Disubstituted pendant-functionalized insulating-conductive block copolymer with enhanced dielectric and energy storage performance

Article abstract of DOI:10.1016/j.reactfunctpolym.2019.03.013

Disubstituted pendant-functionalized block copolymer consisting of insulating polynorbornene and conductive polyacetylene segments was synthesized by tandem metathesis polymerization, and displayed high dielectric constant of 30 and low dielectric loss of about 0.04, while behaved bad film-forming property, which could be improved by incorporating polycyclooctene into polynorbornene backbone. The generated flexible block copolymer with well-defined nanostructure exhibited a relatively high dielectric constant of 23, very low dielectric loss of below 0.0035, and high stored and released energy densities up to 4.52 and 4.02 J cm^?3 at the electric field of 165 MV m^?1, respectively, accompanied with the charge-discharge efficiency of 90%. The enhanced dielectric and energy storage capability were attributed to high permanent dipole moment and the interfacial polarization derived from the unique nanomorphology of block copolymer.

Full text of DOI:10.1016/j.reactfunctpolym.2019.03.013

ReactiveandFunctionalPolymers139(2019)50–59
Contents lists available at ScienceDirect
Reactive and Functional Polymers
journal homepage: www.elsevier.com/locate/react
Disubstituted pendant-functionalized insulating-conductive block
copolymer with enhanced dielectric and energy storage performance
Dandan Zhou, Hengchen Zhang, Cuihong Ma, Huijin Han⁎, Meiran Xie⁎
School of Chemistry and Molecular Engineering, East China Normal University, Shanghai, 200241, China
A R T I C L E I N F O
A B S T R A C T
Keywords:
Disubstituted pendant-functionalized block copolymer consisting of insulating polynorbornene and conductive
polyacetylene segments was synthesized by tandem metathesis polymerization, and displayed high dielectric
constant of 30 and low dielectric loss of about 0.04, while behaved bad film-forming property, which could be
improved by incorporating polycyclooctene into polynorbornene backbone. The generated flexible block co-
polymer with well-defined nanostructure exhibited a relatively high dielectric constant of 23, very low dielectric
loss of below 0.0035, and high stored and released energy densities up to 4.52 and 4.02 J cm−3 at the electric
field of 165 MV m−1, respectively, accompanied with the charge-discharge efficiency of 90%. The enhanced
dielectric and energy storage capability were attributed to high permanent dipole moment and the interfacial
polarization derived from the unique nanomorphology of block copolymer.
Tandem metathesis polymerization
Dielectric constant
Polarization
Energy density
Nanostructure
1. Introduction
electric field, and cause Ohmic heating of the capacitors. Therefore,
enhancing the k is more practical than the breakdown strength to im-
For the next-generation energy harvesting and storage systems,
there is a crucial need to explore dielectric materials conforming to the
ongoing trend toward miniaturization coupled with higher performance
[1,2]. Compared with the ceramic counterparts, polymeric dielectrics
possess some advantages, such as a higher breakdown strength, a lower
dielectric loss, greater reliability and mechanical flexibility, and are
therefore an ideal choice for high-performance dielectric materials.
With regard to linear dielectric polymers, the energy density (Ue) can be
calculated as Ue = 0.5kε0E2, where k is the relative dielectric constant,
ε0 the vacuum permittivity, and E the applied electric field [2]. Biaxi-
ally oriented polypropylene (BOPP), the commercially state-of-the-art
polymer dielectric film, is attractive due to its high breakdown strength
(730 MV m−1) and low dielectric loss (0.02) [3]. Like most polymers,
however, suboptimal Ue attributed to the low k (≈ 2.2) of BOPP re-
sulted in devices occupying too much volume, falling to meet the de-
mand of minimization and high-performance of electronic devices. On
the basis of intrinsic breakdown theory, the challenge going forward is
to move from intrinsic breakdown to engineering breakdown through
inclusion of effects caused by chemical impurities (impurity states in
the bandgap) and defects manufactured in polymer films [4]. Mean-
while, for many widely used linear dielectrics, including BOPP, it was
found that the conduction loss becomes more significant at higher ap-
plied fields. Generally, these losses increase exponentially with the
prove Ue. An effective way to obtain high k while not making any
compromise in increasing dielectric dissipation is to introduce high-k
surface-modified ceramic fillers (e.g., BaTiO3) to polymer matrix [5–7].
In addition, further advances are attained by fabricating multilayer thin
films [8], core-shell structure [9,10], and switching to nanowires [11]
to improve the interfacial properties between ceramic fillers and
polymer matrix. High loading of the ceramic fillers in the polymer
composite, usually over 50 vol%, can raise k by about ten times relative
to the polymer matrix, but inevitably decreases the adhesion of the
composite thus deteriorating the adaptability and flexibility between
the composite and the organic circuit boards [12]. Another promising
approach is to fabricate percolative dielectric composites by dispersing
conductive domains (e.g. metal nanoparticles, graphene oxide, and
carbon nanotubes) in an insulating polymer matrix while preserving
facile flexibility and processability of polymer in the composites at-
tributed to the significantly decreased volume of conducting filler re-
quired to achieve high-k according to percolation theory [13,14].
However, simultaneously, the excessive dielectric loss incurred by
current leakage is undesirable, which imposes much challenge in con-
trolling the percolative composites with the threshold composition in
practical applications. To provide a practical way for the development
of advanced dielectrics with high k value as well as suppressed di-
electric dissipation [15], recently many efforts have been done, and one
⁎ Corresponding authors.
E-mail addresses: hjhan@chem.ecnu.edu.cn (H. Han), mrxie@chem.ecnu.edu.cn (M. Xie).
https://doi.org/10.1016/j.reactfunctpolym.2019.03.013
Received 25 January 2019; Received in revised form 7 March 2019; Accepted 12 March 2019
Availableonline12March2019
1381-5148/©2019ElsevierB.V.Allrightsreserved.

D. Zhou, et al.
ReactiveandFunctionalPolymers139(2019)50–59
direction is to explore some types of “all-polymer” approaches. One is
forming all-polymer percolative dielectric composites by adding con-
ductive π-conjugated polymer as a filler into insulating polymer matrix
[16,17]. Typical conductive π-conjugated polymers included polyani-
line, polythiophene, and polyacetylene (PA). Because of charge accu-
mulation at the interface between the insulating and the conductive
domains, the k and polarization density enhanced. Unfortunately, the
dielectric loss was excessive and thus lowered breakdown field. Un-
desirable macrophase separation was another problem existing in this
method. Still, all-polymer strategy provided dielectrics with greater
reliability, mechanical flexibility, and possibly less cost compared with
inorganic-organic composites. Another is preparing the block copo-
lymer dielectrics strategy, which takes advantage of the direct covalent-
linkage between the conductive π-conjugated segment and insulated
polymer, storing energy combining electronic conduction with inter-
facial polarization. A terthiophene-containing block copolymer mole-
cular architecture was reported by Tang [18], featuring conductive and
insulating blocks. The conductive block forms conjugated, nanoscale
domains with high electronic conductivity, which is dispersed in an
insulating polymer matrix to prevent percolation and interdomain
conduction, leading to strong electronic and interfacial polarizations in
the electric field, ultimately shedding light on large dielectric re-
sponses. Inspired by this work, we have reported interfacial-dominated
block copolymer with insulated polynorbornene (PNBE) and conductive
PA backbone synthesized by ring-opening metathesis polymerization
(ROMP) and metathesis cyclopolymerization (MCP), respectively [19].
