Structure-conductivity relationship for peptoid-based PEO-mimetic polymer electrolytes

Article abstract of DOI:10.1021/ma300775b

Polymer electrolytes offer great potential for application in lithium batteries. In order to systematically optimize the performance of these materials, atomic level synthetic control over the polymer chemical structure is desired. In this study, we designed a series of chemically defined, monodisperse peptoid polymers to explore the impact of side-chain structure on the thermal and electrical properties. A series of comblike peptoid homopolymers with ethylene oxide (EO)_n side chains of varying length were synthesized by a rapid solid phase synthetic method. The electrical properties of these materials with dissolved lithium salt were characterized by ac impedance. The temperature dependence of the ionic conductivity of the polypeptoid electrolytes is consistent with the Vogel-Tamman-Fulcher equation. The optimum ionic conductivity of 2.6 × 10^-4 S/cm achieved for oligo-N-2-(2-(2-methoxyethoxy)ethoxy)ethylglycine-Li salt complex at 100 °C, is approximately 10-fold lower than the analogous PEO-salt complex. It is, however, nearly 2 orders of magnitude higher than previously reported comblike PEO-mimetic polypeptides. The ionic conductivities of these side chain analogs vary by 3 orders of magnitude, but this variation is entirely governed by the proximity of the system to the glass transition temperature. This investigation shows that polypeptoids provide a unique platform for examining the structure-property relationships of solid polymer electrolytes.

Full text of DOI:10.1021/ma300775b

Article
pubs.acs.org/Macromolecules
Structure−Conductivity Relationship for Peptoid-Based PEO−
Mimetic Polymer Electrolytes
Jing Sun,† Gregory M. Stone,‡,§ Nitash P. Balsara,‡,§,⊥ and Ronald N. Zuckermann*,†,§
†Molecular Foundry, Lawrence Berkeley National Laboratory, Berkeley, California 94720, United States
‡Department of Chemical and Biomolecular Engineering, University of California, Berkeley, Berkeley, California 94720, United States
§Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, California 94720, United States
⊥Environmental Energy Technologies Division, Lawrence Berkeley National Laboratory, Berkeley, California 94720, United States
S
* Supporting Information
ABSTRACT: Polymer electrolytes offer great potential for application in
lithium batteries. In order to systematically optimize the performance of these
materials, atomic level synthetic control over the polymer chemical structure is
desired. In this study, we designed a series of chemically defined, monodisperse
peptoid polymers to explore the impact of side-chain structure on the thermal
and electrical properties. A series of comblike peptoid homopolymers with
ethylene oxide (EO)n side chains of varying length were synthesized by a rapid
solid phase synthetic method. The electrical properties of these materials with
dissolved lithium salt were characterized by ac impedance. The temperature
dependence of the ionic conductivity of the polypeptoid electrolytes is
consistent with the Vogel−Tamman−Fulcher equation. The optimum ionic
conductivity of 2.6 × 10−4 S/cm achieved for oligo-N-2-(2-(2-methoxyethoxy)-
ethoxy)ethylglycine−Li salt complex at 100 °C, is approximately 10-fold lower
than the analogous PEO−salt complex. It is, however, nearly 2 orders of magnitude higher than previously reported comblike
PEO−mimetic polypeptides. The ionic conductivities of these side chain analogs vary by 3 orders of magnitude, but this variation
is entirely governed by the proximity of the system to the glass transition temperature. This investigation shows that polypeptoids
provide a unique platform for examining the structure−property relationships of solid polymer electrolytes.
A variety of functional groups other than the ether group can,
in principle, be used to solvate lithium cations. The electrolytes
in current lithium-ion batteries are mixtures of alkyl carbonates.
Nitrogen-containing groups such as nitriles have also been
shown to solvate lithium ions.10,11 Motivated by these
observations, we have designed a family of comblike PEO−
mimetic polymers based on an N-substituted glycine or peptoid
backbone. These polymers contain main-chain tertiary amide
functional groups in addition to the side-chain ether oxygens.
