microcavity by printing a 320 nm Si film
onto a reverted Si inverse opal (7 μm) and subsequently placing a Si invers opal onto the
cavity layer, as shown in Figure 3.10. If the Si cavity layer is replaced by an emitting
layer with high quantum yield in future, strong light-mater interaction can be expected in
such system, due to the strong photonic strength from high dielectric contrast 3D PhCs.
1000 1200 1400 1600 18000.0
0.2
0.4
0.6
0.8
1.0
Re
flect
an
ce
Wavelength (nm)
Figurassem
3.4 C
PhCs
corre
advan
sandw
speci
contr
re 3.10 (a) Smbled by tran
Conclusions
A new typ
s and Si th
sponding ca
nced transfe
wich structu
ifically patte
rol their emis
Schematic annsfer-printin
pe of vertica
in films. Su
avity modes
fer-printing
ure is fabrica
rn cavity lay
ssion proper
nd (b) opticang.
al microcavit
uch structur
and Q-facto
with micro
ated indepen
yer to achiev
rties.
60
al image illus
ty is develop
re is first m
r. The 3D P
o-structured
ndently, this
ve certain mo
strating a 3D
ped by comb
modeled in
PhC microcav
stamps. Si
s design can
odal profile
D Si PhC mic
bining 3D Si
FDTD to
vity is then
ince each
n potentially
or incorpora
crocavity
i inverse opa
calculate th
assembled b
layer of th
y allow us t
ate emitters t
al
he
by
he
to
to
61
3.5 References
[1] G. J. Fochesatto, Choice: Current Reviews for Academic Libraries 2009, 47, 724. [2] K. A. Arpin, A. Mihi, H. T. Johnson, A. J. Baca, J. A. Rogers, J. A. Lewis, P. V. Braun, Adv Mater 2010, 22, 1084. [3] G. S. Solomon, M. Pelton, Y. Yamamoto, Phys. Rev. Lett. 2001, 86, 3903. [4] E. Yablonovitch, Phys. Rev. Lett. 1987, 58, 2059. [5] S. Noda, M. Fujita, T. Asano, Nat Photon 2007, 1, 449. [6] K. J. Vahala, Nature 2003, 424, 839. [7] A. Tandaechanurat, S. Ishida, D. Guimard, M. Nomura, S. Iwamoto, Y. Arakawa, Nat Photon 2011, 5, 91. [8] S. Furumi, H. Fudouzi, H. T. Miyazaki, Y. Sakka, Adv Mater 2007, 19, 2067. [9] F. Jin, Y. Song, X. Z. Dong, W. Q. Chen, X. M. Duan, App Phys Lett 2007, 91. [10] H.-G. Park, S.-H. Kim, S.-H. Kwon, Y.-G. Ju, J.-K. Yang, J.-H. Baek, S.-B. Kim, Y.-H. Lee, Science 2004, 305, 1444. [11] S. Strauf, Nat Photon 2010, 4, 132. [12] R. A. Morgan, Photon Spectra 1990, 24, 89. [13] A. M. Kasten, J. D. Sulkin, P. O. Leisher, D. K. McElfresh, D. Vacar, K. D. Choquette, Ieee J Sel Top Quant 2008, 14, 1123. [14] H. L. Ning, A. Mihi, J. B. Geddes, M. Miyake, P. V. Braun, Adv Mater 2012, 24, Op153. [15] E. Palacios-Lidón, A. Blanco, M. Ibisate, F. Meseguer, C. López, J. Sánchez-Dehesa, Appl. Phys. Lett. 2002, 81, 4925. [16] A. Mihi, C. J. Zhang, P. V. Braun, Angew Chem Int Edit 2011, 50, 5711. [17] H. J. Yang, D. Y. Zhao, S. Chuwongin, J. H. Seo, W. Q. Yang, Y. C. Shuai, J. Berggren, M. Hammar, Z. Q. Ma, W. D. Zhou, Nat Photon 2012, 6, 615. [18] L. L. Tang, T. Yoshie, Opt Lett 2010, 35, 3144. [19] M. Dems, I. S. Chung, N. Peter, S. Bischoff, K. Panajotov, Opt Exp2010, 18, 16042. [20] S. Kim, J. A. Wu, A. Carlson, S. H. Jin, A. Kovalsky, P. Glass, Z. J. Liu, N. Ahmed, S. L. Elgan, W. Q. Chen, P. M. Ferreira, M. Sitti, Y. G. Huang, J. A. Rogers, P Natl Acad Sci USA 2010, 107, 17095. [21] M. A. Meitl, Z. T. Zhu, V. Kumar, K. J. Lee, X. Feng, Y. Y. Huang, I. Adesida, R. G. Nuzzo, J. A. Rogers, Nat Mater 2006, 5, 33. [22] H. Keum, A. Carlson, H. L. Ning, A. Mihi, J. D. Eisenhaure, P. V. Braun, J. A. Rogers, S. Kim, J Micromech Microeng 2012, 22.
62
CHAPTER FOUR
INCORPORATION OF FUNCTIONAL DEFECTS INTO
3D HOLOGRAPHIC PHOTONIC CRYSTALS
4.1 Introduction & motivation
The previous projects in Chapter 2 and Chapter 3 are built on 3D colloidal
photonic crystals (PhCs). Although this type of PhCs is known for their low cost, large
area and ease of fabrication, their structures often possess abundant undesired disorders
and defects that degrade their optical properties. Alternatively, holographic PhCs,
enabled by multi-beam interference inside a photoresist, can produce large-area and
defect-free 3D structures. So far, holographic PhCs with various symmetries have been
demonstrated such as diamond, woodpile, and face centered cubic.[1-3] However, one of
the remaining challenges for this technology is to furnish the 3D structure with needed
functionalities such as point defects for no threshold lasers or cavity QED,[4, 5] line
defects for waveguiding and plane defects for sensing.[6, 7] To date, the only available tool
to introduce defects into holographic PhCs is two-photon polymerization (TPP) direct
writing.[8] When exposed to high-intensity pulsed laser with half the absorption photon
energy, photoresists can be polymerized by absorbing two photons simultaneously. Direct
writing of defects inside holographic PhCs has been realized by coupling such TPP with
confocal microscopes and high-precision translation stages.[9, 10] Registration of line
defects with the crystal lattice has also been demonstrated using advanced fluorescent
imaging.[11] Albeit the TPP direct writing has an excellent flexibility of creating and
positioning arbitrary defects in PhCs, it suffers from Abbe’s diffraction limit and thus
63
exhibits a poor control of the size and shape of the introduced defects. Such structural
factors, however, strongly dictates their optical properties including defect mode, quality
factor and field distribution. In addition, imperfections in laser beam and scattering can
also increase the surface roughness of these defects, which is detrimental to their optical
performance. Another major drawback of the TPP is that it cannot write in media other
than photoresist or incorporate foreign functional materials such as emitters or absorbers,
significantly limiting the potentials of the 3D architectures created by this technique.
Here we develop a new route to introduce defects into holographic PhCs by transfer-
printing. This technique allows us to add a great variety of functionalities that are pre-
defined with high structural quality via fabrication tools like photo lithography, colloidal
synthesis and so on. By placing colloidal quantum dots (QDs) at specific locations inside
3D holographic PhCs, we demonstrate an excellent control of their spontaneous emission
(SE). In particular, we observe both suppressed emission due to photonic band gap and
enhanced emission from a PhC microcavity by altering the photoresist structure around
the QDs.
4.2 Experimental procedures for embedding defects
Figure 4.1 illustrates the fabrication procedures for incorporating nano/micro
defects into 3D holographic PhCs. The process begins with preparing an SU8 photoresist
film on a glass/ITO substrate, followed by forming a defect layer on a PDMS stamp. If
the defect is spin-casted onto the stamp, an oxygen-plasma treatment is needed in
advance to improve the hydrophilicity of the PDMS surface. Upon heating, the particle
layer is transfer-printed onto the bottom SU8 film. The top SU8 film is casted onto
64
another oxygen-plasma-treated stamp and subsequently printed onto the defect layer.
Finally, the entire sandwich structure undergoes the standard process for holographic
lithography including multi-beam interference exposure, post baking, and
development.[12] A detailed fabrication procedure is summarized as follows,
1. Spin-coating the bottom SU8 film (SU8 2010) at 2000 rpm for 30 s and pre-
baking at 65 C for 5 min and 95 C for 5 min;
2. Treating PDMS with oxygen plasma (600 mTorr, 50 W, 20 SCCM, 80 s);
3. Spin-coating nanoparticle solution (1.5 wt% LaF3 or 0.1 wt% 500 nm
polystyrene or silica spheres in IPA) on the treated PDMS;
4. Placing particle-coated PDMS in contact with SU8 film, followed by baking
at 65 C for 20 min and slowly peeling off PDMS stamp;
5. Preparing the top SU8 film on another treated PDMS using the same spin-
coating and pre-baking recipe as the bottom.
6. Transfer-printing SU8 onto the defect layer;
7. Pre-baking the sandwich structure at 65 C for 5 min and 95 C for 5 min;
8. Holographic exposure at 532 nm with a dose of 50 J/cm2 (0.3 s).
9. Post-baking at 85 C for 20 min.
10. Developing in PGMEA for 3 hours.
FigurPhCs
partic
beam
whos
refrac
the s
embe
are la
such
can b
Figur
electr
nanop
radic
Point
re 4.1 Schems.
The defec
cles. The on
ms during th
se refractive
ctive index
ize of nanom
edded in a ho
attice-registe
sandwich st
be inverted
re 4.2b, we
roplating an
particle lay
als in RIE e
t defects are
matic of proc
ct layer can
nly requirem
he holograph
indices are
difference is
metric scatte
olographic P
ered and not
tructure is fa
with high
successfully
nd SU8 dry
er allows b
etching to g
also demon
cedures to in
n exist in a
ment is that t
hic exposur
e close to th
s significant
ers). Figure
PhC as a pla
t disturbed b
abricated on
dielectric m
y inverted a
removal pr
both the el
go through,
nstrated by in
65
ntroduce pre
variety of
their presenc
re, which ca
hat of SU8,
t (because th
e 4.2a show
anar defect, w
by the prese
a transparen
material via
sandwich Ph
rocedures re
ectrolyte du
resulting in
ntroducing c
e-defined def
forms such
ce does not
an be satisf
or reducing
he intensity
ws a thin lay
where the to
ence of the n
nt ITO subst
electrodepo
hC with Cu2
eported prev
uring electr
n a uniforml
colloidal sph
fects into 3D
as films an
scatter the i
fied by ado
g their dime
of scatterin
yer of LaF3
op and botto
nanoparticle
trate, the pol
osition. As
2O by follow
viously.[2, 12]
rodeposition
ly inverted
heres into Ph
D holographi
nd individua
incident lase
pting object
ensions if th
ng scales wit
nanoparticle
om SU8 PhC
e layer. Whe
lymeric PhC
displayed i
wing the Cu2O
] The porou
and plasm
3D structure
hC structure
ic
al
er
ts
he
th
es
Cs
en
Cs
in
O
us
ma
e.
s.