The excellent dielectric performance with high-k and reduced dielectric
loss at a relatively reasonable range is ascribed to percolation theory in
which migration of charge in the conductive domains happened under
an applied electric field, leading to charge accumulation and local di-
pole formation along the boundaries between the conductive and the
insulating domains. Orientational polarization in the PNBE block
bearing difluorophenyl pendant contributes a lot as well. By replacing
difluorophenyl with bulky 3,5-bis(trifluoromethyl)biphenyl to enlarge
the scope of side groups on the PNBE-b-PA backbone, a superhelical
nanotube self-assembled from copolymer was studied, and most im-
portantly, the k enhanced without increasing dielectric loss owing to
attaching stronger polarity groups 3,5-bis(trifluoromethyl)biphenyl to
polymer chains [20]. Such polar groups include cyano [21], tri-
fluoromethyl [22], sulfone-containing moiety [23], and so on. Thanks
to the introduction of dipolar thiourea group to main-chain of poly-
mers, aromatic polythiourea possesses an enhanced k [24]. Easy dipole
rotation of small-sized polar sulfone group in the side chain enables
glassy polymer to exhibit a high k and a low dissipation.
conductive-insulative interfacial area, leading to excellent dielectric
properties ascribed to percolation theory. Meanwhile, the stronger is
the polarity of molecules, the greater is the degree of polarization, and
the k value is likewise higher [22]. Therefore, the rational design of
disubstituted pendants consisting of alternated electron-withdrawing
bis(trifluoromethyl)biphenyl moiety and electron-donating pyrrolidine
and ether moieties on PNBE backbone was adopted to facilitate delo-
calization of the electrons, and increased dipole moment and dipole
polarizations, compared to monosubstituted PNBE, can be expected for
further improving the dielectric properties of block copolymer synthe-
sized by tandem metathesis polymerization (ROMP-MCP), which con-
tributed from the combination of dipolar, electronic, and interfacial
polarizations of polymer chain.
2. Experimental
2.1. Materials
Lithium aluminum hydride (LiAlH4), cyclooctene (COE), tricyclo-
hexylphosphine[1,3-bis(2,4,6-trimethylphenyl)-4,5-dihydroimidazol-2-
ylide-ne]benzylide-ne ruthenium dichloride (Ru-II), tetrakis(triphenyl-
phosphine)palladium(0) (Pd(P(C6H5)3)4), bis(trifluoromethyl)phenyl
boronic acid, dimethyl 5-aminoisophthalate, norbornene-5,6-endo-di-
carboxylic anhydride (NDA), 4-dimethylaminopyridine (DMAP), 4-bro-
mophenol, diisopropyl azodicarboxylate, triphenylphosphine, 1-(3-di-
methylaminopropyl)-3-ethylcarbodiimide hydrochloride (EDCI·HCl), and
bis(trifluoromethane)-sulfonimide lithium (LiTFSI) were commercially
available from Sigma-Aldrich or Energy Chemical Co. and used as re-
ceived without purification. [1,3-Bis(2,4,6-trimethylphenyl)-4,5-dihy-
droimidazol-2-ylidene][3-bromopyridine]2
benzylidene
ruthenium
dichloride (Ru-III) [29], 4-carboxyl-1,6-heptadiyne [30], 3′,5′-bis(tri-
fluoromethyl)biphenyl-4-ol [31] were prepared according to the known
synthetic procedures in the related literatures. Chloroform (CHCl3),
carbon tetrachloride (CCl4), ethyl acetate (EtOAc), hexane, ethanol,
methanol, and magnesium sulfate (MgSO4) were purchased from
Shanghai Chemical Reagents Co, and used as received without pur-
ification. Solvents were distilled over drying agents under nitrogen prior
to use: dichloromethane (CH2Cl2) from calcium hydride; ether, tetra-
hydrofuran (THF) and dioxane from sodium/benzophenone. All reac-
tions were carried out under dry nitrogen atmospheres using standard
Schlenk-line techniques.
2.2. Characterization
1H (500 MHz), 13C (125 MHz), and 19F (471 MHz) nuclear magnetic
resonance (NMR) spectra were recorded on a Bruker DRX500 spectro-
meter with tetramethylsilane as an internal standard in CDCl3 at room
temperature. Ultraviolet-visible (UV–vis) absorption spectra were
measured on a UV-1800 spectrometer with the concentration of
1.0 × 10−3 mg mL−1. Fourier transform infrared spectroscopy (FT-IR)
spectra were recorded on a Nicolet Nexus 670 in the region of
4000–400 cm−1 using KBr pellets. High resolution mass spectrometer
(HRMS) was measured by a Bruker QTOF micromass spectrometer.
Elemental analysis was conducted with an Elementar vario EL. Gel
permeation chromatography (GPC) was used to calculate the relative
molecular weight and polydispersity equipped with a Waters 1515
Isocratic HPLC pump, a Waters 2414 refractive index detector, and a set
of Waters Styragel columns (7.8 × 300 mm with particle size of 5 mm;
pore diameter of 103, 104, and 105 Å). The injection loop, the columns
and the RI detector were thermostated at 35 °C in the same oven. The
eluent was a solution of LiTFSI (0.05 mol L−1) and 1-(n-butyl) imida-
zole (0.05 mol L−1) in THF. The samples were dissolved in eluent so-
lution at 0.2 wt% before being filtered on a polytetrafluoroethylene
(PTFE) filter with 0.22 μm diameter pore size. Number average mole-
cular weight (Mn) and polydispersity index (PDI) were calculated ac-
cording to a calibration curve established with polystyrene standard.
PNBE is known as insulating low-dielectric materials [25–27], and
can readily be functionalized by attaching conjugated aromatic struc-
ture and dipolar groups, such as bis(trifluoromethyl)biphenyl to the
side chain to substantially improve the k value [22]. Meanwhile, the
presence of an endo-N-arylpyrrolidine moiety fused to the PNBE pre-
pared via ROMP appears to be essential to control the stereoregularity
(the stereochemistry of double bonds and the tacticity) and is thought
to have stereoregular microstructure. Thus, the polar groups could be
aligned in one orientation, and the k value substantially raised. Besides,
the superior thermal properties, that is, high glass-transition tempera-
ture (Tg) and degradation temperature (Td) enable PNBE to exhibit high
temperature capability necessary for these demanding high-tempera-
ture capacitors. PA is the simplest conductive π-conjugated polymer
possessing good electrical and thermal conductivity. The third genera-
tion Grubbs catalyst (Ru-III)-initiated MCP of 1,6-heptadiyne mono-
mers provides PA derivative with improved solubility, stability, and
excellent processability [28]. Based on the features of PNBE-b-PA
backbone [19,20], herein, we try to expand the π–π stacking interaction
by attaching strong polar conjugated aromatic 3,5-bis(trifluoromethyl)
biphenyl groups to both insulating PNBE and conductive PA blocks to
explore more possibilities to obtain well-defined nanostructure in
which the unique nanostructure can act as a nanocapacitor with a large
51

D. Zhou, et al.
ReactiveandFunctionalPolymers139(2019)50–59
Thermal gravimetric analysis (TGA) was performed using an
SDTA851e/SF/1100 TGA instrument under nitrogen flow at a heating
rate of 10 °C min−1 from 25 to 800 °C. Differential scanning calorimeter
(DSC) was performed on a Netzsch 204F1 in nitrogen atmosphere. An
indium standard was used for temperature and enthalpy calibrations.
