Observation of negative charge trapping and investigation of its physicochemical origin in newly synthesized poly(tetraphenyl)silole siloxane thin films

Article abstract of DOI:10.1021/ja1108112

A new kind of organic-inorganic hybrid polymer, poly(tetraphenyl)silole siloxane, was invented and synthesized for realization of its unique charge trap properties. The organic portions consisting of (tetraphenyl)silole rings were responsible for negative

Full text of DOI:10.1021/ja1108112

ARTICLE
pubs.acs.org/JACS
Observation of Negative Charge Trapping and Investigation of Its
Physicochemical Origin in Newly Synthesized Poly(tetraphenyl)
silole Siloxane Thin Films
†
‡
§
,‡
,†
Jin-Kyu Choi, Seunghyun Jang, Ki-Jeong Kim, Honglae Sohn,* and Hyun-Dam Jeong*
†
Department of Chemistry, Chonnam National University, Gwangju 500-757, Republic of Korea
‡
Department of Chemistry, Chosun University, Gwangju 501-759, Republic of Korea
§
Pohang Accelerator Laboratory, Pohang University of Science and Technology, Pohang 790-784, Republic of Korea
S Supporting Information
b
ABSTRACT: A new kind of organicꢀinorganic hybrid polymer, poly(tetraphenyl)-
silole siloxane, was invented and synthesized for realization of its unique charge trap
properties. The organic portions consisting of (tetraphenyl)silole rings were respon-
sible for negative charge trapping, while the SiꢀOꢀSi inorganic linkages provided the
intrachain energy barrier for controlling electron transport. The polysilole siloxane
dielectric thin films were fabricated by spin-coating and curing of the polymers,
followed by characterization with spectroscopic ellipsometry (SE), near edge X-ray
absorption fine structure spectroscopy (NEXAFS), and photoemission spectroscopy
(PES). The abrupt increase in density and decrease in thickness of the thin film at a
curing temperature of 100 °C was attributed to a thermodynamically preferred state in
the nanoscopic arrangement of the polymer chains; this was due to cofacial πꢀπ
interactions in a skewed manner between peripheral phenyl groups of the
(
tetraphenyl)silole rings of the adjacent polymer chains. Using the NEXAFS spectrum
to assess high electron affinity, the LUMO energy level of the dielectric thin film cured at 150 °C was positioned 1 eV above the
Fermi energy level (E ). The electron trapping of the dielectric thin films was confirmed from the positive flat band shift (ΔV ) in
F
FB
the capacitanceꢀvoltage (CꢀV) measurements performed within the metalꢀinsulatorꢀsemiconductor (MIS) device structure,
which strongly verified the polymer design concept. From the simple kinetics model of the electron transport, it was proposed that
the flat band shift (ΔV ) or trap density of the negative charges (|F|) was logarithmically proportional to the decay constant (β) for
FB
the electron-tunneling process. When a phenyl group of a silole ring in a polymer chain was inserted into the two available phenyl
groups of another silole ring in another polymer chain, the electron transfer between the groups was enhanced, decreasing the trap
density of the negative charges (|F|). For the thermodynamically preferred state generating the high refractive index, the distance
between the two phenyl groups of the adjacent polymer chains was estimated to be in the range of 0.27ꢀ0.36 nm.
’
INTRODUCTION
inorganic networks. This is generally processed by chemical inter-
4,5
actions through H-bonding and π-bonding. In class II type
materials, part of the organic and inorganic components are linked
through strong chemical bonds (covalent, ionic, dative).
Compared to the materials made entirely from either organic
or inorganic components only, organoꢀinorganic hybrid materials
are materials in which organic and inorganic components are mixed
at a molecular level to give superior mechanical, thermal, optical, and
Polyorganosiloxane polymers are one of the most representa-
tive organicꢀinorganic hybrid materials of class II type in which
organic groups and inorganic ꢀSiꢀOꢀSiꢀ moieties are con-
nected by SiꢀC covalent bonds, as shown in Figure 1a. These
polymers are easily synthesized from monomers containing
1ꢀ3
electrical properties. The nanostructures and physical properties
of the organicꢀinorganic hybrid materials are dependent upon not
only the molecular structure of the organic or inorganic components
themselves but also their synergetic effects. One of the most crucial
points in designing new organicꢀinorganic hybrid materials is to
consider the type of chemical bonds and interfacial area between the
organic and inorganic components from which its physical proper-
ties will be controlled. These materials are basically classified into
3
SiꢀC bonds as they retain the covalent character of the sp
hybridization and maintain the SiꢀC bonds, even from nucleo-
3
philic attack. For example, hydrolysis/condensation polymeri-
zation (denoted solꢀgel polymerization in a materials science
0
1
context) with organo-substituted silicic acid esters (R Si(OR) ),
n
4ꢀn
two types according to the nature of the interface. In class I type
materials, no covalent or ionic bonds are present between the
organic and inorganic components, whereby the material is simply
synthesized by embedding organic components within growing
Received: December 2, 2010
Published: May 04, 2011
r 2011 American Chemical Society
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Journal of the American Chemical Society
ARTICLE
Figure 1. (a) Typical example of organicꢀinorganic hybrid materials: silsesquioxane polymer. The polymer is classified as a polyorganosiloxane. The
SiꢀCH groups generate a molecule-free volume to decrease the dielectric constant, while the ꢀSiO3/2ꢀ structures provide high thermal and
mechanical stabilities. (b) Polarizability of a material is varied as a function of frequency. At lower frequencies, orientation, ionic, and electronic
3
1
4
polarization mechanisms have to be counted using a Debye equation. Around the higher frequency of 4.74 ꢁ 10 Hz (633 nm), electronic polarization is
the main contributor, which is estimated by the LorenzꢀLorentz equation.
or hydrosilylation with vinyl-containing silicone and H-terminated
silicone monomers, are usually employed in the syntheses of such
polyorganosiloxane polymers.
The polyorganosiloxane polymers are basically dielectric
materials (abbreviated as dielectrics). What kinds of materials
are typically assigned to dielectrics? If electric polarization is
provoked inside a material as an electric field is applied through it,
the material is regarded as dielectric. Most insulator materials can
be dielectrics because they demonstrate an electric polarization
phenomenon when applying an electric field. Specifically, most
of the materials containing SiꢀO bonds, such as polyorganosi-
loxane polymers, are not only insulators owing to high band gap
values but also dielectrics because of polarization.
In order to investigate its physical properties and dielectric
applicability, it is necessary to understand dielectric polarization,
which has been treated as relatively less important in the field of
chemistry within recent decades, although being an ever compre-
hensive theme within physics, chemistry, materials science, and
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Journal of the American Chemical Society
ARTICLE
only inherits the advantages of conventional nonvolatile flash
memory but also further improves the stability of information
16
storage within the memory devices. The improved storage
stability is due mainly to electrical isolation between the nano-
dots, where the memory function remains undisturbed by the
charge leakage casually generated between a single nanodot and
the underlying silicon layer, thereby allowing higher density and
lower power operation for the memory device. One of the
widespread concepts for realizing this is the manner in which
silicon or gold nanoparticles are embedded in a floating gate layer
fabricated with SiO dielectrics; this is one of the most recent
2
topics in electrical engineering fields, as shown in Figure 2.
Injection of an electron occurs from the channel region via direct
tunneling when the control gate is forward biased with respect to
the source and drain. The resulting stored charges within the
nanoparticles screen out the electric field coming from the
control gate and modify the threshold voltage (ΔV ) of the
th
device, whose approximate magnitude for a single electron per
16b
nanoparticle is given by:
ꢀ
ꢁ
qnwell
εOX
1 εOX
2 εSi
ΔVth ¼
tcntl þ
twell
ð1Þ
where ΔV is the threshold voltage shift, tcntl, the thickness of
th
the control oxide under the gate, twell, the linear dimensional
of the nanoparticle well, ε’s are the relative permittivities, q is the
magnitude of electronic charge, and nwell is the density of the
nanoparticles.
Figure 2. (a) Schematic of nanodot memory device, where nanodots
are used as charge trap sites. The improved storage stability is due mainly
to electrical isolation between the nanodots. (b) Writing and erasing
operations of the nanodot device. An injection of an electron occurs
from the channel region via direct tunneling when the control gate is
forward biased with respect to the source and drain.
Moreover, the concept of hybridization of nonvolatile flash
memory or nanodot memory with an organic thin film transistor
12ꢀ15
(OTFT) has been also pursued.
This is targeted at extend-
ing the advantages of the OTFT, which include low cost, solution
processability, and a large area fabrication, all of which enable
fabrication of memory devices onto flexible plastic substrates.
Compared to conventional flash memory that employs solid-
state transistors, numerous applications of flexible electronics
require nonvolatile data storage. For example, in Baeg’s ap-
proach, polymer electrets with demonstrated charge trapping
capabilities are used as charging layers within the OTFT device,
imitating floating gate layer trapping charges in the flash
electrical engineering. The reason is that the theory of dielectric
polarization was established prior to 1960; after which, application
of dielectrics drew much of the attention. However, since the
recent trends in the development of electronic and optoelectronic
devices became necessary for highly integrated, low-cost, solution
processability and large area processing, the design of new di-
electrics based on the theory of dielectric polarization has become
6,7
more strongly recommended. The design and synthesis of low
dielectric constant materials (low-k) using polyorganosiloxane
polymers is one of the most representative examples of utilization
of this theory. The presence of organic groups in polyorganosilox-
anes generates molecular-free volumes that provide the decrease in
its refractive index and dielectric constant, which is explained using
12
memory. Moreover, integration of the solution-processed
nanodot layer into the OTFT has been investigated, mimicking
nanodot memory for the solid-state devices, where gold nano-
particles as charge trapping centers are involved with a self-
14
assembled block copolymer.
6
LorenzꢀLorentz and ClausiusꢀMosotti equations, respectively,
The authors have initiated compelling research into charge
trapping polyorganosiloxane dielectrics, which can be utilized as
a solution-processed floating gate layer or nanodot layer in
as summarized in Figure 1b. Moreover, the high dielectric constant
in select silicateꢀsilsesquioxane hybrid materials has been ascribed
to the existence of SiꢀOH groups, which has been explained using
10,11
OTFT-based nonvolatile memory devices.
As a first trial,
8
Debye’s equation. Interestingly, the high refractive indices of
the authors tested a nanocomposite dielectric as the nanodot
layer, wherein porphyrin molecules and ethylene-bridged silses-
quioxane dielectrics were used as the charge trap center and
methyleneꢀbiphenylene-bridged silsesquioxane materials were
also explained by considering the high electronic polarizability of
9
10
the bridged methyleneꢀbiphenylene group.
insulating matrix in a single nanocomposite layer, respectively.
In addition to the dielectric polarization of the polyorganosi-
loxanes, in recent years, charge trap properties have also been studied
as another important aspect of dielectrics within this group.