Depe
4.3a)
FigurPhCs
FigurSU8
ending on th
) or solid def
re 4.2 SEMs and (b) Cu2
re 4.3 SEM holographic
he solubility
fects (Figure
M images of L2O PhCs inv
images of (a PhCs.
y of the sph
e 4.3b) can b
LaF3 nanopaverted from S
a) a polystyr
66
heres in the
be created us
article layerSU8 template
rene sphere
SU8 develo
sing polystyr
r sandwichedes.
and (b) a sil
oper, air def
rene or SiO2
d in (a) SU8
lica sphere s
fects (Figur
2 spheres.
8 holographi
andwiched i
re
ic
in
67
4.3 Light-matter interaction between introduced defects & their hosts
Light-matter interaction has hardly been investigated in 3D holographic PhCs since their
invention.[13] However, there have been extensive efforts focused on the PhCs realized by
other methods, such as colloidal self-assembled opals[14], top-down etched 3D
architectures[15] and so on. Two strategies are often used to introduce emitters into those
PhCs: (1) 3D PhCs are immersed in organic dye or colloidal QD solutions, and therefore
the entire structures are infiltrated with those emitters;[16-21] (2) emitters such as rare earth
ions are directly doped into 3D scaffolds during template-assisted material
conversions.[22] However, accurate studies on the SE manipulation require that the
emitters be placed at well-defined locations in 3D PhCs. Both methods lead to a 3D PhC
containing non-localized emitters that experience varying photonic environments at the
surface and inside the crystal. In comparison, our technique can offer an excellent spatial
control of the defect, allowing us to selectively place colloidal PbS QDs inside the PhC.
The incorporation of the QD emitter follows the same procedure as the particle defect,
except that the particle layer is replaced by a QD-SU8 composite film that is obtained via
dispersing QDs (10 mg/ml in hexane) in dilute SU8 solution (14%). By altering the
exposure type of this additive composite layer, we are able to convert it to either PhC
lattice or a solid plane cavity. In each case, the emission of the embedded QDs is
influenced by different photonic density of states (pDOS).
4.3.1 Control of emission by photonic band gap
The QD-SU8 mixture has a 0.2 wt% solid content of PbS QDs, exhibiting a very
low absorption at 532 nm. Such composite is spin-casted into a ~800 nm thick film and
68
subsequently sandwiched between two 10 μm thick SU8 layers. As illustrated in Figure
4.4a, after holographic exposure and development, the QDs are only confined in the
middle crystal plane. Figure 4.4b presents a cleaved cross section of this structure. The
highly uniform PhC lattice suggests the presence of QDs exerts no disturbance on either
photoresist chemistry or holographic exposure. Reflectance and photoluminescence (PL)
measurements are performed to characterize the optical properties of the QD-doped
holographic PhCs. A 4 X objective with a numerical aperture of 0.1 is used for excitation
as well as data collection in both measurements. Figure 4.5 shows the reflectance
spectrum taken from a holographic PhC with quantum dots embedded in the middle
lattice, where a nearly 80% reflectance peak arises at 1400 nm. The strong reflectance
peak, also called Bragg peak, is collectively contributed by crystal layers that are above
and below the QD-doped region and both have the same stop gap position. Therefore, the
embedded QD emitters, tough introduced extrinsically with respect to the PhC, do not
create any defect states and consequentially should interact directly with the modified
pDOS by the 3D PhC host.
Figurembe
Figurthe m
re 4.4 (a) Sedded with q
re 4.5 Reflemiddle lattice
100.0
0.2
0.4
0.6
0.8
1.0
Ref
lect
ance
Schematic aquantum dots
ction spectrue.
000 12
and (b) cross in the midd
um of a 3D h
200 1Wa
69
s-section SEdle lattice.
holographic
400 1avelength
EM image o
PhC with qu
600 1(nm)
of 3D holog
uantum dots
1800
graphic PhC
s embedded i
2000
Cs
in
70
4.3.2 Control of emission by 3D photonic crystal microcavity
Introducing defects into 3D holographic PhCs by transfer-printing also allow us to
take the advantage of other established fabrication technologies such as conventional UV
photolithography, to create reliable optical functionalities in a massively parallel manner.
To illustrate this capability, we expose the QD-SU8 composite layer before printing with
a photomask that consists of 200 μm 400 μm rectangle arrays. As illustrated in Figure
4.6a, the exposed SU8 defect layer crosslinks during the post-bake and results in a solid
SU8 film containing QDs. The reason that we pattern the defect layer into large islands is
to provide an effective pathway for the developer to reach and develop the bottom PhCs,
leading to a well-defined planar defect sandwiched between 3D PhC lattices (Figure
4.6b). The homogenous QD-SU8 defect layer (within the rectangle region) inside the 3D
PhC breaks the lattice periodicity and forms a vertical PhC microcavity, which
consequentially opens allowed states in the forbidden photonic band. Figure 4.7 shows
the reflectance spectrum from such structure. The dip at 1350 nm inside the Bragg peak is
the cavity resonance mode associated with the defect layer. According to the so-called
Purcell effect,[23] the spontaneous emission of an emitter will be enhanced if such emitter
is placed inside a microcavity and emit near the cavity mode.
Figurvertic
Figurthat c
re 4.6 (a) Sccal cavity wi
re 4.7 Reflecontains quan
100.0
0.2
0.4
0.6
0.8
1.0
Ref
lect
ance
chematic andith quantum
ectance specntum dots.
000 12
d (b) SEM idots embedd
trum of a ho
200 14
Wa
71
image of theded in the de
olographic P
400 1
avelength
e fractured crefect layer.
PhC with an
1600 1
(nm)
ross-section
n introduced
1800
n of a 3D Ph
planar defec
2000
C
ct
72
The approach presented here can be extended to incorporations of arbitrary
defects into 3D PhC building blocks. For instance, photonic circuits including emitters,
microcavities, waveguides and detectors can be defined by standard micro/nano
fabrication technologies prior to transfer-printing. In addition, alternately stacking
multilayers of PhCs and defects can also be realized by repeatedly printing, to achieve
more complex functions and more efficiently utilize the 3D space. Although the
refractive index of the introduced defects is restricted for minimizing the scattering loss
during the holographic exposure, there are a broad variety of photonic materials
compatible with this technique such as most of organic polymers, inorganic oxides,
nanocrystals, and porous media that possess a tunable refractive index based on its
porosity.[24] If the transfer-printing is precisely aligned with the incident laser beams,
where for the phase-mask-based holographic lithography [25] the transferred objects only
need to register with respect to the phase mask, the defects can be incorporated with
registration to the crystal lattice. In addition, by inverting the photoresist PhCs with high
dielectric materials to increase their photonic strength, we can anticipate stronger light-
matter interactions at the confined defects.
4.4 Conclusions
In summary, we have demonstrated that transfer-printing provides an excellent
means to embed a broad variety of artificial defects into holographic PhCs. The
introduced defects can either remain unchanged or be dissolved during the process of
holographic lithography, leading to either dielectric or air defects. We have also placed
colloidal quantum dots at a specific location inside 3D PhCs. By changing their
73
surrounding photonic environments, we observe both suppressed and enhanced emission
from the localized emitters. Our method provides a robust capability for adding
functionalities into 3D PhCs, opening the door for realizing integrated photonic circuits
in 3D holographic structures.
4.5 References
[1] G. Q. Liang, X. L. Zhu, Y. G. Xu, J. Li, S. Yang, Adv Mater 2010, 22, 4524. [2] S. G. Park, M. Miyake, S. M. Yang, P. V. Braun, Adv Mater 2011, 23, 2749. [3] Y. C. Chen, J. B. Geddes, J. T. Lee, P. V. Braun, P. Wiltzius, App Phys Let 2007, 91. [4] A. Tandaechanurat, S. Ishida, D. Guimard, M. Nomura, S. Iwamoto, Y. Arakawa, Nat Photon 2011, 5, 91. [5] G. Khitrova, H. M. Gibbs, M. Kira, S. W. Koch, A. Scherer, Nat Phys 2006, 2, 81. [6] S. A. Rinne, F. Garcia-Santamaria, P. V. Braun, Nat Photon 2008, 2, 52. [7] P. V. Braun, S. A. Rinne, F. Garcia-Santamaria, Adv Mater 2006, 18, 2665. [8] M. Deubel, G. Von Freymann, M. Wegener, S. Pereira, K. Busch, C. M. Soukoulis, Nat Mater 2004, 3, 444. [9] V. Ramanan, E. Nelson, A. Brzezinski, P. V. Braun, P. Wiltzius, App Phys Let2008, 92, 173304. [10] E. C. Nelson, F. Garcia-Santamaria, P. V. Braun, Adv Fun Mater 2008, 18, 1983. [11] J. Scrimgeour, D. N. Sharp, C. F. Blanford, O. M. Roche, R. G. Denning, A. J. Turberfield, Adv Mater 2006, 18, 1557. [12] M. Miyake, Y. C. Chen, P. V. Braun, P. Wiltzius, Adv Mater 2009, 21, 3012. [13] M. Campbell, D. N. Sharp, M. T. Harrison, R. G. Denning, A. J. Turberfield, Nature 2000, 404, 53. [14] J. F. Galisteo-Lopez, M. Ibisate, R. Sapienza, L. S. Froufe-Perez, A. Blanco, C. Lopez, Adv Mater 2011, 23, 30. [15] J. M. van den Broek, L. A. Woldering, R. W. Tjerkstra, F. B. Segerink, I. D. Setija, W. L. Vos, Adv Fun Mater 2012, 22, 25. [16] P. Lodahl, A. Floris van Driel, I. S. Nikolaev, A. Irman, K. Overgaag, D. Vanmaekelbergh, W. L. Vos, Nature 2004, 430, 654. [17] L.-L. L. Zhi-Yuan Li, Z.-Q. Zhang, Phys Rev Lett 2000, 84. [18] N. Vats, S. John, K. Busch, Phys Rev A 2002, 65, 043808. [19] I. S. Nikolaev, P. Lodahl, A. F. van Driel, A. F. Koenderink, W. L. Vos, Phys Rev B 2007, 75, 115302. [20] M. R. Jorgensen, J. W. Galusha, M. H. Bartl, Phys Rev Lett 2011, 107, 143902. [21] H. Yamada, T. Nakamura, Y. Yamada, K. Yano, Adv Mater 2009, 21, 4134. [22] E. Bovero, F. C. Van Veggel, J Am Chem Soc 2008, 130, 15374. [23] E. M. Purcell, Phys Rev 1946, 69, 681.