Polypeptoids are a class of sequence-specific polymer in which
chemically diverse side chains are attached to the amide
nitrogen.12 The iterative solid phase submonomer synthesis
method allows for the efficient synthesis of polypeptoids of
exact monomer sequence from an extremely diverse set of side-
chain functionalities derived from readily obtainable re-
agents.13−15 Unlike the typical polyamides, polypeptoids are
considerably more flexible due to the lack of backbone
hydrogen bonding. This is crucial for solid polymer electrolyte
design due to the requirement of rapid segmental motion for
ion transport. Hydrogen bonding in polypeptides, nylons, and
INTRODUCTION
■
All-solid rechargeable batteries containing a dry polymer
electrolyte are promising due to potential of creating systems
with high energy densities and extended cycle life. Poly-
(ethylene oxide) (PEO)-based materials have received the most
interest for polymer electrolyte applications due to their high
ionic conductivity (around 10−3 S/cm at 80 °C).1−4 In these
electrolytes, the lithium cations are complexed to ether oxygens
of the polymer. Thus, the cations are surrounded by a solvation
shell as is often the case for cations dispersed in conventional
low molecular weight electrolytes, such as water and alkyl
carbonates.5−7 There are two general mechanisms for ion
transport in electrolytes: the ion hopping mechanism wherein
ions hop from one liquidlike region to the next, and the
vehicular mechanism wherein the ion migrates with the
solvation shell more-or-less intact.8,9 In polymer electrolytes,
the vehicular motion is severely suppressed because motion of
the solvation shell requires motion of all of the entangled
chains. Ion transport in high molecular weight polymer
electrolytes is thus facilitated by rapid segmental motion in
the vicinity of the ions. Since polymer segmental motion is
affected by the glass transition temperature (Tg), rapid ion
transport is generally obtained in low-Tg polymers.
Received: April 16, 2012
Revised: May 21, 2012
Published: June 8, 2012
© 2012 American Chemical Society
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(d, 2H, S−CCH−CH), 7.36 (d, 2H, S−CCH−CH), 4.14 (t, 2H,
CH2−CH2−O−Ts), 3.69 (t, 2H, CH2−CH2−O−Ts), 3.60 (m, 6H,
CH3−O−CH2−CH2−O−(CH2)2), 3.52 (m, 2H, CH3−O−CH2−
CH2−O), 3.39 (s, 3H, O−CH3), 2,47 (s, 3H, C−CH3).
aramides typically dominates the physical properties of the
polymers, stabilizing the crystals and raising the melting point.
The N-substitution in peptoid polymers eliminates the
presence of hydrogen bond donors along the backbone,
allowing for a flexible backbone with reduced interchain
interactions and excellent thermal processability.16 Therefore,
the crystallization behavior of the polypeptoids is dominated by
their side chains. In contrast, PEO is a semicrystalline polymer,
and the conductivity of PEO/salt mixtures below 60 °C, the
melting point of PEO,17 decreases as an increasing fraction of
the chains are incorporated into the crystals.
Previous studies have addressed the efficacy of polymer
electrolytes that contain ether oxygens and other functional
groups. Poly[L-methoxytri(ethylene oxide)-L-glutamate] com-
plexes with a variety of alkali metal salts. However, due to the
rigid backbone of polypeptides, the ionic conductivity was
lower than 10−5 S/cm at 100 °C.18 Nagaoka et al. synthesized a
linear copolymer with dimethylsiloxane and ethylene oxide
segments.19 A variety of (EO)n-containing comb-branched
copolymers based on polysiloxane20 and polyphosphazene21
have been synthesized. None of these studies have demon-
strated conductivities comparable to that of PEO-based
electrolytes. The main difference between these studies and
our work is our ability to exert precise control over monomer
and polymer structure and thus to dissect the contribution of
various factors on the polymer properties. All of the polymer
chains examined in this study contain exactly 20 monomers,
and the number of pendant (EO)n units in the monomers is
precisely controlled (n = 1−3). This allows us to systematically
explore the impact of increasing the number of Li+ coordination
sites per chain.
PEO−mimetic materials have been widely investigated and
used for a variety of applications. In addition to a potential solid
polymer electrolyte candidate, peptoids of this class have also
been used as mimetics of poly(ethylene glycol) in biological
applications. For example, monodispersed homopolymers of N-
(2-methoxyethyl)glycine (Nmeg) have been used to enhance
the solubility of therapeutic peptides for drug delivery22 and to
inhibit the fouling of surfaces by the adsorption of proteins.23,24
Although quite effective, the only attempt to optimize the
chemical structure of these materials was by varying the degree
of polymerization. Variety and tunability of side chains enable
the systematic study of the relationship between polymer
structure and activity and to optimize PEO−mimetic materials
based on polypeptoids.