First from this study, the charge trap properties from the
capacitanceꢀvoltage measurements (CꢀV) in the metalꢀ
insulatorꢀsemiconductor (MIS) device were confirmed. In
spite of this, the phase separation problem between the porphyr-
in molecules and the silsesquioxane dielectrics, as well as
poor reproducibility in the charge trapping phenomenon,
remained inevitable using this approach because of the different
chemical natures of the two components. Second, the charge trap
properties of the aminopropyl-silsesquioxane thin films, where
10,11
This
is the main theme of this article. The authors contend the investiga-
tion presented herein to be very important because of the potential
applicability of these charge trap properties toward next-generation
12ꢀ15
memory devices and flexible electronics.
It has been well-known that nanodot layers bearing isolated
charge trap sites can be used for nanodot memory, which not
7
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Journal of the American Chemical Society
ARTICLE
Figure 3. (a) Molecular structure of silole. R , R , R , R , R , and R are alkyls, aryls, or halides. (b) Molecular structure of the poly(tetraphenyl)silole
1
2
3
4
5
6
siloxane polymer (n = 7). (c) Step potential model (SPM) for explaining electron transport through the poly(tetraphenyl)silole siloxane dielectrics. For
simplification, only three polymers, wherein the number of monomers was seven (n = 7) for each polymer, were considered. Tentatively, the energy
barrier of the interchain electron transfer was assumed to be higher than that for an intrachain transfer.
formation of two-fold N(Sit) was proposed to be the origin of
maintained in the polyorganosiloxane, the directly linked organic
moieties for the charge trap properties can be chosen in a variety
ways from the π-conjugated molecules. This is due to the low
HOMOꢀLUMO gap of their π-conjugated organics that usually
allows for high electron affinities and low ionization energies, thus
enhancing charge trapping in real applications of the devices.
Among the various organic moieties this lab has been keenly
interested in is the silacyclopentadiene (silole) ring for use as a
charge-trapping center, whose molecular structure is shown in
Figure 3a. In the silole, the LUMO energy level was much lower
compared to other heteroaromatic molecules, due to mixing of the
σ* orbital of the silicon atom with the π* orbital of the butadiene
2
11
electron trapping, were investigated. The limitation of this
material lies in its difficulty in engineering the charge trap
density; this is due to the defect sites responsible for the electron
trapping being tenuously controlled in an artificial manner.
For this article, in order to circumvent the phase separation
problem and the overall lack in engineering capabilities of charge
trap properties in the above two polyorganosiloxane materials, this
lab adopted a new material architecture that combines the
insulating properties of polyorganosiloxanes and the charge trap
properties of π-conjugated organic moieties, where the two
components were linked through direct covalent bonds as class
II organicꢀinorganic hybrids. This is thought to drastically pre-
vent the phase separation problem and allow for ease of engineer-
ing of the charge trap properties. While ꢀSiꢀOꢀSiꢀ linkages are
17
moiety. It is anticipated that this low-lying LUMO facilitates
electron capture from the Fermi energy level of the adjacent metal
or semiconductor layers. Then, in the case of using materials
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ARTICLE
designed to realize charge trap properties with solution processa-
bility. This was followed by investigation of the charge trap
properties using the MIS device of the polymer, which was
discussed in terms of its electronic structure.
’
EXPERIMENTAL SECTION
2
.1. Theoretical Methods. The geometries of the molecules were
preoptimized using the PM3 semiempirical method in the Hyperchem
.0 package. The ground-state geometries were then fully optimized
using the density functional theory (DFT) at the B3LYP level and
-31G** basis set in Gaussian03.
.2. Materials. All synthetic manipulations were carried out under
8
6
2
an atmosphere of dry argon gas using standard vacuum-line Schlenk
techniques. All solvents, such as diethyl ether, THF, hexane, and toluene
were purchased from Aldrich Chemical Co. Inc. (St Louis, MO, USA)
and distilled from sodium/benzophenone ketyl. All other reagents, such
as diphenylacetylene, lithium wire, silicon tetrachloride, and sulfuric acid
were purchased from Aldrich Chemical Co. Inc. and used without
further purification.
2.3. Synthetic Procedures. Synthesis of Dihydroxy(tetraphenyl)-
silole (2). Synthesis of dihydroxy(tetraphenyl)silole (2) is summarized
schematically in Figure 4. The synthesis of dichloro(tetraphenyl)silole (1) is
previously reported. To synthesize the dihydroxy(tetraphenyl)silole (2),
Figure 4. Schematic of poly(tetraphenyl)silole siloxane synthesis.
19
dichloro(tetraphenyl)silole (1) (2.0 g, 4.5 mmol) was dissolved into 50 mL
of aqueous THF (THF/H
h under an Ar atmosphere. The resulting mixture was then extracted with
ether. The extract was washed with distilled water and brine, dried over
anhydrous Mg SO , and filtered. The filtrate was condensed under reduced
2
O = 4/1) and stirred at room temperature for
containing the silole moiety within the charge trap memory,
substantial enhancement of electron trapping is expected.
In addition to the schematic of the hybridized polymer, poly-
2
2
4
(tetraphenyl)silole siloxane (Figure 3b), the schematic of the
pressure to give pure 2 (1.76 g, 4.2 mmol) in 94% yield as a bright-green
1
tentative one-dimensional potential energy diagram is displayed in
Figure 3c, where electrons inside the film are trapped or transported
through the polymer. The design concept is introduced in this article
for the first time, through comprehensive discussion within a creative
atmosphere between Prof. Jeong and Prof. Sohn, the authors of this
article, who seriously considered the above arguments for the
polyorganosiloxanes and charge trap properties. Shown in Figure 4
is the synthetic scheme of the poly(tetraphenyl)silole siloxane
polymer, which proceeds via synthesis of dichloro(tetraphenyl)silole
solid: H NMR (CDCl ) δ = 6.8ꢀ7.2 (m, 20H), 3.0 (s, 2H) (Supporting
3
Information, Figure S2)
Synthesis of Poly(tetraphenyl)silole siloxane (3). Synthesis of poly-
(tetraphenyl)silole siloxane (3) is summarized schematically in Figure 4.
Dihydroxyl(tetraphenyl)silole (2) (1.0 g, 2.3 mmol), dissolved in 50 mL
of THF, was refluxed with 3.0 mL of sulfuric acid as a catalyst for
polymerization over 2 d. After polymerization, the solvent was dried
under reduced pressure. The residual solid was dissolved in 5.0 mL of
THF and poured into 100 mL of hexane. Polysilole siloxane (3) (0.6 g,
60% yield) was obtained as a light-brown powder after the third cycle of
dissolutionꢀprecipitation.
(
1) and dihydroxy(tetraphenyl)silole (2).
The importance and impact of this work, from the viewpoint
2.4. Analysis of Poly(tetraphenyl)silole Siloxane (3). Gel
of chemistry or materials science, are that advancing considera-
tions of the physical properties or functions required for materials
applications leads to a contrived design of the unique molecular
structure and subsequent first synthesis of the material.
Permeation Chromatography. Molecular weights were measured by gel
permeation chromatography (GPC) using a Perkin-Elmer series 200
equipped with Varian Polymer Laboratories columns (PLgel 5 μm
Mixed-C and Mixed-D) with freshly distilled THF as the eluent.
Molecular weights were calibrated by polystyrene standards (Varian,
Easical PS-1).
As shown in the schematic of Figure 3c, the quantum well
structures with energy barriers of step potential are included
artificially in the hybrid materials for inducing charge trap proper-
ties. Energy barriers originated from the ꢀSiꢀOꢀSiꢀ linkages of
the high band gap, while the midgap-state trapping charges are
created from the HOMO and LUMO energy levels of the π-
conjugated (tetraphenyl)silole rings. Naturally, at this moment, it
is difficult to present a method for theoretical estimation of the
energy barrier and width of the quantum wells for a polysilole
siloxane polymer, although they are pursued through DFT
calculations within this article. Instead, only through successful
synthesis of the materials will it be possible to experimentally
validate the charge trap behavior; capacitanceꢀvoltage (CꢀV)
measurements, a very well-known experimental method in elec-
trical engineering and memory device fields, were taken of the
Nuclear Magnetic Spectroscopy. NMR-grade deuteriochloroform
was stored over 4 Å molecular sieves. All NMR data were collected with
1
Bruker 300 MHz spectrometers (300.1 MHz for H NMR and 75.5 MHz
13
for C NMR) Chemical shifts are reported in parts per million (ppm);
downfield shifts are reported as positive values from tetramethylsilane
(
3
TMS) standard at 0.00 ppm. Samples were dissolved in CDCl unless
13
otherwise stated. C NMR was recorded as proton-decoupled spectra.
Fourier Transform Infrared Spectroscopy. In order to investigate the
molecular structure of the dihydroxy silole and the polysilole siloxane,
Fourier transform infrared spectroscopy (FT-IR) was implemented.
The measurements were conducted on a Nicolet 380 spectrometer
ꢀ1
operated in the mid-IR range of 4000ꢀ400 cm , with all spectra
ꢀ
1
obtained at a spectral resolution of 7.7 cm in the transmittance mode.
Thermogravimetric Analysis (TGA). TGA was performed on a
METTLER TOLEDO SDTA851e. The polysilole siloxane powder
was placed in an aluminum oxide (AlOx) TGA pan. The sample was
1
1,18
metalꢀinsulatorꢀsemiconductor (MIS) device.
In this article, for the first time, this lab synthesized a poly-
tetraphenyl)silole siloxane polymer; its thin films were artificially
(
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ꢀ
1
heated at a rate of 10 °C min from 30 to 800 °C under a nitrogen gas
2.7. Fabrications and Measurements of Device Struc-
tures for Electrical Properties. MIS Structures. Metalꢀinsulatorꢀ
semiconductor (MIS) device structures were fabricated to investigate the
capacitanceꢀvoltage (CꢀV) characteristics. Aluminum metal, as a top
electrode, was thermally deposited onto the polysilole siloxane thin films
ꢀ
1
flow of 50 mL min
.5. Synthesis and Characterization of Poly(tetraphenyl)-
silole Siloxane Thin Films. Spin-Coating Process. A 4 wt % poly-
tetraphenyl)silole siloxane (3) solution in tetrahydrofuran (THF) was
.
2
(
obtained bymixing 0.078 g ofpoly(tetraphenyl)silole siloxane (3) powder
with 1.87 g of THF in an air atmosphere for 5 min in a vial that was
previously cleaned with ethanol and acetone. Then, the 4 wt % solution
was filtered (PTFE, 0.2 um), followed by spin-coating onto a p-type
siliconwaferat2000 rpm for 25s, respectively. The usedSi-waferhadbeen
cleaned with ethanol and acetone prior to use as a substrate. Then, the
spin-coated thin films were heated at 80 °C on a hot plate for 10 min in an
air atmosphere to remove allremaining solvent. Finally, the thinfilmswere
cured in a vacuum tube furnace under a pressure of approximately
where low-doped Si(100) wafers (resistivity: 1ꢀ30 Ω cm, thickness:
3
525 nm) were used as a semiconductor layer. The backside of the Si wafer
in the MIS devices was also coated with Al metal to diminish contact
resistance.