All the samples were first heated from 30 to 200 °C and held at this
temperature for 3 min to eliminate the thermal history, and then, they
were cooled to room temperature and heated again from 30 to 200 °C at
a heating or cooling rate of 10 °C min−1. The hydrodynamic diameter
(Dh) was determined by means of dynamic light scattering (DLS) ana-
lysis using a Malvern Zetasizer Nano-ZS light scattering apparatus
(Malvern Instruments, U.K.) with a He—Ne laser (633 nm, 4 mW). The
Nano ZS instrument incorporates noninvasive backscattering (NIBS)
optics with a detection angle of 173o. The z-average diameter of the
sample was automatically provided by the instrument using cumulate
analysis. Transmission electron microscopy (TEM) images were re-
corded on the JEOL2100F microscopes operating at 120 kV. Samples for
TEM measurement were prepared by depositing a drop of CCl4 solution
with a determinated concentration on the copper grids coated with
carbon, followed by air-drying. Additionally, the samples were not
stained before measurement because the electron density difference
between the two blocks provided sufficient contrast for TEM imaging.
Atomic force microscopy (AFM) observation was performed on an SPM
AJ-III atomic force microscope at a measurement rate of 1.0005 Hz in
the tapping mode, and the AFM image was obtained at room tem-
perature in air. Samples were prepared by drop coating a CCl4 solution
at 0.01 mg mL−1 on a freshly cleaved mica surface, and were then air-
dried at room temperature. Each sample was examined at least twice
under the same conditions, and the images were found to be re-
producible. Dielectric analysis was conducted on a Novo control BDS40
dielectric spectrum analyzer at room temperature, and a contacting
electrode method was used. Measurements were carried out at low
applied voltage (typically 10 mV) and varying frequency (typically
100 Hz to 1 MHz) for 3–5 specimens of each sample to ensure re-
producibility. The edge side (d) of the guarded electrode is 0.3 mm. The
capacitance (Cp) and the dissipation factor (dielectric loss, loss tangent)
of the tested films were recorded and the data were interpreted using a
parallel RC circuit model expected to describe a “leaky” capacitor. The
dielectric constant (εr) can be calculated by Eq. (1):
sonicated for 5 min to give a solution concentration of 5 wt%. The so-
lution was then spin-coated on indium tin oxide (ITO) substrate at
spinning speed of 1000–1200 rpm for 30 s, and dried in air overnight,
followed by continuous heat treatment (40 °C) in vacuum for another
12 h to evaporate all the solvents. Film thickness was measured at
multiple positions with Elcometer (Veeco Dektak 6 M) and found to be
about 1.6 μm, and then, a thin layer (500 Å) of Pt particle was sputtered
on the exposure sides of the sample for dielectric impendence mea-
surement. For COE-contained block copolymers, the films were pre-
pared by casting solution at 2–9 wt% on a copper for dielectric im-
pendence and polarization measurements, affording films with uniform
thickness and good quality. For polarization measurements, the thinner
a film was, the higher the maximum electric field could be imposed
with a fixed voltage about 10,000 V, however, if the film was too thin, it
was more likely to crack. Considering the two aspects and film-forming
property of polymers, the thicknesses of films were controlled to be
about 15 and 60 μm for two polymers by adjusting the solution con-
centration measured at several locations by the Elcometer (SMD-565J-
L), the copper substrate served as the bottom electrode, and silver
electrode was sputtered on the film's top surface, while the film with a
thickness of 60 μm spontaneously peeled away from the copper sub-
strate to afford the free-standing film, and silver electrodes were sput-
tered on both sides of the film. According to the literature [18], there
were not any significant variations in dielectric properties for polymer
with different electrodes.
2.4. General procedure for polymerization
2.4.1. Syntheses of homopolymers
Typically, ROMP of bis-3,5-(3′,5′-bis(trifluoromethyl)biphenyloxy
methyl)-phenyl norbornene pyrrolidine (BTNP) using Ru-III as the
catalyst was conducted in a dry nitrogen-filled Schlenk tube. A solution
of Ru-III (1.8 mg, 2 μmol) in 0.2 mL of THF, which was degassed in
three freeze-vacuum-thaw cycles, was added to a degassed solution of
monomer BTNP (169.5 mg, 0.2 mmol) in 0.8 mL of THF to give a
monomer concentration of 0.2 mol L−1. After stirring for 0.5 h, the
polymerization was quenched by adding 0.2 mL of ethyl vinyl ether
with stirring for 0.5 h, and then precipitated into an excess of methanol.
Polymer was dissolved in CH2Cl2, and precipitated out once again from
methanol. The obtained poly[bis-3,5-(3′,5′-bis(trifluoromethyl)biphe-
nyloxy methyl)-phenyl norbornene pyrrolidine] (PBTNP100) was dried
in a vacuum oven at 40 °C to a constant weight. 1H NMR (500 MHz,
CDCl3, 25 °C, TMS): δ = 7.9–7.38 (m, 10H), 7.19–6.29 (m, 7H),
5.57–5.27 (d, J = 46.9 Hz, 2H), 5.08 (s, 4H), 3.73–2.31 (m, 8H), 1.62
(s, 1H), 1.40 (d, J = 54.3 Hz, 1H); GPC: Mn = 70.1 kDa, PDI = 1.12.
= Cpl/ 0A
r
(1)
where l is the thickness of film and ε0, the permittivity of the free space,
is 8.85 × 10−12 F m−1, A is the guarded electrode area, A = πd2. Po-
larization measurements at higher applied voltages employed a polar-
ization tester (Precision Multiferroic, Radiant, Inc.). The electric dis-
placement (D)-electric field (E) hysteresis loops were recorded for
applied voltage up to 10,000 V with a cycle frequency of 100 Hz. The
electric field can be calculated by equation: E = U/t, where U is the
voltage, and t is the thickness of the film.
2.4.2. Synthesis of block copolymers via tandem ROMP and MCP
The block copolymer was prepared by tandem ROMP-MCP with the
sequential addition of monomers, and a typical procedure was as follows.