Synthesis of 2-(2-Methoxyethoxy)ethyl Azide. 50 g of 2-(2-
methoxyethoxy)ethyl tosylate (0.18 mol) was added to 300 mL of
DMF under N2. Subsequently, 40 g of sodium azide (0.62 mol) was
added into the mixture, and the reaction was stirred at 60 °C. After 36
h, the mixture was diluted with a large amount of water and extracted
with diethyl ether. The organic layer was washed with water, dried
1
over MgSO4, and evaporated under vacuum. Yield: 23.0 g, 90%. H
NMR (500 MHz, CDCl3): δ 3.69 (m, 4H, (CH2)2O), 3.57 (m, 2H,
CH2−CH2−N3), 3.41 (m, 5H, CH3−O and CH2−N3).
2-(2-(2-Methoxyethoxy)ethoxy)ethyl azide was synthesized with
1
the same protocol. Yield: 86%. H NMR (500 MHz, CDCl3): δ 3.68
(m, 8H, (CH2)2O), 3.56 (m, 2H, CH2−CH2−N3), 3.41 (m, 5H,
CH3−O and CH2−N3).
Synthesis of 2-(2-Methoxyethoxy)ethylamine. 20 g of 2-(2-
methoxyethoxy)ethyl azide (0.14 mol) was dissolved in 160 mL of
tetrahydrofuran. Triphenylphosphine (40 g, 0.15 mol) was added, and
the reaction was stirred overnight under an atmosphere of nitrogen.
The reaction was quenched with 220 mL of water, and the mixture was
allowed to stir overnight. The solids were removed by filtration, and
the supernatant was washed with toluene and dichloromethane. After
concentration in vacuo, 11.8 g of free amine was obtained. Yield: 72%.
1H NMR (500 MHz, CDCl3): δ 3.62−3.51 (m, 6H, (CH2)2O and
CH2−CH2−NH2), 3.37 (s, 3H, CH3−O), 2.89 (m, 2H, CH2−CH2−
NH2).
2-(2-(2-Methoxyethoxy)ethoxy)ethylamine was synthesized with
the same protocol. Yield: 84%. 1H NMR (500 MHz, CDCl3): δ 3.66−
3.50 (m, 10H, (CH2)2O and CH2−CH2−NH2), 3.39 (s, 3H, CH3−
O), 2.85 (m, 2H, CH2−CH2−NH2).
Synthesis of Peptoid Polymers. Polypeptoids were synthesized
on an automated robotic synthesizer or a commercial Aapptec Apex
396 robotic synthesizer on 100 mg of Rink amide polystyrene resin
(0.61 mmol/g, Novabiochem, San Diego, CA). All solvents, and
reagents described here were purchased from commercial sources and
used without further purification. The synthesis procedure was a
modified version of methods previously described.25 The Fmoc group
on the resin was deprotected with 20% (v/v) 4-methylpiperidine/
DMF before starting the submonomer cycle. An acylation step was
then performed on the amino resin by the addition of 1.0 mL of 1.2 M
bromoacetic acid in DMF and 0.18 mL of N,N′-diisopropylcarbodii-
mide and mixing for 20 min. Displacement of the bromide with
various submonomers occurred by adding a 1.0−2.0 M solution of the
primary amine in N-methyl-2-pyrrolidone, followed by agitation for 90
min. N-(2-Methoxyethyl)glycine polymers (pNme) were acetylated on
the resin after synthesis using a mixture (2.0 mL per 100 mg of resin)
of 0.4 M acetic anhydride and 0.4 M pyridine in DMF for 30 min. The
crude peptoid products were cleaved from the resin by the addition of
50% (v/v) trifluoracetic acid (TFA) in DCM for 5 min, followed by
evaporation under a stream of N2. Oligo-N-2-(2-methoxyethoxy)-
ethylglycine (pNde) and oligo-N-2-(2-(2-methoxyethoxy)ethoxy)-
ethylglycine (pNte) were cleaved with 95% (v/v) TFA in water for
5 min before acylation. Subsequently, the mixture was precipitated
with an excess of cold diethyl ether under vigorous stirring, followed
by centrifugation. The crude cleaved peptoids were then acetylated by
the addition of 40 mM of acetic anhydride and 40 mM pyridine in
DCM. After 20 min, the reaction mixture was precipitated with an
excess of diethyl ether.
