Synthesis and structure-property relationships of polypropylene-g- poly(ethylene-co-1-butene) graft copolymers with well-defined long chain branched molecular structures

Article abstract of DOI:10.1021/ma200604y

A series of polypropylene-g-poly(ethylene-co-1-butene) (PP-g-EBR) graft copolymers with well-defined long chain branched (LCB) molecular structures were synthesized via the combination of coordination polymerization and anionic polymerization. The structu

Full text of DOI:10.1021/ma200604y

ARTICLE
pubs.acs.org/Macromolecules
Synthesis and StructureꢀProperty Relationships of
Polypropylene-g-poly(ethylene-co-1-butene) Graft Copolymers with
Well-Defined Long Chain Branched Molecular Structures
†
,‡
†,‡
†,‡
†,‡
†,‡
†
,†
Lu Wang, Dong Wan, Zhenjiang Zhang, Feng Liu, Haiping Xing, Yanhui Wang, and Tao Tang*
†
State Key Laboratory of Polymer Physics and Chemistry, Changchun Institute of Applied Chemistry, Chinese Academy of Sciences,
Changchun 130022, China
‡
Graduate School of the Chinese Academy of Sciences, Beijing 100039, China
ABSTRACT: A series of polypropylene-g-poly(ethylene-co-
1
-butene) (PP-g-EBR) graft copolymers with well-defined long
chain branched (LCB) molecular structures were synthesized
via the combination of coordination polymerization and anionic
polymerization. The structureꢀproperty relationships of PP-
g-EBR were systematically investigated. The structural informa-
1
tion was clarified by characterization of H NMR and SEC
equipped with triple detectors. The DMA measurements de-
monstrated that the amorphous EBR branches were miscible
with the PP backbones. The melt behavior of the PP-g-EBRs
was studied with small-amplitude dynamic rheological measurements. Increasing branch length and branch density both led to
increased zero-shear viscosity, more pronounced shear-thinning behavior, elevated value of storage modulus at low shear
frequencies, more significant upshift deviation from linear behavior in the Han plot, and reduced loss angle. Comparatively, branch
length contributed more to the improvement of rheological properties than branch density did. Interestingly, DSC results showed
that the crystallization temperature of PP-g-EBR increased with the enhancement of branch length and branch density; however,
when the LCB level exceeded a certain degree, further increasing the LCB level would reduce the crystallization temperature.
1
. INTRODUCTION
reactive extruded PP are usually generated via radical induced
random chain scission followed by recombination. Thus, these
techniques create branched PP with broadened molecular weight
distributions and very complex topological structures, both of
which lead to the difficulty in the structure characterization of the
products.
Isotactic polypropylene (iPP) is one of the leading and fast
growing thermoplastic polymers in the world due to its out-
standing physical and chemical properties such as high melting
point, low density, high tensile modulus, excellent chemical
resistance, and low cost. However, commercial iPP, produced
with ZieglerꢀNatta or metallocene catalysts, exhibits relatively
low melt strength and no strain hardening behavior in the melt
state because it consists of highly linear chains and has a relatively
narrow molecular weight distribution, which limits its applica-
tions in blow molding, thermoforming, extrusion coating, and
foaming. The melt strength of PP can be improved by many
methods such as increasing the molecular weight, broadening
the molecular weight distribution, or introducing long chain
branches (LCBs), among which the most efficient way is to
introduce LCBs onto PP’s backbone. Apart from mediating
rheological property, the presence of LCBs with different chemical
structures from PP backbone may provide an opportunity to
modify other properties of PP such as mechanical properties.
Since the presence of LCBs is generally known to enhance
There are also some reports on the direct synthesis of LCBPP,
such as the copolymerization of in situ generated PP macro-
monomer with propylene by using a single metallocene catalyst
dimethylsilylbis(2-methyl-4-phenylindenyl)zirconium
1
3
14
dichloride or a binary single-site catalyst system. The addition
1
5
of previously prepared macromonomers such as atactic PP, iPP
and polyethylene (PE), poly(ethylene-co-propylene), or allyl-
16
17
18
terminated syndiotactic PP to the propylene polymerization
system catalyzed by a metallocene catalyst is another general
strategy to synthesize LCBPP. This method can be used to
synthesize LCBPP with relatively well-defined branches. How-
ever, the incorporation of macromonomers is difficult due to
their less mobility to diffuse toward the active centers of catalyst,
especially when the molecular weight of macromonomer is well
above its critical molecular weight for entanglements. As a
1ꢀ3
melt strength of a polymer,
developed to synthesize long chain branched polypropylene
several approaches have been
(
LCBPP). Some of them are based on chemical modification
Received: March 16, 2011
4
ꢀ9
of PP such as reactive extrusion
irradiation.
and electron beam
Revised:
April 27, 2011
1
0ꢀ12
The branched structures in the irradiated or
Published: May 09, 2011
r 2011 American Chemical Society
4167
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|Macromolecules 2011, 44, 4167–4179