74
[24] G. G. Qin, Y. J. Li, Phys Rev B 2003, 68. [25] S.Jeon,J.U.Park,R.Cirelli,S.Yang,C.E.Heitzman,P.V.Braun,P.J.A.Kenis,J.A.Rogers,PNatlAcadSciUSA2004,101,12428.
75
CHAPTER FIVE
ASSEMBLY OF TUNABLE POROUS SILICON MICROCAVITY
5.1 Introduction & motivation
Excitement over the use of porous silicon (PSi) in the field of optoelectronics was
first generated over twenty years ago with the observation of the material’s visible
photoluminescence at room temperature.[1] The years following this discovery were filled
with efforts to realize efficient PSi light-emitting devices,[2, 3] but achieving practical
electroluminescence efficiency and stability has proven much more difficult than
anticipated and,[4, 5] as such, has greatly shifted the application focus. Even so, the
distinction of versatile optical material is still fitting for PSi, which is formed by
electrochemically etching silicon (Si) in a hydrofluoric acid-based electrolyte, with the
resultant porosity (i.e. void fraction) determined by the applied current density. Part of
PSi’s versatility is in the ability to obtain free-standing films by performing an
electropolishing step that allows the detached film to be transferred to a new substrate.
Additionally, a great strength of this porous material system is in the inherent ability to
modulate the refractive index in two distinct manners—porosity variation and porosity
infiltration—which has allowed PSi to make its mark in sensing applications.[6] In
particular, porosity variations induced by time-varying etching currents enable the
formation of high quality superlattices with pronounced optical signatures, such as
vertical microcavities with sharp optical resonances,[7] that can be shifted due to the
infiltration of the porous matrix with foreign media.[8, 9] Although these PSi vertical
microcavities have been exploited more for sensing purposes, the fact remains that they
76
possess the ability to manipulate the emission of highly efficient light-emitting entities
that spatially and spectrally overlap with the cavity.
In spite of the realization of high quality monolithic vertical cavities,[10] the
overall effort to demonstrate the coupling of emitters with PSi-based cavities has been
unconvincing due in large part to the limited class of emitters that can be integrated
inside the cavity and the difficulty in handling the fragile PSi films. As a result, most
efforts have strictly relied on emitters that can either be embedded into the porous
structure[11-13] or have been pre-embedded into the starting Si wafer used to fabricate the
PSi.[14] While a hybrid vertical cavity utilizing PSi distributed Bragg reflectors (DBR)
and a polymer defect layer has been considered,[15] the quality factor (Q-factor) of the
resulting structure is too poor for photonic applications. More eloquent methods of
assembling PSi photonic devices have been proposed, such as dry-removal lithography[16]
and a biofunctionalization-driven self-assembly.[17] But, these techniques have been
geared more towards the formation of PSi-based sensing arrays that lack the optical
quality desired for emission modification applications.[18, 19] Recently approaches based
on transfer-printing have realized a broad variety of heterogeneously integrated
optoelectronic and photonic systems.[20-23] In this work, we demonstrate that a modified
transfer-printing technique can enable the formation of high-quality PSi hybrid vertical
microcavities that can incorporate different types of external emitters, such as dyes,
quantum dots and solid-state thin films. Further, we utilize both routes to index
modulation to show that PSi offers the ability to tune the external emitter in a fashion that
is both coarse and fine in nature.
5.2 P
that b
(poyd
coval
film
modi
appli
oxyg
PDM
visco
proce
film
comp
Figursilani
Printing hyb
Freshly et
bears a posi
dimethalsilo
lent bonds,[2
from PDMS
fy the PDM
ed to the m
en-plasma a
MS and PSi is
oelasticity to
esses.[21] As
after relea
pletely print
re 5.1 Opticized PDMS
brid porous
tched PSi po
tive charge
xane) stamp
24, 25] makes
S. To allevi
MS surface
master in so
and then exp
s significantl
kinetically c
a result, we
sing it from
it onto a cur
cal image ofstamp.
silicon micr
ossesses a h
and results
p. This, alo
s it difficult
iate the inte
following a
oft lithograp
posing it to a
ly reduced. T
control the s
e can succes
m the subs
red SU8 surf
f PSi DBR re
77
rocavity
highly hydro
in a strong
ong with th
to achieve
raction betw
a standard si
phy.[24, 26] By
a fluorinated
The treated s
separation en
ssfully retrie
strate via e
face.
eleased from
phobic, hyd
electrostatic
e irreversib
damage-fre
ween PSi an
ilanization p
y activating
d silane age
stamp, howe
nergy during
eve the entir
electro-polis
m silicon sub
drogen-termi
c attraction t
ble formation
e detachmen
nd the PDM
procedure th
g the PDMS
nt, the adhe
ever, still off
g the pick-up
re free-stand
shing (Figu
bstrate and r
inated surfac
to the PDM
n of Si-O-S
nt of the PS
MS stamp, w
hat is widel
S surface vi
esion betwee
fers sufficien
p and printin
ding PSi DBR
ure 5.1) an
retrieved by
ce
MS
Si
Si
we
ly
ia
en
nt
ng
R
nd
a
78
The assembly of polymer-PSi hybrid microcavity begins with printing a ~ 500 nm
cured SU8 film onto a PSi DBR that consists of 15 pairs of alternating high (~2.6) and
low (~1.7) refractive index layers (Figure 5.2a). Next, another PSi DBR with the same
index contrast—but only 11 lattice periods—is transfer-printed onto the ½ λ SU8 cavity
layer (Figure 5.2b). Figure 5.3 displays the cross section of such a hybrid microcavity,
showing that the printed SU8 layer forms smooth and distinct interfaces with PSi. The
optical properties of a hybrid microcavity, as well as a monolithic PSi microcavity
obtained by a single etching procedure, are characterized by measuring the reflectance
using a 4X objective with a 0.1 numerical aperture. As shown in Figure 5.4a, the cavity
modes of these two types of microcavities both appear as sharp dips around the center
(1500 nm) of the 300 nm broad DBR stop band and yield a similar Q-factor of ~750. This
value, though two or three times larger than those reported previously for hybrid
systems,[13] is significantly smaller than the Q-factor (~7000) of a monolithic structure
measured with an very small spot size and numerical aperture,[27] suggesting that the our
observed Q factor of our hybrid microcavity could potentially be limited by the
reflectivity measurement. The cavity mode across the entire printed SU8 cavity is found
to be 1497.6 ± 6.8 nm (Figure 5.4b), where the small spectral deviation, though likely
caused by the thickness variation of the SU8 cavity, demonstrates the capacity of our
method for assembling large-area, high-quality microcavities.
Figurhybri
Figurclose
re 5.2 Opticid microcavi
re 5.3 Crosr view of SU
cal images oity.
ss-section SEU8 cavity lay
of (a) bottom
EM image oyer.
79
m PSi DBR
of PSi-SU8
printed with
hybrid mic
h SU8 and (
crocavity. In
(b) assemble
nset image:
ed
a
80
Figure 5.4 (a) Reflectance spectrum of PSi-SU8 hybrid microcavity. (b) Spatial cavity mode distribution across the sample.
1000 1200 1400 1600 1800 2000 22000.0
0.2
0.4
0.6
0.8
1.0
Ref
lect
ance
Wavelength (nm)
0 1 2 3 4 5 6 7 8 91320
1360
1400
1440
1480
1520
1560
1600
1640
Cav
ity m
ode
(nm
)
Position (mm)
(a)
(b)
81
5.3 Coarse tuning of microcavity resonance
The ability to form PSi-polymer hybrid microcavities affords the opportunity to
control the emission of a variety of emitters that can be dispersed in the polymer matrix,
such as organic dye molecules, colloidal quantum dots, and rare earth nanocrystals,
among others. In order to utilize such a non-porous polymeric cavity, while still
maintaining the aforementioned porosity-based refractive index modulation, we introduce
an additional PSi layer inside the cavity that contacts the polymer layer, shown
schematically in Figure 5.5. This layer, known as the cavity coupling layer (CCL),
couples with the solid polymer layer to concurrently produce a resonant cavity mode
located spectrally at , where m is the order of the cavity mode
and is the optical thickness (the product of the refractive index and thickness) of the
ith layer. Consequently, the refractive index modulation capabilities of PSi are extended
to the cavity and pave a pathway to tune the cavity mode. Further, both routes to
refractive index modulation can be utilized in the CCL, with porosity variation
introducing a coarse tuning mechanism, while porosity infiltration provides a fine tuning.
The porosity variation coarse tuning is accomplished by forming a gradient
refractive index (GRIN) CCL that results in a large spatial variation of the cavity mode.
In order to achieve a GRIN CCL, it is necessary to induce a spatially-varying current
density during its formation, given that the resulting refractive index is determined by the
local current density.[28] A straightforward method to introduce a spatially-varying
current density is through the electrode.[29] During the formation of the PSi DBR, we
utilize a Pt ring electrode (5 mm diameter) that resides ~25mm from the sample, giving a
very uniform current density distribution. However, for the GRIN CCL, we use a Pt wire
cente
curre
the i
Simil
reson
Figursymm
FigurMeas
ered ~1mm a
nt density. T
interference
larly, the GR
nance observ
re 5.5 Schemetric cavity
re 5.6 (a) Osured spatial
above the s
This variatio
fringes fro
RIN nature
ved across th
ematic illusty, n1= n3 and
Optical ima variation of
ample, whic
on is apparen
m the CCL
of the CCL
he assembled
tration of PSd1= d3.
age of a GRf the cavity m
82
ch gives rise
nt prior to P
L atop the
L is evident
d PSi microc
Si cavity co
RIN CCL lmode from a
n2 d2
n1 d1
n3 d3
e to a strong
PSi vertical
PSi DBR,
by the larg
cavity in Fig
oupling laye
ayer etcheda GRIN micr
g spatial va
microcavity
shown in
ge variation
ure 5.6b.
ers, n1d1 and
d on top of rocavity.
ariation in th
y assembly b
Figure 5.6a
of the cavit
d n3d3. For
f a DBR. (b
he
by
a.