EXPERIMENTAL PART
■
Synthesis of Submonomers. Synthesis of 2-(2-
Methoxyethoxy)ethyl Tosylate. In a round-bottom flask, 25 g of
diethylene glycol monomethyl ether (0.21 mol) was dissolved in 65
mL of tetrahydrofuran. Upon vigorous stirring at 0 °C, 65 mL of 6 M
sodium hydroxide solution was added. To this stirred solution, 50 g of
tosyl chloride (0.39 mol) dissolved in 70 mL of THF was added
dropwise under N2. After 1 h, the mixture was allowed to warm to
room temperature and then stirred for another hour. Finally, 500 mL
of diethyl ether was added, and the organic layer was washed with 1 M
aqueous sodium hydroxide and water, sequentially. After drying over
MgSO4, the organic layer was evaporated to yield a colorless liquid
(55.5 g). Yield: 96%. 1H NMR (500 MHz, CDCl3): δ 7.73 (d, 2H, S−
CCH−CH), 7.27 (d, 2H, S−CCH−CH), 4.11 (t, 2H, CH2−
CH2−O−Ts), 3.62 (t, 2H, CH2−CH2−O−Ts), 3.51 (m, 2H, CH3−
O−CH2−CH2−O), 3.40 (m, 2H, CH3−O−CH2−CH2−O), 3.24 (s,
3H, O−CH3), 2,37 (s, 3H, C−CH3).
The crude products were then dissolved in 1:1 mixture (v/v) of
acetonitrile/water and lyophilized. The peptoids were then purified by
reverse-phase HPLC on a C18 column (Vydac, 10 μm, 22 mm × 250
mm) using a linear gradient of 5−95% acetonitrile in water with 0.1%
TFA over 60 min at a flow rate of 10 mL/min.
Each final product was characterized by analytical reverse-phase
HPLC using a C18 column (10 μm, 50 mm × 2 mm) with 5−95%
gradient at 1 mL/min over 30 min at 60 °C. Peptoid purity was
determined using the analytical reverse-phase HPLC conditions
detailed above, and the molecular weight was determined by matrix-
2-(2-(2-Methoxyethoxy)ethoxy)ethyl tosylate was synthesized with
the same protocol. Yield: 96.7%. 1H NMR (500 MHz, CDCl3): δ 7.80
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assisted laser desorption/ionization mass spectrometry (Applied
Biosystems MALDI TOF/TOF Analyzer 4800) with a 1:1 (v/v)
mixture of peptoid (2 mg/mL in 1:1 acetonitrile: water) and 1,8,9-
dianthracenetriol dissolved in tetrahydrofuran at 10 mg/mL. The final
polypeptoids were then lyophilized prior to subsequent measurements.
Conductivity Measurements. LiTFSI (Li[N(SO2CF3)2, lithium
bis(trifluoromethanesulfonyl)imide)] salt from Sigma-Aldrich, was
heated at 120 °C under vacuum for 1−2 days to get rid of residual
moisture and then stored in an argon-filled glovebox. Polymer
electrolytes were prepared by mixing the polypeptoids with an
appropriate amount of the LiTFSI salt in the glovebox under an argon
atmosphere. The amount of dissolved salt in the polypeptoids is
quantified by the molar ratio of cations to peptoid ethylene oxide
moieties, r, which ranged from 0.02 to 0.12 in this study. After
lyophilization, the mixtures were dried under vacuum at 100 °C for 2−
3 days in the antechamber of the glovebox to remove any trace water
from the atmosphere. All procedures, performed after this step, were
conducted inside the argon-filled glovebox. The polypeptoid electro-
lytes were loaded into the center of Garolite G-10 spacers with 0.125
mm thickness and a central hole with a diameter of 3.86 mm. A
bubble-free disk ∼150 μm thick was obtained in a heated hand press
(100 °C) after 5 min. Stainless steel electrodes were placed on both
sides of the polypeptoid electrolyte disk to create the electrochemical
cell. After measuring sample thickness with a micrometer, the
electrochemical cell assembly was placed inside a Swagelok cell holder
and connected to a Biologic VMP3 for impedance analysis. The
Swagelok cell holder was loaded into a home-built heating chamber.