In a nitrogen-filled Schlenk tube, a solution of Ru-III (1.8 mg, 2 μmol) in
0.2 mL of THF, which was degassed in three freeze-vacuum-thaw cycles,
was added to a degassed solution of the first monomer BTNP (169.5 mg,
0.2 mmol) in 0.8 mL of THF to give a monomer concentration of
0.2 mol L−1. After stirring for 0.5 h at 30 °C, a degassed solution of the
second monomer 4-(3′,5′-bis(trifluoromethyl)biphenyl formate)-1,6-hep-
tadiyne (FBHD) (21.2 mg, 0.05 mmol) in 0.1 mL of THF was added under
nitrogen. After stirring for further 0.5 h, the polymerization was quenched
by adding 0.2 mL of ethyl vinyl ether with stirring for 0.5 h, and the so-
lution was then poured into an excess of methanol. The polymer was re-
dissolved in CH2Cl2, and precipitated out once again from methanol. The
obtained block polymer, poly[bis-3,5-(3′,5′-bis(trifluoromethyl)bipheny-
loxy methyl)-phenyl norbornene pyrrolidine]-block-poly[4-(3′,5′-bis(tri-
fluoromethyl)biphenyl formate)-1,6-heptadiyne] (PBTNP100-b-PFBHD25),
was dried under vacuum for 12 h at 30 °C and stored in a refrigerator. 1H
NMR (500 MHz, CDCl3, 25 °C, TMS): δ = 7.96–7.3 (m, o-ArH-CF3 + p-
ArH-CF3 + m-ArH-O), 7.17–6.31 (m, o-ArH-O + p-ArH-CH2O + o-ArH-
2.3. Sample preparation for dielectric and polarization measurements
Initially, films of homopolymers and COE-free block copolymers for
dielectric impendence measurements were attempted to be prepared by
solution casting method. Typically, polymer was dissolved in CCl4 at a
certain concentration and stirred for 12 h at room temperature, fol-
lowed by sonication for 5 min. After that, the solution was cast on a
copper substrate, and the solvent was removed by evaporation in air
overnight, followed by heat treatment (40 °C) in vacuum for another
12 h to evaporate all the solvents. At last, the film was obtained by heat
pressing. As the solution concentration was at 1–2 wt%, various cracks
happened in the film. When raising the concentration to 3 wt%, the
formed free-standing film was too brittle to handle for subsequent
measurements. As a result, films of homopolymers and COE-free block
copolymers were prepared by spin-coating method. Typically, polymer
was dissolved in CCl4, stirred for 12 h at room temperature, and
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D. Zhou, et al.
ReactiveandFunctionalPolymers139(2019)50–59
CH2O + C=C-CH=CH), 5.57–5.27 (d, J = 48.2 Hz, CH]CHCH), 5.08 (d,
J = 30.0 Hz, CH2O), 3.61–2.60 (m, CHCHCH2 + CHCHCH2N +
CHCH2N + CH-COO + C=C-CH2), 1.63–1.54 and 1.48–1.34 ppm (m,
CHCH2CH); GPC: Mn = 82.4 kDa, PDI = 1.17.
3 min, it was found the integration area ratio of the signal at 6.20 ppm
to the signal at 5.57–5.27 ppm was 1:6.01, indicating 86% of monomer
has converted to polymer, and thus the monomer BTNP is of high ac-
tivity in ROMP and can be served as an ideal candidate as the first block
in preparation of block copolymers.
The COE-contained block copolymer was prepared by tandem ROMP-
MCP with the sequential addition of monomers, and a typical procedure
was as follows. In a nitrogen-filled Schlenk tube, a solution of Ru-III
(1.8 mg, 2 μmol) in 0.2 mL of THF, which was degassed in three freeze-
vacuum-thaw cycles, was added to a degassed solution of the mixed
monomers BTNP (169.5 mg, 0.2 mmol) and COE (22 mg, 0.2 mmol) in
0.8 mL of THF to give a monomer concentration of 0.2 mol L−1. After
stirring for 0.5 h at 30 °C, a degassed solution of the second monomer
FBHD (21.2 mg, 0.05 mmol) in 0.1 mL of THF was added under nitrogen.
After stirring for further 0.5 h, the polymerization was quenched by adding
0.2 mL of ethyl vinyl ether with stirring for 0.5 h, and the solution was
then poured into an excess of methanol. The polymer was redissolved in
CH2Cl2, and precipitated out once again from methanol. The obtained
block polymer, poly[(bis-3,5-(3′,5′-bis(trifluoromethyl)biphenyloxy me-
thyl)-phenyl norbornene pyrrolidine)-random-cyclooctene]-block-poly[4-
As the great impact of solvent on MCP of 1,6-heptadiyne derivatives
was concerned, MCP of FBHD was launched with a molar feed ratio of
FBHD to Ru-III ([M]/[Cat]) of 25:1 in a weakly coordinating solvent of
THF according to the previous reports [32–35]. At first trial, the
transformation of solution color from colorless to dark purple was ob-
served, and a violent reaction with a great deal of heat-releasing was
conducted when the catalyst Ru-III solution was injected into the
monomer solution at the monomer concentration of 0.2 mol L−1 at
30 °C, which was followed by several precipitations from excess me-
thanol to afford PFBHD as a dark purple solid in 95% yield. Thanks to
the good solubility of homopolymer PFBHD in common organic sol-
vents such as CH2Cl2 and THF (Table S1), the solution characterization
of PFBHD by GPC (Fig. 1 and Table 1) could be performed, and found
that PFBHD25 possessed Mn of 14.0 kDa and a narrow PDI of 1.11.
For block copolymer preparation, monomer BTNP was firstly poly-
merized in THF to yield living macro-initiator PBTNP with 100 repeat
units. After fully consumption of BTNP monitored by TLC (within
30 min), the second monomer FBHD (25 equiv) was added and initiated
by this living microspecies, and the solution color immediately changed
from dark yellow to purple, suggesting that MCP of FBHD happened and
the block copolymer PBTNP100-b-PFBHD25 with Mn of 82.4 kDa and PDI
of 1.17 was obtained by tandem ROMP and MCP. Similarly, the block
copolymers PBTNP100-b-PFBHD50 (Mn = 117.3 kDa, PDI = 1.18) and
PBTNP100-b-PFBHD100 (Mn = 128.9 kDa, PDI = 1.18) were further syn-
thesized under the high monomer feed ratios of 100:50 and 100:100, as
shown in Table 1. Comparing three block copolymers with a fixed
PBTNP length of 100 units and elongated PFBHD lengths from 25 to
100 units, Mn enhanced rapidly with the increase of feed ratio, con-
firming monomers BTNP and FBHD are of high activity in tandem
ROMP-MCP despite of the steric hindrance of bis(trifluoromethyl)bi-
phenyl pendants. Meanwhile, three copolymers all shared a narrow PDI
and a relatively good solubility in various organic solvents such as THF,
CH2Cl2. Although block copolymer PBTNP-b-PFBHD offers very pro-
mising dielectric properties, the polymer films were brittle and thus
difficult to handle. To enhance the processability and reliability, flexible
PCOE chain was incorporated into the main chain of PNBE to overcome
the film defects caused by the excessive rigidity of the copolymer.