CapacitanceꢀVoltage Measurements. An HP4284 LCR meter was
used to obtain the CꢀV curves by applying an ac voltage with a frequency
of 1 MHz and an amplitude of 1 V to the top aluminum electrode, with a
dc bias voltage swept over a range of ꢀ35 to 35 V on the MIS devices. The
CꢀV curves were recorded in both forward and reverse directions, and
afterward, in order to investigate the charge trap mechanisms.
ꢀ
2
1.0 ꢁ 10 Torr for 1 h at each temperature of 100, 150, 200, and
350 °C. The final poly(tetraphenyl)silole siloxane thin films will be abbre-
viated as polysilole siloxane thin films in the remaining sections of this article.
X-ray Photoelectron Spectroscopy. Compositional changes according
to different curing temperatures were recorded by X-ray photoelectron
spectroscopy (XPS) measurements conducted on a MultiLab 2000 using
a Mg KR (1253.6 eV) source at a pass energy of 20 eV and under a
pressure of 1.0 ꢁ 10 Torr. An Ar ion gun sputtering of 2 min at a
power of 2 kV and 1.3 uA was used for a brief cleaning of the surface of the
samples prior to measurement.
’
RESULTS AND DISCUSSION
This article comprises various data and interpretations in
20
diverse areas: DFT calculations in Gaussian03; polymer synth-
eses; thin film characterizations; photoemission and photoabsorp-
tion spectroscopy; device physics. This comprehensive, yet
somewhat complex composition, proved unavoidable in order to
accomplish the research goal: invention of poly(tetraphenyl)silole
siloxane and realization of negative charge trapping. Its certain
consequence is complexity in interpretation of the data and
randomness in proving the point of this article.
In order to maintain clearness within a proper scientific
context throughout this section, the authors have adopted the
viewpoint of electron theory, whereby the probability amplitude
and energy of electrons are treated as the main factors in
determining the physical properties of a material, especially
electrical and optical properties, which have recently become
very common in chemistry, materials science, and electrical
ꢀ9
þ
Fourier Infrared Spectroscopy. The changes in molecular structures of
the polysilole siloxane thin films with increasing curing temperature were
investigated by FT-IR. The measurements were conducted on a Nicolet
ꢀ
1
3
80 spectrometer operated in the mid-IR range of 4000ꢀ400 cm , with
ꢀ
1
all spectra obtained at a spectral resolution of 7.7 cm
transmittance mode.
in the
Spectroscopic Ellipsometry Measurements. The thicknesses and
refractive indices of the polysilole siloxane thin films with each curing
temperature were measured by spectroscopic ellipsometry (SE) with an
M2000D (J.A. Wollam Co. Inc., Lincoln, NE, USA).
21
2
.6. Surface Science Spectroscopy. Near-Edge X-ray Fine
engineering fields. This framework is based upon quantum
mechanics and solid-state physics. A different aspect of this
article, compared to conventional electron theory, is that the
nature of the electrons in the poly(tetraphenyl)silole siloxane
thin films (abbreviated polysilole siloxane thin films throughout
the rest of the article) is described only in a qualitative manner,
given the theoretical limitations of the organicꢀinorganic hy-
brids. Another important feature of this article is that energy and
location of the electrons are estimated from the molecular
orbitals of the poly(tetraphenyl)silole siloxanes instead of using
the band theory. This is because the polysilole siloxane thin films
can be regarded as molecular solids, of which the building blocks
are molecules or polymers. Within the authors’ knowledge base,
there exists no conclusive or established theory that accounts for
the electron dynamics within molecular solids. Because of this
theoretical limitation, the authors first calculated the energy level
of the molecular orbitals of the poly(tetraphenyl)silole siloxane
and applied its main results toward explaining the charge trap
properties of the polysilole siloxane thin films.
DFT Calculations of Dihydroxy(tetraphenyl)silole and
Poly(tetraphenyl)silole Siloxane. The DFT calculation results
are summarized in Figure 5. Figure 5a shows the calculated orbital
energy levels of dihydroxysilacyclopentane, dihydroxysilacyclo-
pentadiene (dihydroxysilole), and dihydroxy(tetraphenyl)silole
(2), a series of poly(tetraphenyl)silole siloxanes with a number
of tetraphenylsiloles represented by a select number of silicon
atoms (2ꢀ7), and the polydimethylsiloxane within the energy
range of 2 to ∼ꢀ10 eV. The highest occupied molecular orbitals
(HOMOs) and lowest unoccupied molecular orbitals (LUMOs)
Structure Spectroscopy. The near-edge X-ray fine structure spectrosco-
py (NEXAFS) spectra at the C K-edge were measured at the 2B1 PES
beamline of Pohang Accelerator Laboratory (PAL) in the partial electron
yield (PEY) detection mode with a retarding voltage of ꢀ210 V and
an accelerating voltage of 1.6 kV. Molecular bonding information for the
surface of the polysilole siloxane thin films cured at 150 °C, from
poly(tetraphenyl)silole siloxane (3) solutions of 0.1, 0.4, 1, and 4%, was
obtained from the PEY mode NEXAFS spectra by considering probing
depths less than 10 Å. The data were first normalized by the current of
the clean Au mesh to remove the monochromator structure due to
adsorbed carbon on the optical elements of the beamline. Second, the
edge jump was set to 1 for all spectra to remove contributions from the
emission angle, beam decay, and effective spot size on the sample. The
photons showed a polarization of 85% with an incident photon energy
resolution of 350 meV near the carbon K-edge region.
High-Resolution Core-Level and Valence Band Photoemission
Spectroscopy. High-resolution photoemission spectroscopy (HRPES)
was introduced to investigate the chemical structure of the polysilole
siloxane thin films. All HRPES experiments were performed at the 8A2
beamline at the Pohang Accelerator Laboratory (PAL), which is
equipped with an electron analyzer (SES100, Gamma Data Scienta).
Depending on the wt % of the poly(tetraphenyl)silole siloxane solutions,
the O 1s, C 1s, Si 2p, and valence band spectra were measured
accordingly. The total spectral resolution was less than 200 meV at a
photon energy of 635 eV for Si 2p, O 1s, and C 1s, with a photon energy
of 130 eV for Si 2p and a valence band at 8A2 HR-PES beamline in the
Pohang Accelerator Laboratory (PAL). The binding energy of the C 1s
and O 1s were calibrated by measuring the Si 2p core level of the silicon
substrate. All spectra were recorded in the normal emission mode.
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Figure 5. (a) Molecular orbital energy diagram for optimized geometries. Geometry optimization was performed using DFT calculations in the
Gaussian 03 packages. For the poly(tetraphenyl)silole siloxanes, the density of states is increased as n is increased. In order to estimate the barrier height
due to the ꢀSiꢀOꢀSiꢀlinkage, the energies of the molecular orbitals of the polydimethylsiloxane were also obtained and compared. (b) Shapes of
HOMOs and LUMOs of dihydroxysilole. (c) HOMO and LUMO drawings of polysilole siloxanes (from n = 2 to n = 7). (d) Step potential model of
polysilole siloxane (n = 7). The barrier heights were estimated from the LUMO of polydimethylsilone (n = 7) and unoccupied orbitals of
poly(tetraphenyl)silole siloxane (n = 7). The barrier heights for the interchain electron transfer were arbitrarily set, which may be higher for explaining
the energy barrier of the intrachain electron transfer.
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are respectively denoted by the green and red colors. By compar-
ing dihydroxysilacyclopentane and dihydroxysilole, the HOMO
energy becomes higher and the LUMO energy lower through
addition of π bonds into the silacyclopentyl ring. Here, the low
LUMO energy is partly due to the σ*ꢀπ* conjugation between
carefully counted. This is very simply demonstrated using a step
potential model (SPM), which is typically employed for describing
2
3
the electronic states in conjugated hydrocarbon molecules. As
shown in Figure 5d, the simple step potential model for the
poly(tetraphenyl)silole siloxane (n = 7) was built up by using the
energies and orbital shapes of the seven lowest unoccupied orbitals
from LUMO to LUMO þ6. Through the siloxane chain, there are
seven quantum well structures, each possessing different energy
values. The authors holdthatthese quantumwell structuresare the
origin of the electron trapping of the poly(tetraphenyl)silole
siloxane polymers in the condensed phase, which will be applied
toward interpretation of the following capacitanceꢀvoltage
(CꢀV) results in a metalꢀinsulatorꢀsemiconductor (MIS) de-
vice of the polysilole siloxane thin films.
In addition to the intrachain energy barriers, the interchain
energy barrier was also displayed as a dotted line, as shown in
Figure 5d. This is not actually one-dimensional as it surrounds the
seven quantum wells in three-dimensional space through the
polymer chain. The width and height of the interchain energy
barrier relies mainly on the two nearest located (tetraphenyl)silole
rings of the two adjacent polymer chains. The electron transfer
rate between the two adjacent polymer chains, one of the decisive
control factors of the electron trapping, is surely believed to
be more controlled by the interchain energy barriers than by the
intrachain energy barriers.
An important issue to consider is whether the interchain
energy barrier exists. Even if the interchain energy barrier did
not exist, electron trapping would likely be observed due to the
intrachain energy barrier. Nevertheless, the authors speculate that
nonexistence of the interchain energy barrier in such molecular
solids is forbidden because there would only be van der Waals
interactions between the polymer chains instead of direct covalent
bonds. Conversely, the authors would, without reservation, sug-
gest that existence of the intrachain energy barrier strengthens the
effects of the interchain energy barrier on the charge trapping
phenomena in an additional way. This is because, in the case of
having an intrachain energy barrier, the effective cross-section area
of the electron transfer between the two adjacent polymer chains is
smaller than when there is no intrachain energy barrier. It is the
contention of the authors that this is the real meaning of existence
of the intrachain energy barrier for charge trapping events, even in
the presence of an interchain energy barrier. An effective way to
address this issue is to directly calculate the interchain energy
barriers of the poly(tetraphenyl)silole siloxane and poly-
(tetraphenyl)silole using the Marcus theory and compare with
the other. Here, the poly(tetrapahenyl)silole does not include ꢀ
SiꢀOꢀSiꢀ linkages; thus, it lacks any intrachain energy barriers
through the polymer chain. These kinds of theoretical calculations
are being planned with intent for submission in a separate paper.
Undoubtedly, energetic and spatial criteria in this quantum
well structure, such as barrier heights and well widths, are highly
dependent upon optimized polymer geometry. This argument is
very important for interpretation of the effects of the nanoscopic
arrangements of the polymers upon the electrical properties of
the thin films of the poly(tetraphenyl)silole siloxane, as will be
discussed below. It also allowed the theoretical basis that the
physical properties of the polymer are highly sensitive to the
process conditions for depositing the thin films of the polymer
17
the σ* orbital of the silicon and the π* orbital of the butadiene.