The imaginary and real components of the impedance response, Z′ and
Z″, were measured at 10 °C intervals during heating and cooling scans
between 30 and 100 °C. An ac signal with a peak voltage of 50 mV was
applied in the frequency range (ω) of 10 Hz−1 MHz. An equivalent
circuit comprised of a constant phase element in parallel with a resistor
was used to calculate the electrolyte resistance from the impedance
data. The thickness was measured again after the conductivity
measurement.
Differential Scanning Calorimetry (DSC). Differential scanning
calorimetry (DSC) experiments were performed to determine the
glass transition temperatures of the synthesized peptoids using a TA
Q200 differential scanning calorimeter. In all tests, a scan rate of 10 K/
min was used in the temperature range of −80 to 200 °C for three
heating and cooling cycles.
X-ray Diffraction (XRD). X-ray diffraction (XRD) was performed
at beamline 8.3.1 at the Advanced Light Source located at Lawrence
Berkeley National Laboratory. Samples were thermally annealed in a
vacuum oven at 100 °C for 2−3 days before the measurement. ∼1 mg
of the sample was loaded onto a nylon loop. The signal was radially
integrated to obtain a 1D plot of intensity versus scattering angle.
Thermogravimetric Analysis (TGA). Samples were characterized
using a TA Instruments TGA to investigate degradation temperatures
by mass loss. Approximately 5.0 mg of lyophilized peptoid powder was
placed on an aluminum sample pan. Samples were equilibrated at 30
°C for 20 min and then heated to 500 °C at 5 °C/min under a
nitrogen atmosphere.
Table 1. Model Acetylated Peptoid 20mers and Analytical
Data
a
a
n = 1, Nme = N-(2-methoxylethyl)glycine; n = 2, Nde = N-(2-(2-
methoxyethoxy)ethyl)glycine;
n = 3, Nte = N-(2-(2-(2-
methoxyethoxy)ethoxy)ethyl)glycine.
stable and their thermal behavior (e.g., glass transition
temperature, melting point) are highly tunable and dependent
on the monomer structure and sequence.16 The homopoly-
peptoids studied here show similar thermal stability, resisting
temperatures approaching 300 °C (Figure S2).
DSC traces of all three peptoids are shown in Figure S3. The
lack of melting peak and crystallization exotherm indicates that
these PEO−mimetic peptoids are noncrystalline. Furthermore,
the absence of ordered peaks in the XRD patterns is consistent
with the lack of crystalline domain (Figure S4). The glass
transition temperatures (Tg) as determined by DSC are shown
in Table 2. In general, the Tg values of the polypeptoids
Table 2. Tg of Polypeptoids with and without Salt
Tg (°C)
peptoid
p(Nme)20
p(Nde)20
p(Nte)20
r(Li:O) = 0
38.6
52.2
−6.8
9.9
−26.4
−2.4
r(Li:O) = 0.085
decrease with increasing side-chain EO unit length. Compared
to the Tg (38.6 °C) of p(Nme)20, which is above room
temperature, p(Nde)20 and p(Nte)20 have significantly lower
Tgs, around −6.8 and −26.4 °C, respectively. These values are
higher than the Tg of neat PEO (Mn = 2000), reported to lie
between −30 and −60 °C. It is noted that the Tg of PEO
depends on the fraction of crystalline domains.27,28 There is a
large Tg gap between polypeptoids with one and two EO
residues in the side chain [p(Nme)20 vs p(Nde)20], though the
weight fraction of EO moieties in the neat polypeptoids
increases gradually from 50.0% to 63.6% and 71.4% with the
increasing size of the side chain. It is clear that segmental
motion, which is directly related to Tg, depends not only on the
composition of the monomer but also on the arrangement of
the rapidly moving segments, (EO)n.
As shown in Table 2, the complexation of the polypeptoid
with lithium salts raises the glass transition temperature for all
of these polypeptoids. Furthermore, the Tg increases linearly
with increasing salt concentration, r (Figure 1). This is not
unexpected, as polymer chains coordinate with lithium ions to
form ionic cross-links,29 which limit the mobility of the polymer
chains, resulting in a higher Tg. More interestingly, the slope of
the Tg versus r line is larger for p(Nte)20 than for p(Nde)20 and
p(Nme)20. It is evident that the impact of lithium salt on the
peptoid Tg is greater for p(Nte)20 than for p(Nde)20 and
p(Nme)20. For r values greater than 0.15, the Tg of the complex
based on p(Nte)20 is higher than the one based on p(Nde)20.