ty
a
b)
83
5.4 Fine tuning of microcavity resonance
The addition of a CCL can also provide an excellent means to finely tune the
resonant mode of the pre-fabricated hybrid cavity. This can be done by gradually
infiltrating the porous structure with a nanometer-scale, high aspect-ratio deposition tool
such as atomic layer deposition (ALD). We demonstrate ALD-based tuning by
constructing a 2 λ microcavity consisting of two CCL layers (450 nm optical thickness
each) and a SU8 layer doped with PbS quantum dots (QDs) (Figure 5.7a). The
microcavity resonance strongly influences the photoluminescence (PL) of the embedded
QDs, leading to a significant redistribution of the emission spectrum in the normal
direction, as shown in Figure 5.7b. The original broad PbS QD emission (linewidth: ~
100 nm) is severely suppressed everywhere inside the DBR stopband expect at the cavity
mode. Before ALD, the measured PL position and linewidth from the hybrid microcavity
are 1146 nm and 1.9 nm, respectively, corresponding to a Q-factor of ~ 600. Al2O3 ALD
is subsequently applied to this structure from the top with a 1.2 Å per cycle deposition
rate. The conformal Al2O3 coating gradually increases the CCL optical thickness,
causing the emission peak to red-shift ~ 0.8 nm per cycle in the spectrum. The emission
peak eventually stops at 1163 nm after 20 cycles (Figure 5.8), which indicates that the
porous network has pinched off. The magnitude of this spectral shifting, though
ultimately restricted by pinch-off of the porous network, can be scaled by the refractive
index of the material introduced during the ALD process. For example, a larger tuning
range can be attained by increasing the volume fraction of CCL or infiltrating the CCL
with higher refractive index materials such as HfO2, TiO2, Si, etc. In particular, the total
spect
3.4 an
Figurand tspectrespe
5.5 In
emitt
Figur
layer
perio
interf
top P
emitt
tral shift can
nd the micro
re 5.7 (a) Stwo symmetrum of QDsectively.
ncorporatio
Our techn
ter - such as
re 5.9a illus
s (500 nm),
d PSi DBR
facial SU8 l
PSi DBR (F
ting layer du
n be extended
ocavity consi
Schematic ofetric CCLs s doped in S
on of solid st
nique has th
s highly effi
strates the st
a heterogen
R are seque
ayers not on
Figure 5.9b
ue to the hi
d to 86 nm if
ists of 90% C
f a hybrid mwith 450 n
SU8 film (bl
tate thin film
he ability to
icient group
tructural lay
neous GaAs
entially prin
nly ensure th
b), but they
igh refractiv
84
f the refracti
CCL.
microcavity cnm optical tlack curve)
m emitters
o incorporat
p III-V semi
yout of a 6λ
film (400 μm
nted onto a
he complete
also provid
ve index con
ive index of
consisting ofthickness. (band in a 2λ
te any arbitr
iconductors
hybrid micr
μm 400 μm
15-period b
e printings o
de extra opt
ntrast with r
f the infiltrat
f a QD-dopb) Normalizmicrocavity
rary solid-st
- into PSi m
rocavity, wh
m 1200 nm
bottom DB
f the GaAs
tical confin
respect to G
ed material
ed SU8 layezed emissioy (red curve
tate thin film
microcavitie
here two SU
m) and an 11
R. Here, th
layer and th
ement in th
GaAs. Figur
is
er on e),
m
s.
U8
1-
he
he
he
re
85
5.10a compares the emission spectra of the GaAs film before and after inclusion in a
microcavity. While the stand-alone GaAs material exhibits a relatively broad emission
with 30 nm full width half maximum (FWHM), the GaAs emission from the microcavity
is strongly modified and appears as a sharp peak at the position of the cavity mode. The
measured emission peak linewidth is 1 nm (Figure 5.10b), corresponding to a Q-factor of
~ 900.
Figure 5.8 Emission spectra of QDs from a 2λ hybrid microcavity for different ALD cycles.
Although promising hybrid light-emitting devices from Si and group III-V
semiconductors have been demonstrated,[22, 30-32] they primarily operate below the Si
bandgap ( > 1100 nm) to reduce the absorption loss from Si. Compared to Si, PSi exhibits
a much smaller absorption above the Si bandgap due to the reduced absorbing volume
1130 1140 1150 1160 1170 11800.0
0.2
0.4
0.6
0.8
1.0
1.2
No
rmal
ized
Int
ens
ity
Wavelegnth (nm)
NO ALD 10 cycles 20 cycles 30 cycles
and t
micro
nm fr
the sa
Figurheter
Figurembefrom
0.
0.
0.
0.
0.
1.
1.
Nor
mal
ized
Int
ensi
ty
(a)
the increase
ocavity prese
rom the top
ame thicknes
re 5.9 (a) Scojunction Ga
re 5.10 (a) edded in a 6the 6λ PSi m
800 825.0
.2
.4
.6
.8
.0
.2
ed effective
ents a ~ 900
11-pair DBR
ss (2.4 μm),
chematic andaAs emitter.
Emission spλ PSi micro
microcavity.
850 875
Wavelength (
electronic b
0 Q-factor at
R per round
~ 20% of th
d (b) optical
pectra of a hocavity (red)
5 900
nm)
GaAs bare film GaAs in PSI cav
1 nm
86
bandgap.[33]
t 900 nm sug
trip is less t
he light is ab
image of a 6
heterogeneo). (b) High r
925
vity
8900.0
0.2
0.4
0.6
0.8
1.0
Nor
mal
ize
d In
tens
ity
(b)
In particula
ggests that t
than 3%. Ho
sorbed per r
6λ hybrid m
us GaAs filresolution pl
0 892 894
W
ar, the fact
the absorptio
owever, for a
round trip.
microcavity th
lm (black) alot of the Ga
4 896 89
Wavelength (nm)
that our PS
on loss at 90
a Si film wit
hat contains
and such filmaAs emissio
98 900 90
)
Si
00
th
a
m on
02
87
5.6 Conclusions
We have demonstrated that a modified transfer-printing technique enables the
formation of high-quality, hybrid vertical microcavities that feature any arbitrary light-
emitting layer sandwiched between PSi DBRs. The resonant electromagnetic mode of
this cavity structure couples with the emission spectrum of the light-emitting entity to
severely suppresses emission everywhere in the stopband of the DBR, except at the
spectral position of the cavity mode, where an emission enhancement occurs. We
observed this redistributed emission separately for a PbS QD-doped polymer film and a
GaAs solid-state thin film in a microcavity configuration with PSi DBRs. Additionally,
we have shown that the addition of a PSi CCL extends the inherent index modulation
capabilities of PSi to the cavity. The PSi CCL provides a mechanism for tuning the
hybrid microcavity’s resonant cavity mode and emission spectrum in both coarse and fine
natures by way of intentional porosity variation and porosity infiltration, respectively. We
specifically demonstrated coarse tuning in the form of a spatially varying cavity
resonance through strategically forming a GRIN CCL. We used Al2O3 ALD to infiltrate
the CCL and display our fine tuning capabilities, which have a range that scales
according to the refractive index of the infiltrating material.
5.7 References
[1] L. T. Canham, App Phys Lett 1990, 57, 1046. [2] L. T. Canham, T. I. Cox, A. Loni, A. J. Simons, Appl Surf Sci 1996, 102, 436. [3] O. Bisi, S. Ossicini, L. Pavesi, Surf Sci Rep 2000, 38, 1. [4] G. Korotcenkov, B. K. Cho, Crit Rev Solid State 2010, 35, 153. [5] H. Foll, M. Christophersen, J. Carstensen, G. Hasse, Mat Sci Eng R 2002, 39, 93. [6] G. Korotcenkov, B. K. Cho, Crit Rev Solid State 2010, 35, 1. [7] P. J. Reece, G. Lerondel, J. Mulders, W. H. Zheng, M. Gal, Physica Status Solidi a-Applied Research 2003, 197, 321.
88
[8] V. Mulloni, L. Pavesi, App Phys Lett 2000, 76, 2523. [9] M. A. Anderson, A. Tinsley-Bown, P. Allcock, E. A. Perkins, P. Snow, M. Hollings, R. G. Smith, C. Reeves, D. J. Squirrell, S. Nicklin, T. I. Cox, Physica Status Solidi a-Applied Research 2003, 197, 528. [10] M. Ghulinyan, C. J. Oton, G. Bonetti, Z. Gaburro, L. Pavesi, Journal of Applied Physics 2003, 93, 9724. [11] A. Venturello, C. Ricciardi, F. GiorgiS, S. Strola, G. P. Salvador, E. Garrone, F. Geobaldo, J Non-Cryst Solids 2006, 352, 1230. [12] V. K. Dwivedi, K. Pradeesh, G. V. Prakash, Appl Surf Sci 2011, 257, 3468. [13] H. Qiao, B. Guan, T. Bocking, M. Gal, J. J. Gooding, P. J. Reece, App Phys Lett 2010, 96. [14] P. J. Reece, M. Gal, H. H. Tan, C. Jagadish, App Phys Lett 2004, 85, 3363. [15] F. Y. Sychev, I. E. Razdolski, T. V. Murzina, O. A. Aktsipetrov, T. Trifonov, S. Cheylan, App Phys Lett 2009, 95. [16] D. J. Sirbuly, G. M. Lowman, B. Scott, G. D. Stucky, S. K. Buratto, Adv Matter 2003, 15, 149. [17] T. Bocking, K. A. Kilian, P. J. Reece, K. Gaus, M. Gal, J. J. Gooding, Soft Matter 2012, 8, 360. [18] T. Bocking, K. A. Kilian, P. J. Reece, K. Gaus, M. Gal, J. J. Gooding, Acs Appl Mater Inter 2010, 2, 3270. [19] D. J. Gargas, O. Muresan, D. J. Sirbuly, S. K. Buratto, Adv Mater 2006, 18, 3164. [20] S. Kim, J. A. Wu, A. Carlson, S. H. Jin, A. Kovalsky, P. Glass, Z. J. Liu, N. Ahmed, S. L. Elgan, W. Q. Chen, P. M. Ferreira, M. Sitti, Y. G. Huang, J. A. Rogers, P Natl Acad Sci USA 2010, 107, 17095. [21] M. A. Meitl, Z. T. Zhu, V. Kumar, K. J. Lee, X. Feng, Y. Y. Huang, I. Adesida, R. G. Nuzzo, J. A. Rogers, Nat Matter 2006, 5, 33. [22] H. J. Yang, D. Y. Zhao, S. Chuwongin, J. H. Seo, W. Q. Yang, Y. C. Shuai, J. Berggren, M. Hammar, Z. Q. Ma, W. D. Zhou, Nat Photonics 2012, 6, 615. [23] J. Justice, C. Bower, M. Meitl, M. B. Mooney, M. A. Gubbins, B. Corbett, Nat Photonics 2012, 6, 610. [24] D. Qin, Y. N. Xia, G. M. Whitesides, Nat Protoc 2010, 5, 491. [25] I. Wong, C. M. Ho, Microfluid Nanofluid 2009, 7, 291. [26] Y. N. Xia, G. M. Whitesides, Annu Rev Mater Sci 1998, 28, 153. [27] M. Gal, P. J. Reece, W. H. Zheng, G. Lerondel, Photonics: Design, Technology, and Packaging 2004, 5277, 9. [28] S. Ilyas, M. Gal, 2006 Conference on Optoelectronic and Microelectronic Materials & Devices 2006, 245. [29] B. E. Collins, K. P. S. Dancil, G. Abbi, M. J. Sailor, Adv Fun Mater 2002, 12, 187. [30] A. Lee, H. Y. Liu, A. Seeds, Semicond Sci Tech 2013, 28. [31] H. Park, A. W. Fang, S. Kodama, J. E. Bowers, Opt Exp 2005, 13, 9460. [32] A. W. Fang, H. Park, R. Jones, O. Cohen, M. J. Paniccia, J. E. Bowers, Ieee Photonic Tech L 2006, 18, 1143. [33] S. Datta, K. L. Narasimhan, Phys Rev B 1999, 60, 8246.