This σ*ꢀπ* conjugation can be also understood with the LUMO
electron density isocontour plots of the dihydroxysilacyclopentane
and dihydroxysilole in I-1 and I-3 of Figure 5b. The LUMO of the
dihydroxysilole isprimarily delocalized through π-bonded carbons
and the silicon atom, confirming the σ*ꢀπ* conjugation as
1
7,22
reported.
In addition, the HOMO energy of the dihydrox-
ysilole was higher than that of the dihydroxysilacyclopentane, as
shown in Figure 5a, because the HOMO arises mainly from the π
orbitals of the butadienyl rings as shown in I-4 of Figure 5b,
compared to that of the dihydroxysilacyclopentane from the σ
orbitals (I-2 of Figure 5b).
When the orbital energies of the dihydroxysilole and the
dihydroxy(tetraphenyl)silole (2, n = 1) are compared in
Figure 5a, the HOMO energy increased considerably by addition
of phenyl substituents on the silole ring, while the LUMO energy
only slightly decreased. As shown in I-5 and I-6 of Figure 5b, the
LUMO and HOMO orbitals are spread out mainly on the silole
ring, as well as the phenyl groups.
Variation in the electronic structure of the poly(tetraphenyl)silole
siloxane, according to the number of (tetraphenyl)siloles counted
by the number of silicon atoms (1ꢀ7), was theoretically investi-
gated as well. It is noteworthy that the energies of the HOMO and
LUMO were little affected upon increasing the number of tetra-
phenylsiloles, although many occupied and unoccupied orbitals
were newly generated. As shown in Figure 5c, the HOMOs and
LUMOs for the poly(tetraphenyl)silole siloxanes were similar to
those of the dihydroxy(tetraphenyl)siloles, indicating that the
HOMOs and LUMOs of the (tetraphenyl)siloles in the polymer
do not significantly interact with each other due to the high energy
barrier originating from the SiꢀOꢀSi linkages.
The intrachain energy barrier generated by the σ and σ*
orbitals of the SiꢀOꢀSi linkages in the poly(tetraphenyl)silole
siloxane was indirectly estimated from corresponding orbital
energies of the polydimethylsiloxanes, which have the same
skeleton structure as the poly(tetraphenyl)silole siloxane, as
there are so many energy levels for the polysilole siloxane
polymer that the σ and σ* orbitals were not distinctly found.
As indicated in Figure 5a, the HOMO and LUMO energies for
the dimethylsiloxane were calculated to be ꢀ7.815 and 1.167 eV,
respectively. In addition, the HOMO and LUMO energies for
the polydimethylsiloxane (n = 7) were calculated to be ꢀ7.160
and 0.855 eV, respectively, which are regarded as the energies of
the highest σ and lowest σ* orbitals for poly(tetraphenyl)silole
siloxane (n = 7), as mentioned above. Then, the energy barrier
for an electron occupying the LUMO in poly(tetraphenyl)silole
siloxane (n = 7) was estimated as follows:
energy barrier ðeVÞ
ꢂ
¼
ꢀ
¼
the energy of lowest σ orbitals in polydimethylsiloxane ðn ¼ 7Þ
ꢂ
the energy of π orbitals in polyðtetraphenylÞsilole siloxane ðn ¼ 7Þ
0:855 ꢀ ð ꢀ 1:851Þ ¼ 2:706 ðeVÞ
(
formation of molecular solids).
Synthesis and Characterization of Dihydroxy(tetraphenyl)-
However, the LUMO was localized mainly on the fifth silole ring
on the polymer chain, as shown in Figure 5c,d, indicating that the
energy and spatial variation in the energy barriers have to be
silole and Poly(tetraphenyl)silole Siloxane. The weight-aver-
aged molecular weight (M ) and polydispersity (M /M ) of
w
w
n
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Figure 6. (a) Thermogravimetric analysis results of the poly(tetraphenyl)silole siloxane polymer. The abrupt decrease in weight was observed around
80 °C. (b) Changes in the FT-IR spectra for the poly(tetraphenyl)silole siloxane thin films according to curing temperatures. (c) Change in the peak at
2
ꢀ
1
about 920 cm in the FT-IR spectra, indicating a gradual decrease in intensity of the SiꢀOH groups with an increase in curing temperature.
(d) Changes in the chemical composition of poly(tetraphenyl)silole siloxane thin films, according to curing temperatures, investigated from XPS analysis
results.
polysilole siloxane polymer 3 were measured at approximately
polysilole siloxane thin films. The peaks in the range of
1
ꢀ
3
087 and 1.18, respectively, using GPC (Supporting Information,
1600ꢀ1400 cm were ascribed to the stretching of CdC in
24
Figure S1). Since the molecular weight of the dihydroxy(tetra-
phenyl)silole monomer (2) is 418.6, the poly(tetraphenyl)silole
siloxane (3) consists of nearly seven monomers on average. As such,
the lab focused upon the polymer of seven monomers in Figure 3b
and c and Figure 5a,c,d, and the related results and discussion. The
synthesis of 3 was confirmed by NMR (Supporting Information,
Figure S2).
the phenyl groups, as well as silole rings.
The inorganic sections (SiꢀOꢀSi linkages) were also con-
ꢀ
1
firmed by the peaks at 1200ꢀ1000 cm , attributed to their
ꢀ
1
antisymmetric stretching, and the peak at 800 cm , assigned to
25
their symmetric stretching. The position of the SiꢀOꢀSi anti-
symmetric stretching peak was highly dependent upon the local
bonding geometry of the SiꢀOꢀSi linkages. The peak position is
ꢀ
1
The TGA results of the poly(tetraphenyl)silole siloxane (3)
polymer are shown in Figure 6a. Weight loss began at 90 °C by
evaporation of physisorbed water molecules, with additional
weight loss from approximately 170 °C related to the condensa-
tion reaction of the SiꢀOH groups. The possibility of the
condensation reaction is also supported by the following FT-IR
results of the polysilole siloxane thin films. Above 280 °C, a huge
loss in weight was observed, due to thermal decomposition of the
organic sections, as also mentioned in the below FT-IR data.
Synthesis and Characterization of Polysilole Siloxane Thin
Films. Figure 6b shows the FT-IR spectra of the poly-
maintained at approximately 1100 cm during the curing tem-
peratures, up to 200 °C, which is evidently due to the SiꢀOꢀSi.
Moreover, the terminal SiꢀOH groups were detected at about
ꢀ
1 26
920 cm
,
the intensity of which gradually decreased with
increasing curing temperature, as shown in Figure 6c. As the
curing temperature increased, the peaks indicating the organic
groups gradually decreased, but only slightly. Above 200 °C,
decomposition of the organic species accelerated, consistent with
the TGA results in Figure 6a and the XPS results in Figure 6c.
Finally, the organic species disappeared to a considerable extent
at curing temperatures of 350 °C. Figure 6d shows variations in
the chemical composition of the polysilole siloxane thin films
according to the curing temperature.
This lab measured refractive index using spectroscopic ellip-
sometry according to the curing temperatures shown in
Figure 7a,b. The refractive indices of the thin films were
abruptly increased as the curing temperature increased from 80°
(
tetraphenyl)silole siloxane thin films (abbreviated polysilole
siloxane thin films) at different curing temperatures. The peaks
at 3080, 3060, and 3020 cm are attributed to the stretching of
ꢀ
1
2
ꢀ1
the sp -hybridized CꢀH, while the peaks at 760 and 700 cm
were assigned to the bending of CꢀH in the monosubstituted
2
4
aromatic groups, implying the existence of phenyl groups in the
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Figure 7. Spectroscopic ellipsometry results. (a) Refractive indices of
poly(tetraphenyl)silole siloxane thin films. The abrupt increase in
refractive index was observed as the curing temperature increased from
8
0 to 100 °C. At higher curing temperatures, the refractive indices
gradually decreased with increasing curing temperatures. (b) Changes in
refractive indices and film thicknesses according to curing temperature.
(
as prepared) to 100 °C, while decreasing slightly at higher
temperatures. The refractive index was closely related to the
density of the thin films, which was supported by the following
6,8
Figure 8. (a) Tentative model for nanoscopic arrangements of constitut-
ing poly(tetraphenyl)silole siloxane polymer chains inside the dielectric
thin film. The geometry of the individual polymer chain was initially
optimized using DFT calculations in Gaussian03 packages. Here, the
number of silole rings was set to 7. At low curing temperatures (e80 °C),
the kinetically preferred state is sustained, where the silole rings stays in a
dominant state of “restricted intermolecular rotation”, while, at higher
curing temperatures (g100 °C), the thermodynamically preferred state
emerges where cofacial πꢀπ interactions in a skewed manner between
phenyl rings bonded to the silole rings must exist. (b) Total energy graph
along the interchain distance (s). The origin of the total energy of the
polymer was divided into the two main components, the πꢀπ inter-
molecular interaction betweenphenyl rings of the neighboring siloles in the
polymers, as mentioned above, and total conformational strains of the
polymer, such asangle strains or geometry relaxation of the polymer chains.
LorenzꢀLorentz equation, as briefly explained in the Intro-
duction and Figure 1b:
!
2
n ꢀ 1
∑
Nj RjðelectronicÞ
3
¼
ð2Þ
2
n þ 2
3ε0
where n is the refractive index, N density, R
electronic
j(electronic)
polarizability, and ε vacuum permittivity. As explained in detail
j
0
in the Supporting Information (Figure S3), the refractive index
(
n) is directly proportional to N and Rj(electronic), in the range of
j
the refractive index of the polysilole siloxane thin films.
The abrupt increase in density of the thin film cured at 100 °C
is explained by a chemical kinetics model for nanoscopic
arrangement and interchain interaction of the constituting poly-
silole siloxane polymers in the film, as summarized in Figure 8a,b.
As the polymer solution was spin-coated at room temperature,
the coating solvent evaporated too quickly compared to the time
scale required for the polymers to be arranged on the substrate,
temperature gave rise to huge differences in the refractive index
values, as shown in Figure 7a,b, indicating the significant increase
in density of the films. This can be understood by creating the
nanoscopic model of collective interactions of the polysilole
siloxanes, thermodynamic preferred state, as depicted in Figures 8a
and 13f, which is based upon the unique molecular structure of
the polysilole siloxane.
The total energy graph in Figure 8b is qualitatively drawn to
explain the abrupt increase in the density of the films at a curing
temperature of 100 °C. This is accomplished by introducing two
models, the kinetically preferred state (R) for the film cured at
27
leaving a metastable state of the polymer arrangement. This is
considered to be a kinetically preferred state in terms of chemical
kinetics, as illustrated in Figures 8a and 13f. A curing temperature
of 80 °C for the as-prepared sample is thought to be only for
removal of the remaining solvent molecules trapped between the
polymer chains and subsequently too low to irritate the kinetically
preferred state. Surprisingly, the small increase (20 °C) in curing
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0 °C and the thermodynamically preferred state (β) for the film
ARTICLE
8
energy of R and β states, respectively:
cured at 100 °C, both related to the nanoscopic arrangement of
the polymer, all of which are summarized in Figure 8a. The
authors assert that the abrupt increase in the density of the films is
related to the πꢀπ intermolecular interactions between the
phenyl rings of the neighboring silole groups. A theoretical
evaluation of the intermolecular interaction between benzene
molecules with benzene dimers of various structures, such as
parallel, T-shaped, and slipped parallel cases, according to the
intermolecular (benzene to benzene) distance, has been per-
!