Similar behavior has been observed in (EO)n-containing
comblike polystyrene.30 Molecular dynamics simulations31
RESULTS AND DISCUSSION
■
In this article, we designed and synthesized a series of well-
defined homopeptoids with a controlled numbers of ethylene
oxide units on each side chain. The structure of the synthesized
polypeptoids is shown in Table 1 along with their abbreviations.
Their molecular weight and purity characteristics are given in
Table 1. Their absolute monodispersity was confirmed by
analytical HPLC and MALDI mass spectrometry (see
Supporting Information, Figure S2). We designed 20mers for
the investigation with molecular weights ranging from 2362 to
4128 g/mol. Teran et al. have shown that the ionic conductivity
of PEO/LiTFSI mixtures is independent of PEO molecular
weight when it exceeds 2000 g/mol.26 It has been previously
reported that a wide variety of homopolypeptoids are thermally
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In this expression, σ is the ionic conductivity, A is a constant
proportional to the number of charge carriers, B is equivalent to
the activation energy for ion motion, R is the gas constant, T is
the temperature, and T0 is a reference temperature, which is
Tg(r,n) − 83 K; note that the Tg of the electrolytes depends on
both salt concentration and length of the side chains. Unlike
linear PEO, the PEO−mimetic polypeptoids do not crystallize
in our temperature window (regardless of salt concentration),
and thus the VFT equation fits the entire data set in Figure 2.
The polypeptoid with the largest molecular weight shows the
highest conductivity; it should be clear that the conductivity of
our electrolytes is dominated by segmental motion, not
vehicular motion. The conductivity of p(Nde)20 and p(Nte)20
is significantly higher than that of p(Nme)20. We thus focus on
p(Nde)20 and p(Nte)20 in our discussion of the effect of salt
concentration.
Figure 1. Tg of the polymer electrolytes versus salt concentration r.
Ionic conductivity of PEO−mimetic polypeptoids varies
significantly with the concentration of lithium salt dissolved in
the polymer, shown as in Figure 3. With increasing salt
and neutron diffraction data32 indicate that the Li+ ions are
coordinated with PEO through five ether oxygens of a polymer
chain.33 p(Nte)20 has longer side chains that appear to be more
effective in coordinating Li+. We thus expect the segmental
motion at a fixed temperature to be more rapid in p(Nte)20 in
the 0 < r < 0.15 regime. In the 0.15 < r < 0.24 regime we expect
more rapid segmental motion in p(Nde)20 at the same
temperature. The question of how changes in the strength of
complexation and segmental motion affect conductivity is an
interesting question that we will address shortly.
The ionic conductivities of several polypeptoid/LiTFSI
polymer electrolytes were measured as a function of temper-
ature and salt concentration. Measured values of ionic
conductivity of the three polymer electrolyte systems [p-
(Nme)20, p(Nde)20, and p(Nte)20] at r = 0.085, at temperatures
between 30 and 100 °C, are shown in Figure 2. In each case the
Figure 3. Ionic conductivity plots as a function of salt concentration r
at 50 and 90 °C.
concentration, the conductivity of the electrolytes initially
increases because more charge carriers are incorporated into
the system, despite a modest rise in Tg. However, as the salt
concentration increases further, the ionic conductivities in both
systems levels off and eventually decreases. This behavior can
be explained by considering the opposing influences of chain
mobility (as measured by the Tg) and charge carrier density
(proportional to salt concentration). For the p(Nte)20−Li salt
complexes, the Tg increases rapidly with increasing temperature
(Figure 1). The increasing concentration of charge carriers
results in an increase in conductivity at low r values, but this is
eventually offset by the decrease in segmental motion, resulting
at high r values. This eventually results in lower conductivities
at r > 0.14. Qualitatively similar behavior is seen in the
p(Nde)20−Li salt complexes. In general, the ionic conductivities
of the p(Nte)20−salt complexes are higher than those of the
p(Nde)20 complexes, until r (Li:O) reaches 0.14. Interestingly,
a crossover in conductivity between p(Nte)20 and p(Nde)20
takes place at an r of 0.14, which is approximately the same
point at which their Tg’s cross over (Figure 1). The p(Nde)20
electrolyte conductivity at the highest r values is comparable to
that at low r values at both 50 and 90 °C. In contrast, the
p(Nte)20 electrolyte conductivity at the highest r values is
significantly lower than that at low r values at 50 °C.