Random copolymerization of BTNP (100 equiv) and COE (150 equiv)
was firstly carried out, after confirming the fully conversion of BTNP and
COE in a relatively short period of time (0.5 h) by TLC, the monomer
FBHD (25 equiv) was then added, accompanied by an instant color
change from dark-yellow to purple, suggesting an effective tandem
ROMP-MCP occurred, and afforded the block copolymer P(BTNP100-co-
COE150)-b-PFBHD25 in high yield of 93%, which has an increased Mn of
(3′,5′-bis(trifluoromethyl)biphenyl formate)-1,6-heptadiyne] (P(BTNP100
-
co-COE100)-b-PFBHD25), was dried under vacuum for 12 h at 30 °C and
stored in
a
refrigerator. 1H NMR (500 MHz, CDCl3, 25 °C, TMS):
δ = 7.96–7.3 (m, o-ArH-CF3 + p-ArH-CF3 + m-ArH-O), 7.17–6.31 (m, o-
ArH-O + p-ArH-CH2O + o-ArH-CH2O + C=C-CH=CH), 5.55–5.25 (s,
CH=CHCH + CH2CH=CHCH), 5.08 (d, J = 30.0 Hz, CH2O), 3.94–2.48
(m, CHCHCH2 + CHCHCH2N + CHCH2N + CH-COO + C=C-CH2), 1.02–
2.13 ppm (m, CHCH2CH + CH2-CH2); GPC: Mn = 90.1 kDa, PDI = 1.29.
3. Results and discussion
3.1. Synthesis and characterization of monomers
The functional monomer BTNP was synthesized starting from the
reaction of NDA with 5-amino-isophthalic acid dimethyl ester (1) in the
presence of catalyst DMAP in dioxane at refluxed temperature, firstly
achieved
N-(dimethyl
3,5-benzenedicarboxylate)-norbornene-di-
carboximide (NDI), which was then easily transformed into N-(3,5-di-
hydroxymethylphenyl)-norbornene-pyrrolidine (HNP) by reduction
reaction with LiAlH4 and followed by the etherification reaction with
3′,5′-bis(trifluoromethyl)biphenyl-4-ol (2), providing the functional
monomer BTNP as a white powder solid. Monomer FBHD was readily
afforded by the conventional esterification reaction of 2 and 4-car-
boxyl-1,6-heptadiyne using EDCI·HCl and DMAP as the catalyst system
in CH2Cl2 at room temperature, as fully illustrated in Scheme 1. The
structure characterizations of intermediates NDI and HNP as well as
monomers BTNP and FBHD were conducted and discussed in detail, as
depicted in Supporting Information.
3.2. Synthesis of polymers by ROMP and MCP
93.0 kDa and PDI of 1.27 compared with those of PBTNP100-b-PFBHD25
.
Homopolymer PBTNP was obtained by ROMP of BTNP, followed by
the synthesis of PFBHD via MCP, and block copolymer PBTNP-b-
PFBHD was finally acquired in the way of tandem ROMP and MCP
using Ru-III as catalyst. The synthetic routes are illustrated in Scheme
1. Initially, ROMP of BTNP with the molar feed ratio of BTNP to Ru-III
([M]/[Cat]) of 100 (Table 1) was conducted in THF. The solution color
immediately changed from colorless to dark-yellow when Ru-III was
added to the solution of BTNP. After 30 min, the reaction mixture was
precipitated into excess methanol for several times to give polymer
PBTNP100 with the Mn of 70.1 kDa and PDI of 1.12 in 94% yield.
Moreover, ROMP of BTNP initiated by Ru-III at [M]/[Cat] ratio of
100:1 was traced via the in-situ 1H NMR spectroscopy (Fig. S4). As
ROMP proceeded, the integration area of olefinic protons on the endo-
norbornene ring at about 6.20 ppm decreased, and correspondingly the
integration area of new signal of olefinic protons on the polymer
backbone at 5.57–5.27 ppm increased. When the reaction time was
Considering the low-dielectric and excellent film-forming property of
PCOE, random copolymerization of BTNP (100 equiv) and COE (100
equiv) was also adopted to optimize the dielectric property of copoly-
mers, and the copolymer P(BTNP100-co-COE100)-b-PFBHD25 was gained
in good yield with a lowered Mn (90.1 kDa) and rational PDI (1.29).
3.3. Structure characterization of polymers
Polymer structures were identified by NMR, IR, and UV–vis spec-
troscopy. In the 1H NMR spectrum of PBTNP (Fig. 2a), the signal of
olefinic protons on the endo-norbornene ring at about 6.21 ppm es-
sentially disappeared after ROMP, correspondingly the new signal of
olefinic protons on the backbone was observed at 5.57–5.27 ppm,
which evidenced the successful occurrence of ROMP. In the 1H NMR
spectrum of PFBHD (Fig. 2b), the signals at 2.82–3.69 ppm assigned to
methylene and methenyl protons (Hb,c) on the five-membered ring, and
53

D. Zhou, et al.
ReactiveandFunctionalPolymers139(2019)50–59
Scheme 1. General synthetic procedure of monomers and polymers.
the signals at 6.29–8.05 ppm belonged to the protons on biphenyl
(Hd,e,f,g) and conjugated double bond (Ha) were observed. Because the
acetylenic proton peak at 3.03 ppm for FBHD and the signals at
2.82–3.69 ppm for Hb,c on the five-membered ring of PFBHD came
closely, it was hard to decide whether or not the acetylenyl protons
completely disappeared. However, it was found the integration area
ratio of Sb,c at 2.82–3.69 ppm to Sa,d,e,f,g at 6.29–8.05 ppm was 5:8.93
(Fig. S5b), which was agreed well with the theoretical value of 5:9,
indicating that the acetylenyl group was completely transformed to
conjugated double bonds after a successful MCP. The 1H NMR analysis
conjugated double bonds on the trans-PA backbone at 6.78–6.58 ppm
suggested that the PFBHD backbone should have the trans-configura-
tion [36], which is likely of benefit to encourage positive accumulation
of dipole moments and elevate the dielectric constant. More im-
portantly, the block ratios could be analyzed by integrating the 1H NMR
signals. In the 1H NMR spectrum of PBTNP100-b-PFBHD25 (Fig. 2c), the
peak area ratio of the benzyl protons (Sh) at 5.08 ppm to the alkyl
protons (Sc,d,e,t,u) at 3.61–2.60 ppm was found to be 4.01:8.94 (Fig.
S6a), and the block ratio of PBTNP to PFBHD was calculated as about
1:0.19 (closed to 100:25), suggesting the tested block ratio matched
well with the monomer feed ratio. In the 1H NMR spectrum of P
(BTNP100-co-COE100)-b-PFBHD25 (Fig. 2d), the peak area ratio of the
protons (Sa,v) on polyolefin backbone at 5.55–5.25 ppm to the benzyl
of the representative copolymers PBTNP100-b-PFBHD25 and
P
(BTNP100-co-COE100)-b-PFBHD25 revealed the signals corresponding to
the constituent PBTNP and PFBHD blocks. The appearance of
Table 1
Characteristics for polymerization and the resultant polymers.a
[BTNP]:[COE]:[FBHD]:[Cat]b
Mn (kDa)
PDIc
Yield (%)
c
Run
Polymer
1
2
3
4
5
6
7
PBTNP100
100:0:0:1
70.1
14.0
82.4
117.3
128.9
90.1
93.0
1.12
1.11
1.17
1.18
1.18
1.29
1.27
94
95
97
95
94
95
93
PFBHD25
0:0:25:1
PBTNP100-b-PFBHD25
PBTNP100-b-PFBHD50
PBTNP100-b-PFBHD100
P(BTNP100-co-COE100)-b-PFBHD25
P(BTNP100-co-COE150)-b-PFBHD25
100:0:25:1
100:0:50:1
100:0:100:1
100:100:25:1
100:150:25:1
a
Polymerization conditions: T = 30 °C, Cat = Ru-III, solvent = THF, the initial concentration of BTNP = 0.2 mol L−1. Polymerization time (t): 0.5 h for homo-
polymers, (0.5 + 0.5 h) for all block copolymers.
b
c
The molar ratios in feed of monomers to initiator for homo- and copolymerization.