½Rꢃ
ER=Eβ
ꢄ exp ꢀ
½βꢃ
kT
This implies that the R state does not disappear entirely, even at
100 °C
The authors hold that the kinetically preferred state resembles the
nanoaggregation state in terms of local polymer arrangement, which
29
promotes aggregation-induced emissions (AIEs). For silole-related
molecules showing the AIE, restricted intermolecular interaction
between peripheral phenyl rings perpendicular to the silole rings that
prevent nonradiative energy dissipation has been proposed. On the
contrary, in the thermodynamically preferred state, another different
model for the interchain interactions has to be proposed. The model
for the thermodynamically preferred state has to account for the volume
shrinkage of the polysilole siloxane thin films, which is evident from
the increase in the film density estimated in SE. This is schematically
shown in Figures 8 and 13f, where the cofacial πꢀπ interaction, in a
skewed manner, is enhanced with the phenyl groups perpendicular to
the silole rings. The higher potential energy, accompanied with
distortion of the conformational structure of the polymers, is
compensated by the energy release from the πꢀπ interactions
between the phenyl groups of the silole ring of the two adjacent
polymer chains. The distance between the phenyl rings that can
interact with each other in the thermodynamically preferred state is
much shorter than that in the kinetically preferred state, which
eventually affects the interchain electron transfer; this will be discussed
in the capacitanceꢀvoltage (CꢀV) results below.
28
formed as model chemistry by Tanabe et al. Though calcula-
tion of the intermolecular interaction between benzene
molecules in a skewed manner, as found in our polysilole siloxane
case, was not performed in their study, we proffer that their
viewpoint regarding the intermolecular interaction between two
benzene molecules can be applied toward our interpretation.
The total energy graph is drawn along the interchain distance
(
s) between the phenyl rings of the neighboring silole groups,
defined in Figure 8b as the x-coordinate. The authors speculate
thattheorigin of the totalenergy of the polymer isdividedinto two
main components, the πꢀπ intermolecular interaction between
the phenyl rings of the neighboring siloles in the polymers, as
mentioned above, and the total conformational strains of the
polymer, such as angle strains or geometry relaxation of the
polymer chains. Herein, the πꢀπ intermolecular interaction
2
8
energy drawn (E (s)) was previously cited and fitted by
πꢀπ
eq A.1; the total conformational strain energy (Estrain(s)) was
hypothetically drawn without direct evidence, as fitted by eq A.2:
"
#
ꢀ
ꢁ
ꢀ
ꢁ
n
m
s0
ðπ ꢀ πÞ
s
s
0ðπ ꢀ πÞ
This thermodynamically preferred state has not been proposed in
previous studies of silole-related molecules. This previous rareness
is thought to be related to the sample preparation conditions in
those studies, such that even thin films were formed by simple
coating and solvent evaporation; curing conditions above 100 °C
were not applied in those studies. It has been common knowledge
that, for spin-on glass, tensile stress is enhanced after film curing.
The induced tensile stress provokes some extension of the
polymer chains in the direction parallel to substrate surface,
enhancing the interchain interactions in the direction perpendi-
cular to the substrate surface. This isa drivingforce to overcome an
energy barrier to the thermodynamically preferred state.
The authors can also account for the changes in refractive
index and thickness as the temperature increased above 100 °C,
using the above TGA and XPS results.
(i) 200ꢀ350 °C: Decomposition of the organic species is
accelerated, consistent with the TGA results in Figure 6a
and the XPS results in Figure 6d. This provides the
decrease in the effective volume of the polymer chains,
decreasing film thickness. Furthermore, the gradual dis-
appearance of highly polarizable organic groups decreases
the refractive index constantly.
(ii) 100ꢀ200 °C: The additional weight loss shown in the
TGA results (Figure 6a) is related to the condensation
reaction of theSiꢀOHgroups. Thisprovides a decrease in
the effective volume of the polymer chains, decreasing the
film thickness. This is thought to increase the refractive
index of the films within this temperature range, which is
contradicted with the SE results showing its slight de-
crease (Figure 7b). The authors contend that the increase
in therefractive index, owingtothe condensation reaction,
is negated by the decrease in the refractive index due to
slight decomposition of the organic groups, giving rise to a
E
π ꢀ πðsÞ ¼ A 3
ꢀ
, ðn > mÞ
ðA.1Þ
s
B
EstrainðsÞ ¼
ðA.2Þ
1
þ exp½βðs ꢀ s0ðstrainÞÞꢃ
where s0
and s0(strain) are the specific interchain distances where
(πꢀπ)
the total πꢀπ interaction energy between the phenyl rings and total
conformational strain energy of the polymer become minimum and
half of maximum, respectively, while n and m are integers, β is a
coefficient, and A and B are arbitrary constants. All of the factors are
arbitrarily chosen to suitably describe the speculated physicochemical
situation. This lab hypothesizes that the conformational strain energy of
the polymers is nearly zero when the distance between the phenyl rings
of the neighboring siloles of the polymers is far enough, then it increases
when the interchain distance (s) decreases below a point. After that, the
conformational strain energy is assumed to be constant. Thus, as the
polymer solution was spin-coated onto the substrate at room tem-
perature, the polymer films might remain in the R state (kinetically
preferred state) due to the energy barrier originating from the increase in
the conformational strain energy near s0(strain). As the curing tempera-
ture was increased from 80 and to 100 °C, the polymer chain could
overcome the energy barrier and enter into the β state (ther-
modynamically preferred state), decreasing the interchain distance (s).
The authors hold that the two local minima still exist over
1
00 °C, even though the β state becomes more dominant in the
films as curing temperature increases. With [R] and [β] as the
respective concentrations of finding the R and β states in the
polymer film, under thermodynamic equilibrium conditions
more preferred at 100 °C than 80 °C, the ratio of the probabilities
are given by the following equation, where E and E are total
R
β
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ARTICLE
Figure 9. (a) Near-edge X-ray fine structure spectroscopy. Especially, for the poly(tetraphenyl)silole siloxane thin films fabricated from 4 wt %
solution, the peak due to excitation from the C1s to LUMO level appeared at 285.6 eV. (b) Core level photoemission spectra. In the Si 2p
spectra, the binding energy of the Si 2p electron in the poly(tetraphenyl)silole siloxane thin films was determined to be 102.8 eV by performing
the charge calibration procedure, positioning the binding energy of the Si 2p electron of the silicon substrate to 99.3 eV. For the
poly(tetraphenyl)silole siloxane thin films fabricated from 4 wt % solution, the binding energy of the C 1s electron in the poly-
(tetraphenyl)silole siloxane thin films was determined to be 284.6 eV. (c) Si 2p spectra for the photon energy of 130 eV. (d) Valence band
photoemission spectra in the case of photon energy of 130 eV. (e) Comparison of the valence band photoemission spectrum for the 4 wt % case
with the simulated spectrum (n = 1ꢀ7). (f) Comparison of the photoemission and the NEXAFS spectra with the simulated spectrum for
occupied and unoccupied orbitals (n = 7).
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ARTICLE
peak at 99.9 eV might arise from a Si atom at an imperfect
oxidation state of þ1, in the shallow native-oxide layer between
32
the polysilole siloxane film and Si wafer. In addition, the Si 2p
peak at 103.4 eV from the Si atom, with an oxidation state of þ4
in the native oxide, might be overlapped and thus hidden by the
strong peak at 102.8 eV from the polysilole siloxane thin films.
Considering the same chemical environment of the Si atoms of
the other three films, the Si 2p peaks for 0.4, 1, and 4 wt %
precursor solutions must also be positioned at 102.8 eV. After
applying the binding energy shift values used for the charge
calibrations into the C 1s and O1s spectra, this lab also obtained
the charge-calibrated ones, as shown in Figure 9b.
Figure 10. Energy diagram of molecular orbitals of the poly-
(
tetraphenyl)silole siloxane with respect to the Fermi energy level (E
F
).
It is possible to discuss the oxidation state of Si atoms from binding
energy values. The binding energy for the Si atoms of an oxidation
state of þ2 is 102.1 eV, while that for the oxidation state of þ3is102.8
The relative position of the E was obtained by subtracting the C 1s peak
F
energy (284.6 eV) in the photoemission spectra from the first absorption
peak energy (285.6 eV) in the NEXAFS, from the poly(tetraphenyl)silole
siloxane thin films fabricated from the 4 wt % solution. The LUMO levels
were drawn from the DFT calculation results. The HOMO levels were
also estimated from DFT calculation results.
33
eV. It is interesting that the binding energy for the Si atom of the
polysilole siloxane thin films is the same as that for the þ3 oxidation
state, since the formal charge of the silicon atom is þ2. The authors
speculate that this is due to the high electron affinity of the
slight decrease in the refractive index. Accordingly, the
authors tentatively assume there is a little decomposition
of the silole rings, even though this is not evidently
detected in the TGA results (Figure 6a).
(
tetraphenyl)silole ring, decreasing the electron density of the
adjacent silicon atom. In a materials science sense, electron affinity
χ) is defined as the energy required to remove an electron from the
(
bottom of the conduction (E ), whereby χ = E ꢀ E , in which E is
C
0
C
0
18a
NEXAFS and HRPES Studies. NEXAFS is used to monitor the
electronic transitions from the core level of a specific atomic
species of a molecule to its unoccupied molecular orbitals (π* and
the vacuum energy level. The E of the polysilole siloxane thin films
C
originates from the LUMO energy of the polysilole siloxane, which is
discussed in the above DFT Calculations. The high electron affinity of
the (tetraphenyl)silole ring, owing to the low-lying LUMO, is
favorable for capturing electrons from the Fermi energy level of
adjacent metal or semiconductor layers when the polysilole siloxane
thin film is inserted between the metal and semiconductor layers. For
the polysilole siloxane, due to the intrachain energy barrier from the ꢀ
SiꢀOꢀSiꢀ bond, the electron transferred from the Fermi energy
level of the metal or semiconductor layer is trapped in the LUMO
energy level of the silole group.
The important aspect of the C1s spectra is the binding energy
value itself for the film fabricated from the 4 wt % precursor
solution, 284.6 eV. This is because it has to be compared to the
NEXAFS data, as discussed above, in order to provide the
physicochemical origin for the negative charge trap properties of
polysilole siloxane thin filmsspin-coatedfrom the4 wt% precursor
solution. In the presented system, the high electron affinity is
crucial in explaining the charge-trapping properties. The C 1s
binding energy increased at the higher electron deficiency of the C
atom, while the LUMO energy can be also affected by electron
deficiency. Then, binding energy in XPS or charge trapping in the
following capacitanceꢀvoltage (CꢀV) measurement could po-
tentially be affected by electron deficiency or electron affinity. The
authors surmise that the C 1s binding energy in the polysilole
siloxane thin films is related to the charge trap properties.