Figure 2. Ionic conductivity plots as a function of temperature at a salt
concentration of r = 0.085. The lines through the data points are VTF
fits.
ionic conductivity increases with increasing temperature, as is
typical of polymer electrolytes. The Vogel−Tamman−Fulcher
(VTF) equation,34 which is typically used to describe the steep
dependence of viscosity on temperature in the vicinity of the
glass transition, is often used to describe the temperature-
dependent conductivity data of polymer electrolytes.
σ(T) = A exp{−B/(R(T − T ))}
0
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The relative importance of complexation and segmental
motion of conductivity is addressed in Figure 4 where the
the ionic conductivity and the chemical structure of the
polymer electrolytes. In particular, we showed that the large
changes in conductivity seen in the systems at a fixed
temperature and salt concentration were correlated to changes
in Tg and not changes in other factors such as complexation and
length of the pendant (EO)n chains. Although, it is worth
mentioning that the Tg can be dependent on the strength of
complexation and length of pendant side chains. We hope that
such systematic studies will enable the design of polymer
electrolytes with higher conductivity than PEO and the
incorporation of other functional groups that can help rapid
transport of the lithium ions. Furthermore, the incorporation of
the homopolypeptoid blocks into a diblock copolymer with a
mechanically robust block is likely to yield polymer electrolyte
materials with high conductivity and high modulus.35 The
precise tunability of polypeptoids offer a platform for the
fundamental study of the structure−conductivity relationship of
block polymer electrolytes in a facile way.
Figure 4. Ionic conductivity plots as a function of 1000/(T − T0)
where T0 = Tg − 83 K. Note that Tg depends on both salt
concentration (r) and side chain length (n).
ASSOCIATED CONTENT
■
S
* Supporting Information
1H NMR spectra of the submonomers, analytical HPLC and
MALDI spectra, differential scanning calorimetry data, TGA,
XRD for polypeptoids. This material is available free of charge
via the Internet at http://pubs.acs.org.
conductivity of all 3 polymers at r = 0.085 is plotted versus 1/
(T − Tg(r,n) − 83 K) on a semilog scale. All of the data
collapse on to a single line. A fit through the data yields
AUTHOR INFORMATION
Corresponding Author
*E-mail: rnzuckermann@lbl.gov.
Notes
■
σ(T) = 0.485 exp{−1400/(T − T )}
0
The collapse in Figure 4 indicates that the conductivity of all of
the electrolytes is dominated by segmental motion, not by
complexation. The large changes in monomer structure as n is
varied from 1 to 3 have no impact on ion transport other than
the effect of this on segmental motion. This suggests that the
introduction of amide linkages in the polypeptoid backbone do
not significantly affect the coordination between the polymer
and the ions (Li+ and TFSI−). Thus, the poor conductivity of
p(Nme)20 seen in Figure 2 is attributed entirely due to the slow
segmental motion and not the inability of the single pendant
EO unit to coordinate with Li+. Similarly, the interesting
crossover in conductivities of p(Nde)20 and p(Nte)20 seen in
Figure 3 is also entirely due to changes in segmental motion.
Plotting all the other ionic conductivities with different r and n
as a function of 1000/(T − T0) showing the same slope (Figure
S5) confirms this conclusion. These data lie on different curves
because the number of charge carriers related to A is changing
with different r, but irrespective of n (Table S2).
The authors declare no competing financial interest.
ACKNOWLEDGMENTS
■
Funding for this work was provided by the Soft Matter Electron
Microscopy Program, supported by the Office of Science,
Office of Basic Energy Science, U.S. Department of Energy,
under Contract DE-AC02-05CH11231. The work was carried
out at the Molecular Foundry at Lawrence Berkeley National
Laboratory, supported by the Office of Science, Office of Basic
Energy Science, U.S. Department of Energy, under Contract
DE-AC02-05CH11231. We thank Dr. Daniel Hallinan, Scott
Mullin, and Alexander Teran for helpful advice and Babak Sanii
for help with the XRD experiments.
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