Determined by GPC using THF as eluent.
54

D. Zhou, et al.
ReactiveandFunctionalPolymers139(2019)50–59
PBTNP100
PFBHD25
absorbance of PBTNP at 240–350 nm corresponded to the phenyl
pendant, which was similar to that of monomer BTNP, as shown in Fig.
S7a. For PFBHD, the narrow absorbance at 220–300 nm was attributed
to the phenyl pendant, and the new absorption at 350–650 nm was due
to the conjugated PA backbone in contrast to that of monomer FBHD,
which convinced that MCP of FBHD was succeeded. In the UV–vis
spectra of copolymers, the distinct absorbance at 240–350 nm for
PBTNP block and the broad peak at 350–650 nm for PFBHD block were
observed in CHCl3 (Fig. S7b), which affirmed that PBTNP block was
connected to the PFBHD block.
PBTNP100
PBTNP100
PBTNP100
-
b
b
b
-
-
-
PFBHD25
PFBHD50
PFBHD100
-
-
P BTNP100-co-COE100
P BTNP100-co-COE150
-
b
-
PFBHD25
PFBHD25
-
b
-
3.4. Nanostructure and morphology of block copolymer assembly
Compared with microsized domains, the well-defined nanodomains
possess larger interacting areas with the insulating polymer matrix, and
thus have access to the potential to develop sufficiently higher polar-
ization levels, breakdown strengths, and mechanical property en-
hancements, offering more acceptable possibilities to engineer and
optimize the properties of all-polymeric molecular dielectrics. More
importantly, the small size of the nanoparticles makes it possible to
reduce the dimensions of polymer nanocomposite-based devices, con-
forming to the ongoing trend toward miniaturization electronic devices.
The self-assembly behavior of polymers was determined by DLS ana-
lysis, which showed the aggregates existed in solution with various
4
5
6
7
8
9
10 11 12 13 14 15 16 17 18 19
Retention time (min)
Fig. 1. GPC traces of homo- and copolymers.
protons (Sh) at 5.08 ppm was 0.9:1 (Fig. S6b). Considering the theore-
tical ratio of Ha on PNBE backbone to Hh on the benzyl group is 1:2, the
loading ratio of PBTNP to PCOE can be calculated as 0.5:0.4 (closed to
1:1). While the peak area ratio of Sh at 5.08 ppm to Sc,d,e,t,u at
3.61–2.60 ppm was 1:2.32 or 4:9.28 (Fig. S6b), and the block ratio of
PBTNP to PFBHD was calculated as about 4:1 (i.e. 100:25). All of the
points showed the loading ratios of PBTNP to PCOE and PFBHD were
1:1:0.25 (i.e. 100:100:25), which was almost consistent with the feed
ratios for the preparation of copolymers. In general, block polymers
were synthesized as expected structure via a tandem ROMP-MCP pro-
cess. In addition, the stretching vibration of unsaturated C ≡ C-H at
3273 cm−1 and C^C at 2116 cm−1 disappeared in FT-IR spectrum of
PBTNP100-b-PFBHD25 (Fig. S3c) and P(BTNP100-co-COE100)-b-
PFBHD25 (Fig. S3d), and both curves consisted of the CeN stretching
vibration at 1390 cm−1 in PBTNP and carbonyl stretching vibration
absorption at 1755 cm−1 in PFBHD, indicating the successful pre-
parations of copolymers PBTNP100-b-PFBHD25 and P(BTNP100-co-
COE100)-b-PFBHD25 via a tandem ROMP-MCP process.
number-average Dh. Actually, PBTNP100-b-PFBHD25 and P(BTNP100
-
co-COE100)-b-PFBHD25 displayed a single molecular chain distribution
in THF and CHCl3 without any assembly formation at the concentration
of 1 mg mL−1 (Figs. S10a,b and S11a,b). Interestingly, when changing
the solvent to CCl4, which is a good solvent for PBTNP block while a
poor solvent for PFBHD block (Table S1), the existence of self-assembly
nanostructures with different sizes could be verified by the DLS tech-
nique (Figs. S10c and S11c), and the results pointed out that PBTNP100
-
b-PFBHD25 exhibited homogeneous size of nanostructure with the
dominating Dh of 85 nm in CCl4 at the concentration of 1 mg mL−1. For
comparison, the DLS analysis of P(BTNP100-co-COE100)-b-PFBHD25 re-
vealed two distinct populations with the Dh of 90 nm for single as-
sembly and 300 nm for a combination of several assemblies at the same
concentration of 1 mg mL−1, indicating an enhanced self-assembly
ability for P(BTNP100-co-COE100)-b-PFBHD25 in a selective solvent.
Meanwhile, the π–π stacking interaction and van der Waals interactions
between the rigid trans-PA backbones and aromatic side chains as well
Additionally, UV–vis analysis may provide information about the
structure of polymers. After ROMP of BTNP, the only change was the
transformation of the double bond on endo-NBE to polyene backbone,
which had no contribution to UV–vis absorption. Therefore, the
Fig. 2. 1H NMR spectra of (a) PBTNP100, (b) PFBHD25, (c) PBTNP100-b-PFBHD25, and (d) P(BTNP100-co-COE100)-b-PFBHD25 in CDCl3.
55

D. Zhou, et al.
ReactiveandFunctionalPolymers139(2019)50–59
Fig. 3. TEM images of PBTNP100-b-PFBHD25 (a, b) and P(BTNP100-co-COE100)-b-PFBHD25 (c, d) at 0.3 mg mL−1 in CCl4.
as the steric bulk effect may serve as the driving force. Further insight
into the well-defined nanomorphology of the copolymer film formed by
the solvent evaporation in air was conducted through the TEM analysis,
where no staining for the copolymer film was integrant in consideration
of the electron density diversity between the PBTNP segments and
electron-rich PFBHD segments providing sufficient contrast. TEM
images of PBTNP100-b-PFBHD25 (Fig. 3a) and P(BTNP100-co-COE100)-b-
PFBHD25 (Fig. 3c) dictated the formation of regular nanosphere mor-
phology with average diameters of 82 nm and 89–120 nm at
0.3 mg mL−1 in CCl4, respectively. Magnified images revealed that
PBTNP100-b-PFBHD25 (Fig. 3b) and P(BTNP100-co-COE100)-b-PFBHD25
(Fig. 3d) shared a similar nanomorphology, the alternating gray and
black concentric circles. The gentle difference in Dhs evaluated by two
analyses may be explained as that the flexible PBTNP and PCOE blocks
with good solubility in selective solvent tended to pack loosely around
the rigid PFBHD block, so that the copolymers with more flexible
chains can engineer the nanostructures with much larger Dhs in solu-
tion. The vivid morphology of polymer chains was obtained by AFM
analysis (Fig. S12). Under diluted conditions (0.01 mg mL−1), the self-
assembly of copolymer was driven by the solubility discrepancy of
blocks, and P(BTNP100-co-COE100)-b-PFBHD25 formed into the sphe-
rical nanostructure with an average height of 3.8 nm and dispersed
diameter of about 37 nm in CCl4 at 0.01 mg mL−1. Notably, unlike the
conventional self-assemblies inevitably requiring various additional
post-synthetic treatments, the self-assembly process of block copoly-
mers in selective solvents could be carried out in much simpler manner,
which provided more acceptable possibilities to engineer and optimize
the properties of nanodielectric polymers.
materials. Initially, films of homopolymers and COE-free block copo-
lymers for dielectric impendence measurements were attempted to be
prepared by casting solution on a copper. However, the films with
cracks or brittle were unable to handle for subsequent measurements
(Fig. S13a-d). Instead, films were prepared by spin-coating the solution
on ITO substrate with a typical thickness of about 1.6 μm for dielectric
impendence measurements as a function of frequency at room tem-
perature, and the photographs of films were presented in Fig. S13e-h.