3
0
σ* orbitals). The transition probability of these excitations are
dictated by dipole and symmetry selection rules. In this article, the
focus was on determination of the LUMO energy level, which was
very crucial in explaining the electron trapping mechanism in
the MOS device, and will be discussed in the following subsection.
The NEXAFS spectrum of thepolysilolesiloxanethin film nearthe
carbon K-edge is shown in Figure 9a. For the polysilole siloxane
thin film spin-coated from the 4 wt % precursor solution, a narrow
peak was positioned at 285.6 eV, which was overlaid by a preband
at the low-energy side and a more complex structure at higher
ranges. The narrow peak was assigned to a electronic transition
from the C 1s core level to the LUMO of the molecule, whose
molecular orbital shape is modeled like II-11, as shown in
Figure 5c. However, as discussed below in the HRPES results
for core levels, the binding energy of the C 1s core level for the film
fabricated from the 4 wt % precursor solution was found to be
2
84.6 eV, the energy difference of a C 1s core level and Fermi
energy level (E ). Since the narrow peak of 285.6 eV in the
F
NEXAFS corresponds to an electronic transition from the C 1s
corelevel to the LUMO, the energy of theLUMOis determined to
be 1.0 eV above with respect to the E , as depicted in Figure 10.
F
In the core-level photoemission spectroscopy of the polysilole
siloxane thin films using a photon energy of 630 eV (Figure 9b),
the binding energies of the electrons in the Si 2p and C 1s orbitals
were investigated. The first obstacle confronted was the charging
problem that usually occurs in photoemission studies of an
insulator. Then, the binding energy of the Si 2p electron was
calibrated to a known value. As shown in the Si 2p spectrum in
Figure 9b, the polysilole siloxane thin films fabricated from 0.1 wt
As shown in Figure 9c, the binding energy of the Si 2p peak, in the
case of 130 eV photon energy, was also charge calibrated to
102.8 eV. This calibration allows the valence level spectra obtained
for the 130 eV photon energy to be charge calibrated to Figure 9d.
Valence band photoemission spectroscopy was used to investigate
the valence band of the molecular orbitals of the molecular solids. A
broad and strong peak appeared near 6 eV, with a weak and broad
peak near 16 eV. In order to investigate the origin of the valence
spectrum, this lab simulated the photoemission spectra of the
polysilole siloxane (n = 1ꢀ7) from the DFT calculation results
%
precursor solution showed two prominent peaks, where the
former peak was due to the Si atoms in the thin films and the
latter to the Si atoms of the silicon substrate, owing to its
relatively low thickness. Because the binding energy of the Si
31
34
2
p electrons from the silicon substrate -bulk silicon- is 99.3 eV,
using Gausssum software. Then, the valence level spectrum for the
the latter peak was calibrated to 99.3 eV and, hence, positioned
the former peak to 102.8 eV without further consideration. The
4 wt % case was compared with the simulated spectra, as shown in
Figure 9e. Here, all of the simulated spectra (n = 1ꢀ7) were
7
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Figure 11. (a) Schematic of the vertical structure of the MIS device. (b) Energy band diagram for three typical cases. (c) Capacitanceꢀvoltage (CꢀV)
characteristics for the poly(tetraphenyl)silole siloxane thin films according to the curing temperature.
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Table 1. Thicknesses and Dielectric Constants of the Poly(tetraphenyl)silole Siloxane Thin Films (abbreviated polysilole siloxane
a
thin films) Cured at 80, 100, 150, 200, and 350 °C
reverse sweep
|F|(C cm
forward sweep
|F|(C cm )
ꢀ3
ꢀ3
curing temp. (°C)
thickness (nm)
dielectric constant
ΔVFB,2(V)
)
ΔVFB,2(V)
3
3
8
0
351
290
264
205
107
3.00
3.02
3.28
3.32
3.39
>35.8
15.9
9.7
>0.015
0.010
0.008
0.028
0.052
>35.8
14.4
3.5
>0.015
0.009
0.003
0.011
0.016
1
1
2
3
00
50
00
50
19.7
9.9
7.7
3.0
a
Flat band shift (ΔVFB,2) due to the charges trapped over the film thickness and trap densities of negative charges (|F|) for reverse sweep (starting from
positive bias conditions) and forward sweep (starting from negative bias conditions) directions were also summarized.
calibrated by positioning the LUMO to 1.0 eV, as mentioned in the
NEXAFS results. For instance, since the band gap value
it are dependent on the frequency and amplitude of the applied
voltage and doping concentration of the silicon wafers, it remains
difficult to determine ddep exactly. Therefore, the capacitance in
theaccumulation regionfrom eq 3 isusedto calculate the dielectric
permittivity in the dielectric thin films, which are shown as the flat
region in the CꢀV graphs.
(
HOMOꢀLUMO gap) in the polysilole siloxane (n = 7) was
estimated at 3.193 eV from the DFT calculation, the binding energy
scale in the simulated valence band spectrum was calibrated to
position the HOMO level to ꢀ2.193 eV, if the LUMO level was set
to 1.0 eV above the E . Similar to the valence level spectrum for the 4
Further important information obtained from the CꢀV curves
F
wt % case, the simulated spectra showed two peaks at 7 and 17 eV,
with a 10 eV difference, while the shapes of the simulated photo-
emission spectra were nearly identical regardless of the number of
is the flat band shift (ΔV ). In the MIS structures, the band
FB
structure of the silicon substrate very close to the insulator is not
flat, but ratherbentdown atthe p-type semiconductor, asshown in
the case of Φ < Φ (unbiased) in Figure 11b; this is known as
(tetraphenyl)silole groups (n). In Figure 9f, the photoemission and
M
SC
NEXAFS spectra of the 4 wt % polysilole siloxane thin film are
compared with the simulated spectra for occupied and unoccupied
orbitals (n = 7). The positioning of relatively low unoccupied and
high occupied orbitals is summarized in Figure 10.
band bending. When a negative bias voltage is applied at the top
electrode, the band structures at the regions become flat, as shown
inthe caseof Φ < Φ (flatbandcondition) inFigure11b. Inthe
M
SC
MIS structure containing an aluminum electrode and p-type
Charge Trap Properties of Polysilole Siloxane Thin Films.
To investigate the dielectric and charge trap properties of the
polysilole siloxane thin films, capacitanceꢀvoltage (CꢀV) curves
were obtained in typical MIS devices at which thermally depos-
ited Al and low-doped p-type Si wafers were used as the metal
and semiconductor, respectively; this is schematically illustrated
in Figure 11a. A bias voltage was applied to the aluminum top
electrode with respect to the silicon wafer in the MIS structures.
Under the negative bias voltage at the aluminum electrode, the
holes that were major carriers in the p-type silicon wafer tended
to move toward the interface between the insulator layer and the
silicon wafer. This condition is called the “accumulation region”
in the typical capacitanceꢀvoltage (CꢀV) graphs, Figure 11b.
Here, the thickness of the insulator is directly used for the
thickness value (d), and for calculating the dielectric permittivity
semiconductor, ΔV , specifically ΔVFB,1, is estimated to be
FB
0.8 eV, as the Fermi energy of 4.1 eV of aluminum (Φ ) is lower
M
35
than 4.9 eV, that of the semiconductor (ΦSC).
However, the presence of some charges, positive or negative,
in the insulator can also change the extent of band bending. As a
result, in order to render the band flat, the voltage applied at the
aluminum electrode has to be different, indicating that the ΔVFB
value changed. Generally, the presence of positive charges shifts
the ΔVFB in the negative direction, while the presence of negative
charges shifts the ΔVFB in the positive direction.
The CꢀV curves for the polysilole siloxane thin films, accord-
ing to the curing temperature, are shown in Figure 11c. With the
capacitance in the accumulation region of the CꢀV curves and
the thickness from the SE results, dielectric constants were
calculated using eq 3, as summarized in Table. 1.
(
ε), as shown in the following equation:
It is very reasonable that, in all samples, there were only
positive ΔVFB values, indicating the trapping of negative charges
(electrons). This is because the poly(tetraphenyl)silole siloxane
dielectrics have LUMO energy levels of high electron affinity, as
explained in the Introduction and the NEXAFS results. Even in
the as-prepared sample (80 °C) showing only a flat curve, it is
A
ε0 ¼ 8:854 ꢁ 10ꢀ F=m
12
C ¼ εε0
ð3Þ
d
Conversely, in the positive bias voltage at the aluminum electrode,
the holes are subjected to move down to the bulk of the silicon
wafer, and the depletion layer, where the holes are deficient, is
newly formed at the interface between the insulator layer and
silicon wafer. This conditionisreferred toas the “depletion region”
in the CꢀV graph, as shown in Figure 11b. The thickness of the
evident that its ΔV
(total flat band shift) is greater than 35
FB, tot
V. This is because their capacitance value was surely due to the
accumulation region, as compared to that for 100 °C. From the
ΔV
values of these graphs, the volume densities of trapped
FB,tot
18b
depletion layer (d ) has to be added to calculate the dielectric
charges (F; charge trap density) are calculated using eqs 5ꢀ9:
dep
permittivity, as shown in the following equation:
The ΔVFB,tot is divided into four specific terms:
A
ε0 ¼ 8:854 ꢁ 10ꢀ12 F=m
ΔVFB;tot ¼ ΔVFB;1 þ ΔVFB;2 þ ΔVFB;3 þ ΔVFB;4
ð5Þ
C ¼ εε0
ð4Þ
d þ ddep
where ΔVFB,1 is due to the work function difference of the
aluminum and silicon substrates, which was estimated to be ꢀ0.8
Although the thickness of the depletion layer and the concentra-
tion of the minor carriers (electron in p-type silicon wafer) within
eV, as explained above. The value ΔV
is due to the charges
FB,2
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Figure 12. Representation of the step potential model in the energy diagrams of the MIS device, explaining the CꢀV hysteresis phenomena. At
a negative bias, the electron is injected from the aluminum electrode into the poly(tetraphenyl)silole siloxane thin film, inducing a positive flat
band shift (ΔVFB). At a positive bias condition, the electron is injected from the silicon substrate, inducing a more positive flat band shift
(
ΔVFB).
trapped over the film thickness, ΔVFB,3 and ΔVFB,4 are due to
structural defects in the silicon during oxidation and interface-
trapped charges, respectively, and ΔVFB,3 and ΔVFB,4 are tenta-
tively assumed to be very small. Thus:
Here, F(x) is also assumed to be uniform according to x, along the
film depth, F(x) = F:
Z
t
F
F
1
t2
ΔVFB;2 ¼ ꢀ
xdx ¼ ꢀ
ð8Þ
ε0 εOX
ε0 εOX 2
3
0
3
ΔVFB;2 ¼ ΔVFB;tot þ 0:8 V
ð6Þ
Thus, the trap density of negative charges (|F|) is estimated from
ΔV
The relationship between ΔVFB,2 and the negative charge trap density
:
FB,2
(
F) has to be considered as follows, where t is film thickness:
2
ε0 εOX
3
Z
jFj ¼
ðΔVFB;2
Þ
ð9Þ
t
2
FðxÞ x dx
t
3
3
ΔVFB;2 ¼ ꢀ
ð7Þ
ε0 εOX
0
3
In both cases of reverse and forward sweep directions, the
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estimated charge trap densities for the negative charges are
summarized in Table. 1. In the low curing temperature ranges
of up to 150 °C, the trap density of negative charges (|F|)
decreased as the curing temperature increased. The authors
hold that this is related to the changes in the nanoscopic
arrangement of the polymer chains according to curing
temperature. As illustrated in Figure 8a, the thermodynami-
cally preferred state becomes dominant at relatively higher
curing temperatures, within a range of up to 150 °C, where
the cross section of the cofacial πꢀπ interaction in the
nearest phenyl rings is greatly enhanced. In this nanoscopic
arrangement, the interchain electron transfer rate increased,
provoking a decrease in the trap density of the negative
charges (|F|), which will be qualitatively explained in the
following subsection.