89
CHAPTER SIX
HIGH POWER LITHIUM ION MICROBATTERY FROM
3D HOLOGRAPHIC LITHOGRAPHY
6.1 Introduction & motivation
3D bicontinuous porous electrodes can enable rapid charge and discharge for
lithium ion batteries because of their shortened pathways for both liquid-phase and solid-
phase ion diffusions.[1] Recently this type of electrode has been integrated in
interdigitated microbatteries (MBs),[2] exhibiting 2X greater energy density and 2000X
greater power density compared to previous works.[3-5] Such MBs were realized by
independently electroplating anode and cathode active materials on interdigitated arrays
of 3D porous nickel scaffolds. The Ni scaffold was originally electro-deposited on a
patterned gold substrate through colloidal crystals (opals), followed by removing the opal
template. Although this work has set a few new records for high power MBs, there are
several important issues to be addressed. (1) The Ni current collector grew isotropically
inside opal templates during the bottom-up deposition. This led to hemispherically
shaped electrodes that did not fully utilize the device volume, which consequentially hurt
their energy density. (2) MBs often require tall electrodes (~100 μm) to achieve high
areal energy density.[6] However, in the current fabrication scheme the electrode width
scaled with its height due to the isotropic growth. Raising electrode height caused its
width to increases simultaneously, which degraded their power performance and limited
their areal density. (3) The growth of self-assembled colloidal opals is usually defective
and also not compatible with microelectronic fabrications. (4) Probably the most
90
important issue is that the capacity of the battery decreased to 70% after 15 cycles.
However, the battery should cycle for at least hundreds of times above 80% for practical
applications. In this work, my goal is to address all the above problems using porous
electrodes fabricated by 3D holographic lithography.
6.2 Microbattery assembly
6.2.1 Fabrication of microbattery templates
The fabrication of the microbattery template involves both 3D holographic
lithography and conventional photolithography. The former produces periodic 3D
structures with submicron-sized features, and the latter defined the 2D interdigitated
electrode pattern. In principle, these two types of exposure can be performed on a single
photoresist film in sequence. However, when made using SU8 negative photoresist, the
3D structure suffers from volume shrinkage of ~ 40% during development, while the 2D
pattern only shrinks ~ 7%, resulting in a large distortion in the structure. For example, the
SU8 photoresist in Figure 6.1 is exposed in both manners and the distortion is clearly
visible at the interface between the 2D and 3D patterns. In contrast, the 3D structure from
positive-tone resist is immune to volume shrinkage. 2D and 3D photolithographies have
been combined to make complex patterns in DNQ-based positive resists - AZ9620.[7]
However, the DNQ-based photoresist often possesses a large absorption coefficient,
which restricts the thickness of the 3D structure to ~10 μm. To address all the problems
outlined in the introduction, I use SU8 resist to fabricate thick uniform 3D structures and
subsequently infiltrate the SU8 network with AZ9620 to pattern the interdigitated battery
electrode.
Figurhologinside
as rep
absor
temp
photo
in Fig
total
them
cross
the N
re 6.1 Crosgraphic phote the 3D pat
The 3D S
ported previ
rption at the
late to nicke
oresist AZ96
gure 6.2a. F
area is 4 mm
are 35 μm
-section view
Ni to only gro
s-section SEtolithographitern caused
SU8 structure
iously.[8, 9] T
laser wavel
el by electro
620, the batt
Figure 6.2b
m2. The wid
and 15 μm
w of such a
ow vertically
EM image oies in a singthe structure
e is created o
The ITO glas
length (532 n
odeposition.
tery electrod
is an optica
dth of the in
m, respective
structure, w
y in the next
91
of a SU8 stgle photoresie to distort n
on ITO glas
ss (ITO thick
nm), and als
After infiltra
de is defined
al image of t
ndividual ele
ely. The SE
where the stra
step.
tructure creaist film. Thenear 2D featu
s by 4-beam
kness: 40 nm
so allows inv
ating the SU
d photolithog
the resulting
ectrode finge
EM image in
aight positiv
ated by bothe larger voluures.
m interferenc
m) has low r
version of th
U8 structure
graphically,
g battery tem
ers and the g
n Figure 6.2
ve resist wall
h 2D and 3Dme shrinkag
ce lithograph
reflection an
he photoresi
with positiv
as illustrate
mplate, whos
gaps betwee
2c shows th
ls will restric
D ge
hy
nd
st
ve
ed
se
en
he
ct
FigurresistCross
6.2.2
few o
reach
this p
electr
treatm
hundr
for th
re 6.2 (a) St is infiltrates-section SE
Electrodepo
Ni can b
ohms to a fe
hing a few m
problem, the
rode for 5
ment, a thin
red microns
he significan
Schematic aned in 3D S
EM image of
osition of cu
e easily elec
few kilo ohm
microns thick
e ITO is elec
s in 0.01 M
dark layer a
on the oxid
nt adhesion
nd (b) opticSU8 structurf the patterne
rrent collect
ctroplated on
ms. Howeve
k due to the p
ctrochemical
M Na2SO4
appears on th
dized ITO wi
improvemen
92
cal image ofres and defied AZ9620 r
tors & active
n ITO substr
r, the depos
poor adhesio
lly oxidized
and 0.1 M
he ITO surfa
ithout adhes
nt is that N
f the batteryines the interesist in the S
e materials
rates with re
sited Ni film
on between N
at 2.98 V v
H2SO4 sol
ace, allowin
sion failure.
Ni forms stro
y template. erdigitated eSU8 network
esistances ra
m often dela
Ni and oxid
versus a plat
lution. After
ng Ni to grow
One possibl
ong chemica
The AZ962electrode. (ck.
anging from
aminates afte
de. To addres
tinum counte
r the surfac
w up to a few
le explanatio
al bonds wit
20 c)
a
er
ss
er
ce
w
on
th
indiu
only
remo
SCCM
ITO
(cond
insula
optic
walls
width
Figur(c) C
um and tin v
absorbed on
After Ni i
val with rea
M CF4, 60 m
conductive
dition: 30 m
ate the anod
al image of
s, the individ
h is independ
re 6.3 (a) Sross-section
ia the introd
nto the oxide
inversion, A
active ion et
min), reveal
layer betw
mTorr, 30 W,
de and cathod
f the interdig
dual electrod
dent of its he
chematic an SEM image
duced oxyge
e surface by V
AZ9620 resis
tching (RIE
ling ~10 μm
ween the 3D
, 22 SCCM
de current co
gitated Ni cu
de fingers p
eight, as show
nd (b) opticae of a single
93
en atoms afte
Van der Wa
st is dissolve
condition: 5
m thick inter
D nickel sc
methane, 45
ollectors. Fig
urrent collec
possess a pri
wn in Figur
al images ofporous Ni e
er the oxidiz
als force wit
ed in aceton
500 mTorr,
rdigitated N
caffolds is
5 SCCM H2
gure 6.3a an
ctors. Due to
ismatic shap
re 6.3c.
f 3D interdigelectrode that
zation treatm
thout it.
ne followed b
200 W, 20
Ni scaffolds.
then etched
2, 12 min), t
nd b show s
o the vertica
pe, and also
gitated curreat has a prism
ment, but wa
by SU8 resi
SCCM O2,
The expose
d using RI
to electricall
schematic an
al photoresi
the resultin
ent collectormatic shape.
as
st
2
ed
IE
ly
nd
st
ng
rs.
94
As shown in Figure 6.4a, Ni-Sn and MnO2 are sequentially electroplated onto the
Ni scaffold as anode and cathode, respectively. The electrodeposition methods follow the
procedures described previously,[1, 2] except that the pulsed voltage routine is modified to
be 0.2 s on and 10 s off, to ensure the conformal coating through the whole 3D structure.
Figure 6.4b (c) shows the cross-section micrograph of ~ 100 nm thick MnO2 (Ni-Sn)
film conformally coated on Ni scaffold after 15 cycles of pulsed depositions. The sample
is then immersed in LiOH and LiNO3 molten salts at 300 ºC for 30 min to lithiate the
MnO2. Finally, the cathode and anode are independently charged to 3.8 V and 0.01 V
versus lithium metal at 0.5 C, respectively. The electrolyte is a 1:1 ethylene carbonate :
dimethyl carbonate and 1 M LiClO4. Finally, the microbattery is capped with a PDMS
cover.
6.3 Electrochemical testing of microbatteries
The battery testing is carried out by galvanostatically charging and discharging
the cell between 3.2 V and 1.4 V at various C rates. 1 C stands for charging/discharging
the battery in an hour. At a C rate of N, the cell is charged/discharged at N times the 1 C
current.
FigurSn anconfo
6.3.1
been
its ca
often
large
interp
electr
preve
re 6.4 (a) Scnd MnO2 aormal MnO2
Improving m
Recently,
independen
apacity gradu
n drops to 90
capacity lo
phase (SEI)
rolyte decom
ent further el
chematic of s anode and coating and
microbattery
the half-cel
ntly studied.[
ually fades to
0% within th
oss in the be
layer. This
mposition at
lectrolyte re
battery elecd cathode, r
d (c) conform
y cycle life
ll cyclabiliti
1] The MnO
o 90% after
he first 5 cy
eginning is c
layer comp
the negative
duction duri
95
trodes after respectively
mal Ni-Sn co
es of MnO2
O2 cathode ge
50 cycles. H
ycles and th
caused by th
prises of ino
e electrode.[1
ing the cycli
independent. Cross-sect
oating on 3D
2 and Ni-Sn
enerally exh
However, the
hen maintain
he formation
organic and
10, 11] Once f
ing by block
t electrodepotion SEM im
D current coll
on 3D Ni sc
hibits a good
e capacity of
ns a steady d
n of the sol
organic pro
formed, the
king the elect
osition of Nmages of (blectors.
caffolds hav
d cycle life a
f Ni-Sn anod
decrease. Th
lid electrolyt
oducts of th
SEI layer ca
tron transpo
Ni-b)
ve
as
de
he
te
he
an
ort
96
while only allowing Li ions to travel through. In our microbatteries, where the nano-
porous electrodes inherently have a large surface area, the initial SEI formation can
significantly degrade the limited electrolyte, leading to a great lithium-ion loss in the
battery. Thus, it is important to produce the SEI layer at the anode before assembling the
full cell. In this work, we cycle the Ni-Sn electrode separately for 6 times before the full
cell tests. Figure 6.5 shows the improved cyclability of such microbattery, where the cell
possesses ~ 80% retained capacity after cycling for 100 times at 2 C.