As frequency varied from 1000 Hz to 1 MHz, homopolymers PBTNP100
and PFBHD25 displayed high-k values of 21–19 and 15–12 (Fig. S14a),
respectively, and the dielectric losses (tanδ) were lower than 0.05 (Fig.
S14b), however, negative values were observed at the low frequency
range due to poor signal-to-noise ratio (SNR) resulting from uneven
surface of the sample or microcracks in the sample. Sequentially, block
copolymers PBTNP100-b-PFBHD25 and PBTNP100-b-PFBHD100 with
well-defined nanomorphology exhibited enhanced k values of 30–25
and 25–22, respectively, which were one order of magnitude higher
than that of common polymer dielectrics (k = 2–3) [40,41]. Mean-
while, the dielectric loss values of PBTNP100-b-PFBHD25
(tanδ = 0.05–0.01) and PBTNP100-b-PFBHD100 (tanδ = 0.066–0.001)
stayed at a reasonable low level in a wide frequency range from
1000 Hz to 1 MHz (Fig. S14b). The competitive dielectric performance
of PBTNP100
,
PBTNP100-b-PFBHD25
,
and PBTNP100-b-PFBHD100
mainly came from the contributions of electric and strong dipolar po-
larizations. The rational design of alternating electron-withdrawing bis
(trifluoromethyl)biphenyl moiety and electron-donating pyrrolidine
and ether moieties in PBTNP being bonded covalently was adopted to
facilitate delocalization of the electrons and increase the dipole mo-
ment. Meanwhile, the presence of N-arylpyrrolidine moiety endo-fused
to the PNBE prepared via ROMP seemed to endow the main chain with
stereoregular microstructure [19,20,22], affording the positive accu-
mulation of dipole moments in polar pendants under an electric field,
and the k value substantially raised accompanying with the dielectric
loss at a reasonable range. According to percolation theory [13,14],
when the conductive π-conjugated PFBHD segments were directly
covalent-bonded to the insulating PBTNP chains, in particular, bene-
fiting from a well-defined nanostructure, the interfacial polarization
3.5. Dielectric and electrical energy storage properties of polymers
As well known to us, the dielectric performance of polymers was in
close contact with their various polarizations and the chain micro-
structure [22,37–39]. A well-defined core-shell nanomorphology will
be remained in polymer film formed by the solvent evaporation in air,
which would be of benefit to elevate the dielectric properties. High-k
and low loss are important indicators of high-performance dielectric
56

D. Zhou, et al.
ReactiveandFunctionalPolymers139(2019)50–59
between the two blocks plays an important part in the improved di-
electric responses of the materials for the reason that the large inter-
facial area between the conductive domain and the insulating domain
promotes the accumulation of space charge and high interfacial polar-
ization, forming many microcapacitors under an electric field
[18–20,37–39], thus enhancing the dielectric performance. Compared
is an amorphous, glass-phase dipolar polymer with weak coupling
among the dipoles, consequently producing little polarization hysteresis
loss even at high electric field, which is different from the strongly
coupled dipolar polymers such as semicrystalline PVDF [24]. Appar-
ently, the covalent linkage between the conductive and insulating
blocks and the formed unique nanostructure may be beneficial to de-
crease the leakage current. Last but not least, the good solubility and
excellent film quality of polymer resulted in the improved dielectric
properties. Additionally, another advantage of P(BTNP100-co-COE100)-
b-PFBHD25 was the stability of dielectric constant and almost in-
dependence of the frequency ranging from 100 Hz and 1 MHz, which is
attributed to the nano-scale domains of the conductive PFBHD seg-
ments to permit a sufficiently fast polarizable response even under high
frequency electric fields.
to PBTNP100-b-PFBHD25
,
PBTNP100-b-PFBHD100 bearing longer
PFBHD segment possessed a slightly reduced k value and a varied di-
electric loss between 1000 Hz and 1 MHz, which was explained that
lower k may mainly be related to the decresed content of strong po-
larized insulating PBTNP block in the copolymer and the reduced
quality of the polymer film attributed from more PFBHD block with bad
film-forming property. Unfortunately, the uneven surface or micro-
cracks on thin film samples made it hard to conduct the dielectric im-
pedance test. The interference of the instrument expressed as the ap-
pearance of a “valley” in the curves of dielectric constant and frequency
dependence in dielectric loss at the low frequency range. For these
polymers, due to their bad film-forming property, no further measure-
ment for polarization could be performed. To overcome the film defects
caused by the excessive rigidity of the copolymer, the flexible PCOE
segments were attached into the main chain of PBTNP, endowing the
copolymer with better processability and reliability. Expectedly, when
the insulating low-dielectric COE units (100 and 150 equiv) were
bonded to PBTNP100-b-PFBHD25, the film-forming property of the
corresponding P(BTNP100-co-COE100)-b-PFBHD25 and P(BTNP100-co-
COE150)-b-PFBHD25 was improved (Fig. S13i,j), while their k values
decreased to 23 and 15 (Fig. 4a), respectively, which were inferior to
that of PBTNP100-b-PFBHD25, however, they were still close to the best
values reported for PNBE-b-PA backbone-based percolative materials
[19,20]. Considering that the high dielectric loss of dielectrics may
lower the charge-discharge efficiency, cause dielectric heating, and
limit the frequency response of the capacitors, P(BTNP100-co-COE100)-
b-PFBHD25 may be the best candidate with a low and varied dielectric
loss of 0.0035–0.0015 (Fig. 4b), which was at a low level in the di-
electric polymer family and much lower than those of similar polymers
(tanδ = 0.08–0.02) [19,20]. All the k and dielectric loss values of
polymers are listed in Table S2. The high-k of the copolymers mostly
came from a combination of the actions of dipole polarization in the
PBTNP block and interfacial polarization between the two blocks.