Within high curing temperature ranges above 150 °C, the
trap density of negative charges (|F|) increased as the curing
temperature increased, as shown in Table 1. This is thought to
be related to the formation of a new charge trap center due to
the thermal decomposition of the (tetraphenyl)silole rings in
the polymer chains, mentioned in the above TGA and FT-IR
results.
of the polymer through a tunneling mechanism. These
electrons are also transported to unoccupied molecular
orbitals or vibronically excited states of the LUMO level
of the adjacent polymer through a resonant-tunneling me-
chanism, which are electronically or vibrationally de-excited
into the LUMO level within a very short time. The inter-
chain electron transfer continues in order for the electrons
to reach the polymer next to the aluminum metal, where the
electron is transported into the Fermi energy level of the
aluminum.
The origin of the anticlockwise CꢀV hysteresis over
the entire curing temperature range is related to why
the electron trapping in the reverse sweep condition
(positive bias condition, 4) is higher than that in the
forward sweep condition (negative bias condition, 2, which
is confirmed by comparison of the trap density of negative
charges (|F|) in Table 1. The authors propose that this is
attributed to formation of an additional conducting channel
between the polymer and native SiO of approximately
2
15 Å on the silicon substrate. The SiꢀOꢀSi linkage is
formed from the terminal SiꢀOH of the polymer chain;
the SiꢀOH’s existed on the native oxide. Tunneling through
the σ bond is known to be much faster than through
The appearance of the positive flat band shift is more
succinctly explained by modifying the step potential model
in Figure 5d and applying the molecular orbital energy
diagram discussed in Figure 10 toward the electronic energy
diagram of the MIS structure, as shown in Figure 12. This
model is to be restrictively applied toward only the samples
cured at low temperature ranges (e150 °C), with no thermal
decomposition of the (tetraphenyl)silole rings. Here, the
energy barriers of the interchain electron transfer were
displayed, instead of those of the intrachain electron transfer.
For the simplification, electron energy levels displayed with-
in energy barriers were taken from only few energy levels of
HOMO, HOMO þ 1, LUMO, and LUMO þ 1. The relative
3
6,37
space.
The higher |F| values for the reverse sweep
conditions, compared with forward sweep conditions, is
due to the fast electron pumping from the silicon substrate
to the polymer dielectrics thin film via electron tunneling
through the σ bond (Supporting Information, Remarks S1).
This is why a relatively low-energy barrier at the interface of
the polymer dielectric and silicon substrate is displayed in
the electronic energy diagram of the MIS structure, as shown
in Figure 12. This trend is much more strengthened as the
curing temperature is increased. This is certainly because
the SiꢀOH condensation reaction forming the interfacial
SiꢀOꢀSi linkages is more promoted at higher curing
temperatures.
proximity of the LUMOs to E in terms of electronic energy,
F
compared to HOMOs, resulted in the phenomenological
situation whereby the electron is the owner of the charge
trapping of the polysilole siloxane dielectric thin films, rather
than the hole.
Schematic Demonstration for the Negative Charge
Trapping in the Poly Tetraphenylsilole Siloxane for Sam-
ples Cured at Low Temperature Ranges (e150 °C). This
interpretation mainly focuses on the explanation of the ΔVFB,2
and |F| values in the forward sweep direction for the samples
cured at low temperature ranges (e150 °C), whereby the
electron trapping and electron transfer phenomena were
found to be more closely related. This also enabled the authors
to recognize that investigation of the electron tunneling
process between the adjacent π organic groups was very
important for understanding the electron transfer kinetics of
polymers having quantum well structures. The similar de-
monstration for the reverse sweep direction was summarized
separately (Supporting Information Figure S4), clarifying the
interfacial SiꢀOꢀSi linkages to explain the anticlockwise
CꢀV hysteresis.
At negative bias condition 2, electrons from the metal
electrodes are injected into the LUMO level of the polymer
through a tunneling mechanism. These electrons are trans-
ported to unoccupied molecular orbitals or vibronically
excited states of the LUMO level of another adjacent
polymer through a resonant-tunneling mechanism, which
are electronically or vibrationally de-excited into the LUMO
level within a very short time. The interchain electron
transfer enables the electrons to reach the polymer next to
the silicon substrate, where the electron is transported into
the conduction band of the p-type silicon. This progressing
electron transport through the poly(tetraphenyl)silole si-
loxane dielectrics maintains the steady-state condition of the
charge density within the unit volume. When the negative
bias condition is turned off, some of the electrons are stored
in the dielectrics, noted as electron trapping, generating
the positive flat band shift, as illustrated at 3 and 4 in
Figure 12.
In the interpretation of CꢀV results of the above subsec-
tion, the charge density (F) was held constant over the film
thickness. This was based on the following evident arguments for
the kinetics regarding electron transfer and trapping. The simple
model of the kinetics is proposed in Figure 13a. Here, the
dielectrics layer is treated with a series of volume elements, the
number of which is n, and whose area is A, and thickness is
approximately t/n. The values F , F , and F respectively denote
At positive bias condition 4, the electron transport and
trapping proceed in a reversed manner. The electrons in the
doping level and the topmost energy levels of the valence
band in the depletion region are injected into the LUMO level
f,n
m
s
the negative charge trap density for each element at negative
bias condition, metal, semiconductor, while k , k , k , ..., k
,
nꢀ1
m
1
2
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Figure 13. Origin of interchain electron transfer. (a) An electron transfer kinetics model at a negative bias voltage for the forward
sweep direction (b) Geometric details for describing the cofacial interaction in a skewed manner of two phenyl rings at the interfaces of
the adjacent polymer chains (c) The model of the intermolecular interaction considering van der Waals thickness of the phenyl rings
(
d) Logarithmic relation of ΔVFB and thickness (e) The optimized geometry of the polymer (n = 7) obtained using DFT calculations
with Gaussian03 were applied for packing the volume between the aluminum and silicon parts. (f) The changes in the interchain distance
s) and the distance of the phenyl rings (d) at the interface of the polymer chains for the kinetically preferred state and thermodynamically
preferred state.
(
and k are the rate constants for electron transfer in the
direction for the aluminum to silicon substrate under negative
bias conditions.
Forward Sweep Direction Case. Under negative bias con-
ditions, for the first element, the steady-state approxima-
tion for the charge trap density of the first element (Ff,1)
ꢀs
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Journal of the American Chemical Society
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is applied:
The rate constant for the electron transfer inside the polysilole
siloxane dielectrics (k ) is surely thought to originate from the
1
km
k1
interchain electron transfer rate of the poly(tetraphenyl)silole
siloxane polymer chains. This interchain electron transfer occurs
thorough electron tunneling through the two nearest
(tetraphenyl)silole rings, which is classified into as electron
transfer through space and whose rate constants have been
F a F f F
ð10Þ
ð11Þ
m
f,1
f,2
kꢀm
dF
f,1
¼
k F ꢀ k
F
ꢀ k1Ff,1 ¼ 0
m
m
ꢀm f,1
dt
23,36,37
known to exponentially decay.
Then, with a proportion
km
F
¼
Fm
ð12Þ
constant R:
f,1
kꢀm þ k1
ꢀ
β d
k ¼ R e
3
ð25Þ
1
3
Without any evidence, the electron transfer rate inside the
poly(tetraphenyl)silole siloxane dielectrics is assumed to be
higher than that from the polymer dielectrics to the aluminum
1
ΔVf
t2
¼ C 3
ð26Þ
FB;2
ꢀβ 3 d
electrode, k . k
:
R e
1
ꢀm
3
km
where d is surely the interdistance between phenyl rings on the
silole rings:
F
¼
Fm
ð13Þ
f,1
k1
ΔVf
t2
For the second element:
FB;2
ꢀ1 β d
¼
C R
e 3
ð27Þ
3
3
k1
k2
Ff;1 f F f F
ð14Þ
ð15Þ
f;2
f;3
Nonetheless, collection of the n volumes arranged perpendicular
to the substrate surface through the dielectric films are modeled
as shown in Figure 13b, where l is volume height, perpendicular
to the substrate surface, and s the distance between the volumes
in the same direction, in which the number of C atoms is
considered. Here, “a volume” denotes a three-dimensional,
hypothetical box containing a poly(tetraphenyl)silole siloxane
chain.
dFf;2
dt
¼
k1Ff;1 ꢀ k2Ff;2 ¼ 0
k
1
Ff;2 ¼ k Ff;1
ð16Þ
2
Assuming that the rate constants for the electron transfer over
all volume elements inside the polysilole siloxane dielectrics
t ¼ nl þ ðn ꢀ 1Þs = nl þ ns
ð28Þ
are the same (k = k = ... = k ):
1
2
nꢀ1
Ff;2 ¼ Ff;1
ð17Þ
t
s ¼ ꢀ l
n
ð29Þ
As a result of the above assumptions, the volume densities
In Figure 13b, even though each volume element is tentatively
depicted by a polymer chain, this is not yet confirmed. As shown
again in Figure 13b, assuming that the interface between the
volume elements consists of the interactions of the peripheral
phenyl rings, the geometrical relationship between d and s can be
derived:
of the negative charges for all elements are also the same (F
f,1
=
F
= F = ... = F
), except for the nth element, which
f,nꢀ1
f,2
f,3
was located next to the silicon substrate. For the nth element:
kn ꢀ 1
kꢀs
F
f F f Fs
ð18Þ
ð19Þ
f;n ꢀ 1
f;n
dFf;n
qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi
¼
k
F
ꢀ k F ¼ 0
n ꢀ 1 f;n ꢀ 1
ꢀs f;n
d ¼ sðs þ p þ p cos χÞ þ C0
dt
kn ꢀ 1
kꢀs
qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiꢀ
ꢀ
ꢁ
p
Ff;n
¼
F
ð20Þ
0
2
f;n ꢀ 1
d ¼ sðs þ p þ p cos χÞ þ C0 C ¼ qðq þ pÞsin χ þ a q þ sin χ
2
!
!