Figure 6.5 Capacity retention of a microbattery for the first 100 cycles. The charge and discharge rates are ~ 2 C.
6.3.2 Power performance of the microbattery
To study the power performance, a microbattery cell that consists of 35 μm wide
electrode fingers and a PDMS cover is charged at 2 C and discharged at various C rates.
Figure 6.6 shows the discharge curves at those currents. At 1 C, the battery possesses a
0 10 20 30 40 50 60 70 80 90 1000.0
0.4
0.8
1.2
1.6
2.0
2.4
Nor
mal
ized
Ca
pac
ity
Ca
pac
ity (
h c
m-2
m-1 )
Cycle Number
0.0
0.2
0.4
0.6
0.8
1.0
97
volumetric energy density of 4.5 μWh cm-2 μm-1. At 1000C, it delivers 0.6 μWh cm-2 μm-
1 (energy density) and a 3.6 mW cm-2 μm-1 (power density). The capacity retention of this
microbattery versus C rates is presented in Figure 6.7, where the capacity was
normalized to a 1 C discharge. The cell exhibits excellent power performance and
cyclability at high C rates. Nearly 20% of the capacity is extracted from the battery at
1000 C. However, as I will demonstrate later, the capacity retention at high C rates is
closely related to the gap between the battery and the PDMS cover, as it governs the
effective ion pathways in the electrolyte. After being cycled for 5 times at each high C
rate, the battery only shows a small capacity loss (Figure 6.7).
Figure 6.6 Galvanostatic discharge of a microbattery at various C rates. The width of the electrode is 35 μm and the battery is covered by PDMS.
0 1 2 3 4 51.4
1.6
1.8
2.0
2.2
2.4
2.6
2.8
3.0
3.2
1000C500C 100C
20C
Vol
tage
(V
)
Energy Density (Wh cm-2 m-1 )
1C
98
Figure 6.7 Capacity retention of the microbattery cycled for 5 times at each C rate. The capacity of the cell at high C rates was normalized to the 1 C discharge.
6.3.3 Tuning electrode finger width
The desired height for a microbattery is often on the order of ~ 100 μm to achieve
high areal energy density. Recent advances in photolithography techniques have enabled
photoresist features with aspect ratios greater than 15, producing structures of a few
hundred microns tall and a few microns wide.[12-14] In principle, the electrode spacing
should be small such that the space in microbatteries is most efficiently utilized. To
simulate a practical microbattery, we keep the electrode spacing at 15 μm and vary the
electrode width to investigate the ion transport in 3D porous electrodes. Figure 6.8a
shows the discharge capacity retention of cells with 35 μm, 60 μm and 110 μm wide
electrodes, but without PDMS covers. Although increasing the electrode width should
lead to longer ion diffusion pathway inside electrodes, all three cells are found to have
1 6 11 16 21 260.0
0.2
0.4
0.6
0.8
1.0 2C
1000C500C
100C
20C
Nor
mal
ized
Cap
acity
Cycle Number
1C
simil
at 10
Figur
prima
disch
Figurwidthcells
ar and excep
00 C is near
re 6.8b. Wh
arily travel i
harge, as the
re 6.8 (a) Chs dischargewithout the
ptionally hig
rly 40% for
hen the batte
in the free s
total diffusio
Capacity reteed at variousPDMS cove
gh capacity r
all of them
eries are not
space above
on pathway
ention of uns C rates; (ber.
99
etention, e.g
. The reason
capped and
the cells oth
is shorter in
ncovered mib) schematic
g. the fractio
n for this ob
d have a thic
her than the
n the former
icrobatteriesc illustration
on of the reta
bservation is
ckness aroun
e lateral dire
case.
s with differn of ion tran
ained capacit
illustrated i
nd 8 μm, ion
ections durin
rent electrodnsports in th
ty
in
ns
ng
de he
100
The battery electrodes in a commercial product are generally tall (> 100 μm) and
fully packaged. Because there is little excessive electrolyte above the electrodes, ions are
forced to only diffuse in the lateral directions. To simulate this case, we cap the battery (8
μm tall) with PDMS slabs to make ions travel laterally. Figure 6.9a shows the optical
image of such covered battery electrodes immersed in electrolyte. Here we coat half of
the PDMS with gold in order to visualize the interface between the PDMS cover and the
electrode. Under a 50X objective, the gap between the PDMS and the electrode is
estimated to be ~ 5 μm, which however can still guide the majority of ions to travel in
this channel. Figure 6.9b shows the same sample configuration but with an external
pressure applied to the PDMS. The slightly curved interface suggests a tight contact
between the PDMS and the electrode. The tests in Figure 6.10 are performed on the cells
that are packaged in this manner. The data from the cell with 110 μm electrode width in
Figure 6.8a is also plotted here, which can be used to assess the performance of a
microbattery with 8 μm wide electrodes given the nature of the ion diffusion in that cell.
When raising the electrode width, the retained capacity at a constant discharge rate is
observed to decrease significantly due to the increased diffusion length in electrodes. At
1000C, only 3% of the capacity is extracted from the cell with 35 μm wide electrodes,
while in the previous uncapped cells nearly 40% of capacity is available as ions shuttle in
a shorter diffusion pathway.
FigurvisuacoatePDM
6.3.4
of th
respe
can i
holog
with
the h
electr
60%
situat
re 6.9 Optialize the coned with gold
MS to ensure
Tuning the
The ion d
e electrode
ectively. Sinc
indirectly co
graphic litho
~ 40% and ~
holographic
rodes with l
porous batte
tions.
cal images ntact betweend. (a) No prgood contac
electrode po
diffusion in e
following D
ce the battery
ontrol the e
ography. Tw
~ 60% pore
lithography
arger porosi
ery retains m
of battery en electrode aressure applct.
orosity
electrolyte p
D =D0ε/τ, w
y electrodes
electrode po
o microbatte
volume are
y. Figure 6
ity can offer
more capacit
101
electrodes seand PDMS, lied to PDM
phase also de
where ε and
are the inve
orosity via
eries (height
fabricated b
6.11 shows
r shorter diff
ty at high C
ealed by PDthe top half
MS; (b) exte
epends on th
τ stand for
erse of their
exposure co
t: 8 μm and
by changing
their powe
fusion pathw
rates in both
DMS coversf of the PDMernal pressu
he porosity a
r porosity an
photoresist t
onditions du
electrode w
the exposur
er performan
ways in all d
h covered an
s. In order tMS surface ure applied t
and tortuosit
nd tortuosity
templates, w
uring the 3D
width: 35 μm
e dose durin
nce. Becaus
directions, th
nd uncovere
to is to
ty
y,
we
D
m.)
ng
se
he
ed
Figurelectrin the
re 6.10 (a)rode width de cells with P
) Capacity discharged aPDMS cover
retention oat various C rs.
102
of PDMS-crates; (b) sc
covered micchematic illu
crobatteries ustration of i
of differenion transport
nt ts
103
Figure 6.11 Capacity retention of microbatteries of different porosities tested w/o PDMS covers.
6.4 Simulation and optimization of microbattery in COMSOL
A simple isothermal model for lithium ion batteries is developed in this section
based on an assumption that ionic charges in our PDMS-capped batteries only travel in
one dimension. The modeling is carried out in COMSOL which is well-known for its
strength at solving complex differential equations using finite element analysis.[15, 16]
6.4.1 Evaluation of lithium ion solid state diffusion
Before constructing a detailed simulation that includes lithium-ion diffusions in
both liquid phase and solid phase, we first simplify the task by assuming that charge
species travel infinitely fast in electrolyte and only experience resistances from solid-state
diffusion. Such analysis can help us assess the lithium ion diffusion in active materials at
1 10 100 10000.0
0.2
0.4
0.6
0.8
1.0 60% porosity_uncovered 40% porosity_uncovered 60% porosity_packaged 40% porosity_packaged
Nor
mal
ized
Cap
acity
C rate
high
100 n
and
diffus
condu
exper
indep
the p
positi
open
Figur
differ
drops
discharge ra
nm apart fro
60 nm Ni-
sivity for Mn
uctivity and
rimentally
pendently im
percentage o
ive electrod
circuit volta
re 6.12 SEM
Figure 6.
rent discharg
s below 1.4 V
ates. This mo
om each othe
Sn, for the
nO2 positive
diffusivity f
measured o
mported into
f the full sta
es are 0.95
age 3.2 V.
M image of M
.13 displays
ge rates. Th
V. At low C
odel conside
er and coated
e cathode a
e electrode a
for Ni-Sn are
open circui
COMSOL
ate of charg
and 0.05, c
MnO2 layer t
the simulat
e model def
rates, from
104
ers two infin
d with 140 n
and anode,
are 5.56 S/m
e 9.00 10
it voltage
as a functio
ge). The init
orrespondin
that is confor
ted discharg
fines the end
1 C to 100 C
nitely large p
nm MnO2 (a
respectively
m and 5.00
S/m and 8
profiles fo
on of state o
tial SOC val
ng to a fully
rmally coate
ge curves as
d of dischar
C, the capac
parallel electr
as shown in
y. The cond
10 m2/s
.00 10
or both el
of charge (SO
lues for the
y charged ba
ed on 3D Ni
a function o
rge when the
city retention
rodes that ar
Figure 6.12
ductivity an
s,[15] while th
m2/s.[17] Th
lectrodes ar
OC, which
negative an
attery with a
scaffold.
of capacity a
e cell voltag
n is very high
re
2)
nd
he
he
re
is
nd
an
at
ge
h,
105
while at high C rates such as 500 C and 1000C the capacity retentions are 60% and 35%,
respectively, suggesting that solid-state diffusion only limits the discharge at high C rates
and the relatively poor capacity retentions observed at low C rates in experiment (Figure
6.8a) might be caused by other mechanisms such as liquid-phase diffusion. The
calculated lithium concentration in electrode (Figure 6.14a) also supports the fact that
lithium ion diffusion in active materials is limited at high rates since a large fraction of
lithium ions still remains in the negative electrode when the discharge ends. The lithium
salt concentration in electrolyte at the end of discharge is shown in Figure 6.14b. The
slightly polarized concentration profile is consistent with the assumption of our model.