Particularly, the excellent dielectric performance with high-k and low
dielectric loss is ascribed to the percolation phenomenon in which the
migration of charge in the conductive domains accelerated under an
applied electric field, accompanied by the charge accumulation and
local dipole formation along the boundaries between the conductive
and the insulating domains, leading to the improved interfacial polar-
ization [19,20]; meanwhile, inhibited by the insulating PBTNP domain,
the conductive PFBHD domain was hard to gather together, restraining
the dark current and lowering the dielectric loss effectively. Another
reason for the tiny loss may be that P(BTNP100-co-COE100)-b-PFBHD25
The block copolymers with well-defined nanostructures not only
showed good dielectric properties, but also performed an improved
breakdown field and energy density, which were tested by electric
displacement (D)-electric field (E) hysteresis loop measurement for
polymer films. As shown in Fig. 5a,b, P(BTNP100-co-COE150)-b-
PFBHD25 and P(BTNP100-co-COE100)-b-PFBHD25 maintained linear
polarization behavior (narrow linear D-E loops) at 165 MV m−1. No-
tably, the linear polarization behavior and low hysteresis were an im-
portant indicator of the low level of energy dissipation, which was
agreed well with the low dielectric loss in the impendence test. For P
(BTNP100-co-COE150)-b-PFBHD25, when the maximum applied electric
field was further enhanced, it exhibited slightly tilted up D-E loops, and
was penetrated at an electric field up to 254 MV m−1, affording the
maximum polarization of 2.2 μC cm−2 at 254 MV m−1. Expectedly, the
observed maximum polarization for P(BTNP100-co-COE100)-b-PFBHD25
was enhanced to 5.0 μC cm−2 at 165 MV m−1 compared to that of P
(BTNP100-co-COE150)-b-PFBHD25 even if the breakdown field fell.
Considering the film thickness (t) of 60 μm and the maximum applied
voltage (Umax) about 10,000 V, the maximum applied electric field
(Emax) can be calculated about 166 MV m−1 by equation Emax = Umax/t,
and thus no further polarization measurement for P(BTNP100-co-
COE100)-b-PFBHD25 could be performed. The maximum polarization
for either P(BTNP100-co-COE150)-b-PFBHD25 or P(BTNP100-co-COE100)-
b-PFBHD25 was at a relatively high level in the polymer dielectrics at
the same applied electric field [18,24,37–39], and much higher than
the best value (1.5 μC cm−2 at 200 MV m−1) for the PNBE-b-PA
backbone-based nanodielectric polymers [19,20]. High electric dis-
placement for the copolymer films resulted from their high k, which
should be favorable for high energy density.
The energy density (Ue) of a capacitor was given by the equation:
Ue = ʃE(dD), where E and D were the applied electric field and electric
displacement, respectively. The stored energy density (Us) or released
energy density (Ur) of the polymers as a function of the maximum
electric field strength might be attained by the integration of the charge
or discharge curve of the D-E loop according to this formula [42].
Fig. 4. Dielectric constant (a) and dielectric loss (b) of polymers versus frequency with different COE contents.
57

D. Zhou, et al.
ReactiveandFunctionalPolymers139(2019)50–59
Fig. 5. Polarization (D-E) loops (a, b) and energy density (c, d) for (a, c) P(BTNP100-co-COE100)-b-PFBHD25 and (b, d) P(BTNP100-co-COE150)-b-PFBHD25
.
Charge-discharge efficiency (η) is also an index to measure the energy
storage performance of dielectric materials representing the efficiency
of energy conversion of capacitors, which was given by the equation:
η = (Ur/Us) × 100%. The energy that is unable to be released by elec-
tric energy would generally be converted into heat energy, which may
not only cause the waste of energy, but also be responsible for the re-
leased heat of capacitors because of the poor thermal conductivity of
polymers. Once the temperature exceeds the service temperature of the
capacitor, the capacitor will easily be destroyed. As the COE units in
copolymers reduced from 150 to 100, the Us and Ur for P(BTNP100-co-
COEx)-b-PFBHD25 increased from 3.40 and 2.50 J cm−3 at 254 MV
m−1 to 4.52 and 4.02 J cm−3 at 165 MV m−1 (Fig. 5c,d), respectively.
Obviously, considering the electric field remained at 165 MV m−1, the
polarization still increased as the low-dielectric PCOE content decrease.
Fortunately, this is accompanied by a decrease in energy dissipation
(the area enclosed by the charge–discharge cycle), which provided P
(BTNP100-co-COE150)-b-PFBHD25
and
P(BTNP100-co-COE100)-b-
Fig. 6. Energy density versus efficiency of P(BTNP100-co-COE100)-b-PFBHD25
at 165 MV m−1, PTNP75-b-PFHD25 at 270 MV m−1, and some other polymer
dielectrics at 400 MV m−1 in literatures.
PFBHD25 with η values of 74% at 254 MV m−1 and up to 90% at 165
MV m−1 (Fig. S15), respectively. Compared with P(BTNP100-co-
COE150)-b-PFBHD25 and some other all-polymeric dielectrics
[20,43–45], as illustrated in Fig. 6, P(BTNP100-co-COE100)-b-PFBHD25
showed a combination of high energy density and efficiency level under
the applied electric field, which seems to be a better candidate for di-
(BTNP100-co-COE100)-b-PFBHD25 with well-defined nanostructure ex-
hibited a high-k value of 23, low dielectric loss of below 0.0035, high
stored/released energy density of 4.52/4.02 J cm−3 at 165 MV m−1
,
electric materials. Furthermore, it should also be noted that P(BTNP100
-
and high η value of 90%, which were superior to those of the common
commercial polymer dielectric materials and the similar PNBE-b-PA
backbone-based percolative materials. Therefore, it was a promising
dielectric material which was attractive for practical applications in
thin film capacitor and other dielectric devices.
co-COE100)-b-PFBHD25, immune to moisture attributed to the hydro-
phobic aromatic rings and trifluoromethyl group as well as easy to
soluble in common organic solvents, is very attractive for practical di-
electric applications, especially in solution-processed organic electronic
devices.
4. Conclusions
Acknowledgements
In summary, by combining the strong dipolar and interfacial po-
larizations as well as stereoregular chain microstructure, the re-
presentative dipolar glass insulated-conductive block copolymer P
This research was financially supported by the National Natural
Science Foundation of China (No. 21574041, 21704025).
58

D. Zhou, et al.
ReactiveandFunctionalPolymers139(2019)50–59
Appendix A. Supplementary data
[24] S. Wu, W. Li, M. Lin, Q. Burlingame, Q. Chen, A. Payzant, K. Xiao, Q.M. Zhang,
Aromatic polythiourea dielectrics with ultrahigh breakdown field strength, low
dielectric loss, and high electric energy density, Adv. Mater. 25 (2013) 1734–1738.
[25] B.-G. Shin, T.-Y. Cho, D.Y. Yoon, Structure and properties of polynorbornene de-
rivatives: poly(norbornene dicarboxylic acid dialkyl ester)s and poly(norbornene
dimethyl dicarboxylate)s, Macromol. Res. 15 (2007) 185–190.
Supplementary data to this article can be found online at https://
doi.org/10.1016/j.reactfunctpolym.2019.03.013.
[26] B. Liu, Y. Li, A.S. Mathews, Y. Wang, W. Yan, S. Abraham, C.S. Ha, D.W. Park,
I.L. Kim, Synthesis of vinyl-type functionalized polynorbornenes with cyclic pen-
dant imide side groups by using palladium-based catalyst for low dielectric constant
materials, React. Funct. Polym. 68 (2008) 1619–1624.
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