As mentioned earlier, there are certainly SiꢀOꢀSi linkages at the
2
2
p
1
a
2
þ
cos χ þ 2 p þ
ð30Þ
interface between the (tetraphenyl)silole siloxane polymer dielec-
2
2
trics and the silicon substrate, resulting in higher electron transfer
nꢀ1
rates than those between π organic groups, k . k . Accordingly:
As shown in Figure 13c, if half (1.70 Å) of the van der Waals
38
ꢀ
s
thickness of the waiting phenyl groups is considered, the
minimal value of s is estimated to be approximately 0.85 Å. From
the previous report on benzeneꢀbenzene interaction, the edge-
to-face type configuration allows for a shorter intermolecular
distance by approximately 0.40 Å, compared with that obtained
Ff;n f 0
ð21Þ
Overall:
km
k1
Ff ꢂ Ff;1 ¼ Ff;2 ¼ Ff;3 ¼ ::: ¼ Ff;n ꢀ 1
¼
Fm ð22Þ
3
8
using a van der Waals thickness (2.50 Å versus 2.90 Å). In
addition, it was also reported for benzene dimer that the
intermolecular interaction potential energies were calculated to
be the largest (most negative) when the intermolecular distances
were approximately 3.79 and 5.00 Å for the T-shaped and the
Equation 8 is modified for a forward sweep direction:
Ff
1
ΔVf = ꢀ
t2
ð23Þ
ð24Þ
FB;2
ε0 εOX 2
3
28
parallel structures, respectively. Therefore, the intermolecular
distance (d) between the phenyl groups of the silole ring and the
approaching phenyl ring shown in Figure 13c would be in the
1
km F
1
2
1
ΔVf
3
m
2
t2
= ꢀ
t ¼ C
3
FB;2
3
k ε ε
k1
1
0 3 OX
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Table 2. In the Low Curing Temperature Ranges (e 150 °C), the Interchain Distance (s) and the Interdistance (d) of Phenyl
a
Rings, Defined in Figure 13b, were Estimated Using s = (t)/(n)ꢀl and d = 0.835s þ 0.346 (nm), respectively
s (d) (nm)
n = 6
n = 7
PM3 (l = 1.24)
n = 8
temp. (°C)
t (nm)
(t)/(205) (nm)
DFT (l = 1.37)
0.342 (0.643)
PM3 (l = 1.33)
DFT (l = 1.39)
PM3 (l = 1.30)
8
0
351
290
264
1.712
1.415
1.288
0.382 (0.681)
0.085 (0.411)
0.322 (0.624)
0.025 (0.362)
-0.102 (0.273)
0.472 (0.767)
0.175 (0.490)
0.048 (0.381)
0.412 (0.709)
0.115 (0.437)
1
00
50
0.045 (0.378)
1
ꢀ0.082 (0.285)
ꢀ0.042 (0.312)
ꢀ0.012 (0.334)
a
The height (l ) of a polymer chain was measured from the optimized geometry obtained from the DFT calculations, as shown in Figure 13e.
range of 3.79ꢀ5.00 Å corresponding to the s values of 0.48ꢀ1.87 Å
calculated by eq 30 (Supporting Information Figure S5), respec-
tively. Then, the s is tentatively assumed to decrease further to
approximately ꢀ0.1 nm. The maximal value of s is set to 0.2 nm,
which is thought to be sufficient for investigating the intermo-
lecular interaction of the phenyl rings.
volume elements stacked into the vertical direction of the films (n)
is estimated to be approximately 205. This implies that the total
height ((t)/(n)) of a volume element and the interchain distance
to the verticaldirection of the film is approximately 1.29 nm for the
150 °C-cured samples, as summarized in Table 2. This is close to
the effective height of the polymer (l , approximately 1.39 nm)
estimated from the optimized geometry using the DFT calcula-
tions for the molecular structure of the poly(tetraphenyl)silole
siloxane polymer (n = 7), as shown in Figure 13e. Then, very
naturally, each of the volume elements was found to be occupied
by a polymer chain of n = 7. Using eqs 29 and 32, the s and d values
were also estimated for the 80, 100, and 150 °C cases, as
summarized in Table 2. For the case of 80 °C, depicted as the
kinetically preferred state, the s and d values were estimated to be
In the range of ꢀ0.100 nm < s < 0.200 nm (Supporting
Information Figure S5):
d = 0:803s þ 0:346 ðnmÞ
Substituting eq 29 into eq 31:
ð31Þ
ð32Þ
ꢀ
ꢁ
t
d = 0:803 ꢀ l þ 0:346
n
0
.32 and 0.62 nm, respectively, while for the case of 100 °C,
Substituting eq 32 into eq 27 and taking the logarithm:
corresponding to the thermodynamically preferred state, they de-
creased to 0.025 and 0.362 nm, as shown in Figure 13f. In the latter
case, the molecular orbitals of the two phenyl rings are thought to
overlap each other, since the s value (0.025 nm) is slightly lower
than (0.085 nm) the van der Waals contact (Figure 13c). In the
case of 150 °C, the s and d values werefurtherdecreased to ꢀ0.102
and 0.273 nm, respectively, preferring the thermodynamically
preferred state.
"
#
ΔVf
FB;2
t2
ꢀ1
ln
¼ lnðC R Þ þ β d
ð33Þ
3
3
"
#
ꢀ
ꢁ
ΔVf
t
FB;2
t2
ꢀ1
ln
= lnðC R Þ þ β 0:803 ꢀ l þ 0:346
3
3
n
ð34Þ
0:803 β
Nevertheless, the distance decay constant for electron tunnel-
ing β is given by the effective barrier height (ΔE ), through
eff
"
#
39
ΔVf
t2
which electron tunneling proceeds:
FB;2
ꢀ1
3
t
ln
= ½lnðC R Þ ꢀ 0:803β l þ 0:346ꢃ þ
pffiffiffiffiffiffiffiffiffiffi
3
3
n
β ¼ ð10:25 nmꢀ1 eVꢀ1=2Þ ΔE
ð36Þ
eff
ð35Þ
As shown in Figure 13d, ln[(ΔV )/(t )] is plotted according
Then, ΔEeff is estimated to be approximately 1.4 eV, somewhat
close to 1.646ꢀ1.963 eV, as the energy difference between the
energy levels of the LUMO orbitals and vacuum level proposed
in the step potential model of Figure 5d. This is consistent with
the argument that the cofacial πꢀπ interactions between the
peripheral phenyl rings are mainly responsible for the interchain
electron tunneling.
f
FB,2
2
to the film thickness (t), which is compared to eq 35. From this
ꢀ
1
graph, (0.803 β)/(n) was estimated to be 0.047ꢀ0.020 nm in
3
the low-temperature ranges. As mentioned in the CꢀV results, the
ΔV
for the 80 °C-cured sample was undervalued, confirming
FB,2
only the lower limit, because the accumulation region in the CꢀV
graphs swept over the entire sweep voltages. Then, it is thought to
The driving force for this investigation initially came from
whether the poly(tetraphenyl)silole siloxane polymer really
shows the negative charge trapping phenomena as the authors
planned and designed after being motivated by its low-lying
be more reasonable to take the tentative (0.803 β)/(n) value
0.047) from the ΔVFB,2 values for the 100 and 150 °C-cured
3
(
samples.
The theoretical calculation of decay constant β for electron
LUMO orbital. The positive flat band shift (ΔV ) in the CꢀV
FB
tunneling between the interchains has not yet been executed.
However, the authors tentatively muse that the tunneling process
occurs between the (tetraphenyl)silole rings. More specifically, the
process involved the phenyl rings accompanying cofacial πꢀπ
interactions. Fortunately, in recent years, the β value for electron
tunneling through phenyl rings was determined to be near
graphs evidently proved that the presented design concept is
working. It became apparent that investigation of the relationship
between the macroscopic and microscopic properties was very
important for understanding the electronic and optical properties
of this kind of molecular solids. As discussed in the above
subsections, the changes in intermolecular distances of the
phenyl groups of adjacent polymer chains generated the variation
in the electron transfer rate between them, resulting in the
ꢀ
1
36,37
1
2 nm from a luminescence decay kinetics experiment.
Using the β value, for (0.803 β)/(n) = 0.047, the number of
3
7
783
dx.doi.org/10.1021/ja1108112 |J. Am. Chem. Soc. 2011, 133, 7764–7785

Journal of the American Chemical Society
ARTICLE
alteration of electron trapping of the silole rings. The concept of
quantum well structures was very useful to interpret the electron
trapping phenomena, wherein the intrachain energy barriers play
as an origin of the charge trapping and the interchain energy
barriers act as critical factors to control the electron transfer.
The authors therefore proffer that further experimental and
theoretical studies are needed to understand more concretely
the electron behavior of this kind of dielectrics that consists of
the polymer chains of the quantum well structures.
’ ASSOCIATED CONTENT
Supporting Information. Figure S1 (GPC results), Fig-
ure S2 ( H- and C NMR results), Figure S3 (demonstration of
proportional relationship between refractive index (n) and
density (N) of films), Figure S4 (schematic demonstration for
the charge trapping in reverse sweep direction), Figure S5 (plot
of distance between phenyl rings (d) versus interchain distance
S
b
1
13
(s)), Remark S1, and complete refs 16b and 20. This material is
available free of charge via the Internet at http://pubs.acs.org.
’
SUMMARY
’
AUTHOR INFORMATION
In this article, the authors present the first results of compel-
Corresponding Author
hdjeong@chonnam.ac.kr; hsohn@chosun.ac.kr
ling research centered upon negative charge trapping and inves-
tigation of its physicochemical origin in synthesized poly(tetra-
phenyl)silole siloxane thin films. This lab has investigated the
charge trap properties of siloxane-based dielectrics for a solution-
processed floating gate layer or nanodot layer in the OTFT-based
nonvolatile memory devices, where the phase problem and the
lacking engineering capabilities were found. Then, a new material
architecture combining the insulating properties of polyorganosi-
loxanes and the charge trap properties of π-conjugated organic
moieties was adopted, where the two components were linked
through direct chemical bonds. This is thought to prevent the
phase separation problem and allow for ease of engineering of the
charge trap properties. Through comprehensive discussions and
analysis, the novel hybridized polymer, poly(tetraphenyl)silole
siloxane, has been designed, synthesized, and characterized for the
first time. For a low-lying LUMO of the silole (silacyclopen-
tadiene) ring used as a charge-trapping center, the low-lying
LUMO was thought to facilitate electron capture from the Fermi
energy level of the adjacent metal and semiconductor layers. Then,
in the case of using dielectrics containing the silole moiety,
substantial electron trapping enhancement was expected. The
negative charge trapping behavior was experimentally validated
’ ACKNOWLEDGMENT
This research was supported by the Basic Science Research
Program through the National Research Foundation of Korea
(
NRF) funded by the Ministry of Education, Science and
Technology (No. 2010-0008824). This work was also supported
by the PLSI supercomputing resources of KISTI (Korea Institute
of Science and Technology Information). The experiments at the
PLS were supported in part by MEST and POSTECH.
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