Figure 6.13 Simulated galvanostatic discharge of thin film batteries. The coatings of active materials on negative and positive electrode are 60 nm and 140 nm, respectively, and the gap between two electrodes is 100 nm.
0.00 0.01 0.02 0.03 0.04 0.051.4
1.6
1.8
2.0
2.2
2.4
2.6
2.8
3.0
3.2
Vol
tage
(V
)
Capactiy (A.U.)
1 C 20 C 50 C 100 C 500 C 1000 C
106
Figure 6.14 Simulated (a) lithium concentration in the electrode and (b) electrolyte concentration of thin film batteries at the end of discharge.
0 40 80 120 160 200 240 2800
5000
10000
15000
20000
25000
30000
35000
40000
45000
1000 C
500 C
100 C50 C20 C
Lith
ium
Co
ncen
tra
tion
(mo
l/m3 )
Position (nm)
1 C
(a) (b
0 40 80 120 160 200 240 280999.0
999.5
1000.0
1000.5
1001.0
Ele
ctro
lyte
Con
cent
ratio
n (m
ol/m
3 )
Position (nm)
1 C 20 C 50 C 100 C 500 C 1000 C
(b)
107
6.4.2 1D isothermal modeling of microbatteries
Here we adopt Newman’s approach in FEA modeling.[18] The goal is to gain
insights on electrochemical processes of fast charging/discharging in our microbatteries,
while at the same time pointing out the design strategies for optimizing their
performance. Because of the difficulties in defining the 3D holographic electrodes and
the conformal active material coatings in the model, we approximate the 3D Ni scaffold
and its active material coatings as an ensemble of spherical particles. As illustrated in
Figure 6.15, the model contains the following elements:
1) Ni current collectors that accounts for electronic conduction in electrodes;
2) Conductive particles and active material spheres that form the electrodes. The
diameter and filling fraction of Ni-Sn spheres in the negative electrode is 60 nm
and 30%. For the positive electrode, the diameter and filling fraction of MnO2
particles is 140 nm and 30%. The electrode width in the simulation represents half
of the dimension in experiment, because each anode (cathode) interacts with two
adjacent cathodes (anodes).
3) Electrolyte made of 1M LiClO4 in 1:1 EC : DMC. The effects of concentration on
ionic conductivity are introduced from experimentally measured values;[19]
4) Two electrodes are separated by 15 μm;
5) Ionic charge transports between electrodes and electrolyte;
6) Equilibrium potential obtained experimentally from discharge curves to entail
Butler-Volmer electrode kinetics.
During the simulation, the electrical potential in the electron conducting phase is
calculated using Ohm’s law based on charge balance. For the porous electrodes, effective
liquid
ε and
wher
summ
fully
(Ni-S
Figur
d-phase diffu
d tortuosity τ
e and
maries all th
charged, w
Sn) and at mi
re 6.15 Sche
usion coeffic
τ into accoun
/ ,
/ ,
are the d
he parameter
when the Li+
inimum in th
ematic illustr
cient Deff an
nts, they are
diffusivity an
rs used in th
concentrati
he positive e
ration of 1D
108
d conductiv
nd conducti
he simulatio
ion is at its
electrode (M
isothermal
ity σeff are u
ivity of the
ons. The mo
maximum i
MnO2).
microbattery
used. Taking
e electrolyte
odeled batte
in the negat
y model.
g the porosit
(6.1
(6.2
e. Table 6.
ry is initiall
tive electrod
ty
1)
2)
.1
ly
de
109
Table 6.1 Values of parameters used in COMSOL modeling.
6.4.3 Optimization of electrode width
The influence of the porous electrode width on the battery performance is studied
in the section. The capacity retention is simulated at various discharge rates for
microbatteries of different electrode widths and presented in Figure 6.16. The capacity at
high C rates is normalized to that of 1C, where the cell is considered to undergo a quasi-
equilibrium discharge, as all the capacity of the battery can be extracted at such rate.
There exists a critical dimension for the electrode width in each C rate curve, below
which the retained capacity only changes slightly versus electrode width, indicating that
Symbol Description Value
Ds_neg Solid phase Li‐diffusivity Negative 8e‐12[m^2/s]
Ds_pos Solid phase Li‐diffusivity Positive 5e‐16[m^2/s]
R_neg Particle radius Negative 30e‐9[m]
R_pos Particle radius Positive 50e‐9[m]
Ks_neg Solid phase conductivity Negative 9e6[S/m]
Ks_pos Solid phase conductivity Positive 5.56[S/m]
Dl Salt diffusivity in Electrolyte 1.2e‐10[m^2/s]
Εs Solid phase electrode vol‐fraction 0.3
Εl Liquid phase electrode vol‐fraction 0.3
i_1C 1C discharge current density 3.6[A/m^2]
τ Tortuosity of electrode 2.7
110
these batteries have been discharged to a similar state of charge. However, above this
value, the capacity dramatically decreases when the electrode width increases. As I will
discuss later, this critical length is closely related to the ion diffusion length in liquid
phase. Compared to the experimental data in Figure 6.10a, the simulation follows the
same trend that the high-rate capacity retention generally becomes worse with wider
electrodes, though they only have a good agreement at high discharge rates. When the
electrodes are very narrow, the cells experience less resistance from ion diffusion in
electrolyte. In extreme cases, e.g. at the intercept points in Figure 6.16, the capacity
retention reaches maximum and the battery is only limited by solid-state ion diffusion.
Figure 6.16 Simulated capacity retention for cells with different electrode width at various discharge rates.
20 40 60 80 100 1200.0
0.2
0.4
0.6
0.8
1.0
Nor
mal
ized
Cap
acity
Electrode Width (m)
1000C
500C
100C
20C
1C
111
For interdigitated microbatteries that have a fixed footprint or volume, cells with
wider electrodes can provide larger energy density. This is because the number of gaps
between the electrodes (Figure 6.17a) is reduced, even if less fraction of the total energy
is discharged. We plot the simulated energy density for cells with different electrodes at
various discharge rates in Figure 6.17b (the curve is normalized to the energy density
measured experimentally from the cell with 35 μm wide electrodes). For 1 C discharge,
the cell can release all the energy, and therefore the energy density grows monotonically
as the electrode gets wider, obeying W/(W+G), where W and G are the electrode width
and the spacing between them. At high C rates, the energy density initially increases but
reaches the peak at some critical length. The lithium concentration profile in the
electrodes and the salt concentration profile in the electrolyte are simulated for cells with
21 μm, 56 μm and 84 μm wide electrodes at the end of a 20 C discharge. These three
electrode widths correspond to three different regimes in Figure 6.18a. For W = 21 μm,
the electrolyte is only slightly polarized and the lithium concentration reaches the
maximum in the positive electrode and the minimum in the negative electrode at the end
of discharge. Thus, the cell can be fully discharged when W < 56 μm. At W = 56 μm,
there is a large gradient in the lithium salt concentration across the cell and the
concentration nearly drops to zero at the center of the positive electrode. However, in this
cell ion diffusion is still sufficient to deliver charges through the entire electrodes so that
a complete lithiation is achieved in the cathode. When W = 84 μm, the salt is completely
depleted at the position X > 70 μm as lithium ions are not able to travel this far in the
electrode during discharging, which ends the electrochemical processes prematurely and
only extracts 40% of the total capacity.
Figurwhenelectr
estim
tortuo
t is th
Acco
6.1 a
summ
rates,
of hig
re 6.17 (a) Sn raising elerode width.
The ion d
mated as √
osity of the
/
he discharge
3600
ount to Eqn.
and compare
marized in T
, which supp
gh-power mi
Schematic ilectrode wid
diffusion len
, where D i
e electrode,
/ ,
time that ca
0/C_rate ,
6.3 and Eq
e this value
Table 6.2, th
ports our hyp
icrobatteries
llustration thdth; (b) sim
ngth in the p
is the effect
an be evaluat
qn. 6.4, we c
to the criti
hese two qu
pothesis that
s. The result
112
hat the arealmulated ener
porous electr
tive diffusiv
ted from the
calculate √
ical electrod
uantities exh
t liquid-phas
s in Table 6
l/volumetric rgy density
trode at a gi
ity corrected
discharge ra
from the p
de width at
hibit great si
se diffusion c
.2 also provi
energy denfor cells w
iven diffusiv
d with the p
ate as,
parameters li
each discha
imilarity at
can limit the
ide a simple
sity increasewith differen
vity D can b
porosity an
(6.3
(6.4
isted in Tabl
arge rate. A
all discharg
e performanc
e guidance fo
es nt
be
nd
3)
4)
le
As
ge
ce
or
desig
shoul
Figurdischare ca
gning optima
ld be made t
re 6.18 (a) harge rate. Thalculated for
al microbatt
the same as t
Calculated he lithium cor cells with v
teries that f
the effective
energy densoncentrationvarious elect
113
for a given
ion diffusio
sity as a funn in electrodetrode widths
discharge r
on length in l
nction of ele and salt co: (b) 21 μm;
rate the ele
liquid phase
lectrode widoncentration (c) 56 μm;
ectrode widt
e.
dth at a 20 in electrolyt(d) 84 μm.
th
C te
114
Table 6.2 Comparisons between simulated critical electrode width and calculated electrolyte-phase diffusion length at various discharge rates.
6.5 Conclusions
We develop high-power lithium-ion micro batteries based on 3D holographic
lithography, a technique that has excellent design flexibility and great potential for on-
chip applications. The photolithographically defined electrode patterns allow the porous
electrode to grow vertically, which is important for making tall (high areal energy
density) and narrow (high power density) electrode arrays. The cycle life of our batteries
is much improved by pre-cycling the anode to form the necessary solid-electrolyte-
interphase layer. The battery exhibits excellent capacity retention at both low and high
discharge rates. The design parameters for optimizing the energy density are calculated at
certain power based on experiments and modeling. The liquid-phase ion diffusion is
responsible for the decreased energy density in the cells with wide electrodes.
Calculations indicate that for a certain discharge rate, the value of the optimal electrode
width is close to the liquid-phase ion diffusion length in the electrode.
C‐rateElectrode widthat peak energy density (μm)
Diffusion length
(μm)
20 56.2 48.8
50 34.7 30.9
100 24.2 21.8
500 10.0 9.8
1000 7.0 6.9
115
6.6 References
[1] H. G. Zhang, X. D. Yu, P. V. Braun, Nat Nanotechnol 2011, 6, 277. [2] J. H. Pikul, H. G. Zhang, J. Cho, P. V. Braun, W. P. King, Nat Commun 2013, 4. [3] F. Chamran, Y. Yeh, H. S. Min, B. Dunn, C. J. Kim, J Microelectromech S 2007, 16, 844. [4] H. S. Min, B. Y. Park, L. Taherabadi, C. L. Wang, Y. Yeh, R. Zaouk, M. J. Madou, B. Dunn, J Power Sources 2008, 178, 795. [5] M. Nathan, D. Golodnitsky, V. Yufit, E. Strauss, T. Ripenbein, I. Shechtman, S. Menkin, E. Peled, J Microelectromech S 2005, 14, 879. [6] T. S. Arthur, D. J. Bates, N. Cirigliano, D. C. Johnson, P. Malati, J. M. Mosby, E. Perre, M. T. Rawls, A. L. Prieto, B. Dunn, Mrs Bull 2011, 36, 523. [7] J. Park, S. D. Wang, M. Li, C. Ahn, J. K. Hyun, D. S. Kim, D. K. Kim, J. A. Rogers, Y. G. Huang, S. Jeon, Nat Commun 2012, 3. [8] M. Miyake, Y. C. Chen, P. V. Braun, P. Wiltzius, Adv Mat 2009, 21, 3012. [9] Y. C. Chen, J. B. Geddes, J. T. Lee, P. V. Braun, P. Wiltzius, App Phys Lett 2007, 91. [10] P. Verma, P. Maire, P. Novak, Electrochim Acta 2010, 55, 6332. [11] M. B. Pinson, M. Z. Bazant, J Electrochem Soc 2013, 160, A243. [12] H. Lorenz, M. Despont, N. Fahrni, J. Brugger, P. Vettiger, P. Renaud, Sensor Actuat a-Phys 1998, 64, 33. [13] M. C. Peterman, P. Huie, D. M. Bloom, H. A. Fishman, J Micromech Microeng 2003, 13, 380. [14] A. del Campo, C. Greiner, J Micromech Microeng 2007, 17, R81. [15] C. W. Wang, A. M. Sastry, J Electrochem Soc 2007, 154, A1035. [16] V. Zadin, H. Kasemagi, A. Aabloo, D. Brandell, J Power Sources 2010, 195, 6218. [17] E. Hosono, H. Matsuda, I. Honma, M. Ichihara, H. S. Zhou, J Electrochem Soc 2007, 154, A146. [18] M. Doyle, J. Newman, A. S. Gozdz, C. N. Schmutz, J. M. Tarascon, J Electrochem Soc 1996, 143, 1890. [19] K. Xu, Chem Rev 2004, 104, 4303.
116
CHAPTER SEVEN
CONCLUSIONS AND FUTURE WORK
7.1 Conclusions
This thesis research has focused on using 3D periodic structures to enhance light-
matter interaction and energy storage. It began with investigation of fundamentals on
controlling spontaneous emission via 3D photonic crystals in Chapter 2, which
demonstrated that the photonic band structure of silicon inverse PhCs can be specifically
and finely tailored using ALD at an intermediate step between template fabrication and
silicon inversion. This was coupled with the incorporation of rare earth nanoparticle
emitters into the silicon inverse PhCs at a well-defined location provided by a simple
experimental procedure. The narrow emission linewidth of rare earth nanoparticles and
the photonic band gap tuning enabled the study on the effect of the photonic DOS on
spontaneous emission. Time-resolved experiments revealed that the emission rate of
embedded emitters can be strongly manipulated by the stopgap of silicon inverse opals;
up to a 61% change of emission decay rate was observed between the enhanced and the
inhibited SE.
To fully unitize the optical properties of 3D photonic crystals, functional defects
such as microcavities must be created. In Chapter 3, a new type of vertical microcavity
was developed by combining 3D Si inverse opal PhCs and Si thin films. Such structure
was first modeled in FDTD to calculate the corresponding cavity modes and Q-factor.
The 3D PhC microcavity was then assembled by advanced transfer-printing with micro-
structured stamps. Since each layer of the sandwich structure was fabricated
117
independently, this design can potentially allow us to specifically pattern cavity layer to
achieve certain modal profile or incorporate emitters to control their emission properties.
The previous two projects were based on colloidal self-assembled photonic
crystals, which are easy to fabricate but have plenty of undesired defects. In comparison,
3D holographic lithography can achieve structures with large area and free of synthetic
defects. Chapter 4 has demonstrated that this technique combined with transfer-printing
provided an excellent means to embed a broad variety of artificial defects into
holographic PhCs. The introduced defects can either remain unchanged or be dissolved
during the process of holographic lithography, leading to either dielectric or air defects.
Moreover, colloidal quantum dots were also placed at a specific location inside 3D PhCs.
By changing their surrounding photonic environments, both suppressed and enhanced
emission were observed from the localized emitters. Our method provides a robust
capability for adding functionalities into 3D PhCs, opening the door for realizing
integrated photonic circuits in 3D holographic structures.
Chapter 5 has demonstrated that a modified transfer-printing technique enabled
the formation of high-quality, hybrid vertical microcavities that feature any arbitrary
light-emitting layer sandwiched between PSi DBRs. The resonant electromagnetic mode
of this cavity structure coupled with the emission spectrum of the light-emitting entity to
severely suppress emission everywhere in the stopband of the DBR, except at the spectral
position of the cavity mode, where an emission enhancement occured. This redistributed
emission was observed separately from a PbS QD-doped polymer film and a GaAs solid-
state thin film in a microcavity configuration with PSi DBRs. The addition of a PSi CCL
extended the inherent index modulation capabilities of PSi to the cavity. The PSi CCL
118
provided a mechanism for tuning the hybrid microcavity’s resonant cavity mode and
emission spectrum in both coarse and fine natures by way of intentional porosity
variation and porosity infiltration, respectively.
So far, I have showed that 3D periodic structures can control light-matter
interactions. However, their applications are not only limited to photonics. In fact, the 3D
mesoporous network can significantly facilitate electrochemical processes. This was
demonstrated in Chapter 6, where a high-power lithium-ion micro battery was developed
from 3D holographic structures. In particular, the photolithographically defined electrode
patterns allowed the porous electrode to grow vertically, which was important for making
tall (high areal energy density) and narrow (high power density) electrode arrays. The
cycle life of the batteries was much improved by pre-cycling the anode to form the
necessary solid-electrolyte-interphase layer. The battery exhibited excellent capacity
retention at both low and high discharge rates. The design parameters for optimizing the
energy density were calculated at certain power based on experiments and modeling. The
liquid-phase ion diffusion was found to be responsible for the decreased energy density in
the cells with wide electrodes. Calculations indicated that for a certain discharge rate, the
value of the optimal electrode width was similar to the liquid-phase ion diffusion length
in the electrode.
7.2 Future work
Transfer-printing has enabled the formation of high-quality hybrid vertical
microcavities that consist of porous silicon DBR and extrinsic light emitters. Although
such effort has led to realizations of strong emission manipulation of the emitting entities
119
located inside the microcavity, the anticipated lasing behavior has not been observed.
One possible reason is the lack of proper gain media. In particular, the quantum yield of
the PbS quantum dots embedded in porous silicon microcavities was less than 40%,
which is far from the desired quantum efficiency for a laser. The heterogeneous GaAs
thin film, though possessing high quantum yield, does not have the correct structural
design. The strong absorption at cavity nodes results in a very lossy microcavity. In order
to realize lasing in a vertical microcavity, the gain media must satisfy two conditions: 1)
they must have near unity quantum efficiency; 2) the absorption at cavity nodes must be
reduced. Two promising candidates are III-V compound quantum dots and quantum
wells. These emitters are highly efficient due to the quantum confinement. Also, because
of the accurate spatial control in epitaxial growth, gain media can be structured so that
they only appear at antinode positions. I believe that with appropriate gain media, porous
Si-III-V hybrid coherent light sources can be realized in a new wavelength regime
between 900 nm and 1100 nm.
The success of incorporating functional defects in 3D holographic photonic
crystals provides an excellent opportunity for sensing applications. For example, a thin
porous silicon film can be embedded between two 3D holographic photonic crystals to
form a vertical microcavity. The cavity mode of such structure will shift as the extrinsic
chemical species change the refractive index of porous silicon. While sensors based on
porous silicon microcavities have been widely used for this purpose, the hybrid porous
silicon SU8 structure can potentially outperform them for the following reasons: 1) the
sub-micron pores in 3D SU8 photonic crystals can facilitate gas or chemical to easily
reach the porous silicon cavity; 2) SU8 exhibits a highly different surface from porous
120
silicon (e.g. SU8 is highly hydrophobic.), which can allow foreign species to selectively
attach to porous silicon only. However, in monolithic porous structures, gas or chemical
have to diffuse through a thick mesoporous DBR before reaching the cavity and also coat
the entire structure at the same time. Therefore, hybrid structures require less amount of
material infilling to shift the refractive index of the cavity and thus can provide a better
sensitivity.
So far photonic crystals have proved to be a powerful tool to control light-matter
interactions. Previous works have focus on utilizing them to modify either emission or
absorption. However, it is possible to realize more complex functions in photonic crystals
by adding both emitters and absorbers. For example, two coupled microcavities can be
created with emitters located in one cavity and absorbers in the other one. If the two
cavities have the same cavity mode and quality factor Q, the emission and absorption are
both enhanced by a factor of Q. However, since the emission from the emitters can also
be absorbed by the absorbers, the enhancement of the system response is Q2. This
experiment can be achieved with both 3D photonic crystals and porous silicon DBRs
using the transfer-printing technique.
The microbatteries developed in this thesis have shown excellent power
performance and cycle life. However, since the current collector is directly inverted from
the holographic template, it exhibits a high volume filling fraction (40%~60%), which
limits the growth of the active materials and thus the energy density of the battery. This
problem can be solved either by developing a new type of 3D bicontinuous template that
has high volume fraction (70% ~ 90%) or by fabricating the current collector via metal
CVD. Moreover, the height of the current microbatteries is ~ 10 μm, which is primarily
121
limited by poor adhesion between SU8 and ITO glass (thick SU8 film often delaminates
during development due to the stress caused by volume change.) and the difficulty in
removing the SU8 in the presence of nickel current collect (because nickel also reacts
with O2 plasma and significantly reduces their diffusion length). In this case, a 3D
template that is made of low-absorption positive photoresists is preferred due to the fact
that they exhibit little volume change upon exposures and can be easily removed by
organic solvent or strong basic solutions.
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