Macromolecules
ARTICLE
consequence, the branch density of the resultant LCBPP is
limited. It is also difficult to separate the unreacted macro-
monomer from the resulting LCBPP. On the other hand, as
we know, the incorporation of comonomer will reduce molecular
weight of PP backbone. Therefore, it is difficult to obtain LCBPP
samples possessing a similar PP backbone but differing in branch
densities for the macromonomer copolymerization method.
Another important in situ in-reactor method for preparing
LCBPP is the copolymerization of propylene with a non-
formed benzylic anion effectively initiated anionic polymeriza-
tion of 1,3-butadiene at room temperature. At last, the poly-
butadiene (PB) side chains were subjected to a noncatalytic
hydrogenation reaction, yielding comblike LCBPP samples con-
taining iPP backbone and poly(ethylene-co-1-butene) (EBR)
LCBs. A series of PP-g-EBR graft copolymers with controlled
branch density and branch length were synthesized. Structureꢀ
property relationships of PP-g-EBRs were systematically
investigated. One of the most important aims in this work is to
give an insight into the influences of branch length and branch
density of long chain branched polymers on their rheological
properties.
19ꢀ23
conjugated diene comonomer.
The branch density of thus
13
obtained LCBPP could be calculated by the C NMR spectrum.
Nevertheless, the branch length could not be determined, and a
high concentration of diene copolymerized with propylene
would yield polymer gels. Lu and Chung prepared LCBPP with
a well-defined comblike structure through a graft-onto reaction
between a maleic anhydride-grafted PP (PP-g-MA) and several
2. EXPERIMENTAL SECTION
24
2.1. Materials. Toluene (AR grade, from Beijing Chemical Works)
was refluxed over metallic sodium with benzophenone and distilled
under argon atmosphere prior to use. Cyclohexane and hexane were
refluxed over metallic sodium/potassium alloy with benzophenone and
distilled under argon atmosphere before use. p-Toluenesulfohydrazide
amine-terminated PP (PP-t-NH2). Recently, Langston et al.
reported a new method to synthesize LCBPP via the metallocene-
mediated polymerization of propylene with T-reagent p-(3-
2
5
butenyl)styrene.
Although extensive studies have been reported in the literature
about the synthesis and rheological characterization of LCBPP,
the systematic study of the relationships between molecular
structure and the rheological behavior of LCBPP has been
seriously limited in the past compared to LCBPE. This is ascribed
to the difficulty of preparing LCBPP samples with well-defined
long chain branched molecular structures. Chung and co-workers
reported a novel chemical route for preparing PP graft copolymers
(
TSH, from Aldrich) was recrystallized from methanol. 4-Methylbenzyl
chloride (from Alfa Aesar), tripropylamine (TPA, from Sinopharm
0
0
Chemical Reagent Co., Ltd.), and N,N,N ,N -tetramethylethylenediamine
TMEDA, from J&K Chemical Ltd.) were vacuum-distilled over calcium
(
hydride prior to use. Polymerization grade propylene (from YanShan
Petrochemical Corp.) was used as received. Polymerization grade 1,
3
-butadiene was purchased from Jinzhou Petrochemical Corp., purified
by passing through four columns packed with 4 Å molecular sieves and
solid KOH before use. Allylmagnesium chloride (2.0 mol/L solution in
THF from Aldrich), diethyl ether (anhydrous, from Beijing Chemical
Works), methylaluminoxane (MAO: 10 wt % in toluene, Ethyl Corp.),
and rac-Me Si(2-MeBenz[e]Ind) ZrCl (MBI) (from Mitsubishi
Chemical Co.) were used as received. sec-Butyllithium (sec-BuLi: 1.3
mol/L solution in cyclohexane/hexane (92/8)) was purchased from
Acros and used as received. All oxygen- and moisture-sensitive manipu-
lations were carried out under dry and oxygen-free argon atmosphere
using standard Schlenk techniques.
26
with a controlled molecular structure. The poly(propylene-co-
p-methylstyrene) backbone was synthesized using hetero-
geneous ZieglerꢀNatta catalysts and subsequently lithiated in
the p-methyl groups of p-methylstyrene (p-MS) units with sec-BuLi
in order to form anionic active sites for initiating anionic
polymerization of several monomers. However, the multisite
heterogeneous ZieglerꢀNatta catalysts produce copolymers
with wide molecular weight distribution and poor control of
the copolymer composition. Chung and co-workers also succeeded
in preparing poly(propylene-co-p-methylstyrene) copolymer
2
2
2
2.2. Synthesis of p-(3-Butenyl)toluene (p-BT). To a dry
00 mL three-necked round-bottom flask equipped with an isobaric
27
5
with metallocene catalysts, but they found that catalyst activity
decreased sharply with increasing the content of p-MS in the feed
and the molecular weight of the resultant copolymer was not
high. The decrease of catalyst activity was speculated to result
from a steric jamming during the consecutive insertion of 2,
dropping funnel, a condenser pipe, and a magnetic stir bar was
transferred 100 mL (0.2 mol) of allylmagnesium chloride solution.
Then, 20 mL (0.15 mol) of 4-methylbenzyl chloride diluted with 50 mL
of diethyl ether was added dropwise to the flask through isobaric
dropping funnel at ice bath temperature. After complete addition of
27,28
1
-inserted p-MS and 1,2-inserted propylene.
merization of propylene with p-MS catalyzed by rac-SiMe2
2-Me-4-Ph-Ind) ZrCl /MAO predominantly gave p-MS-termi-
The copoly-
4
-methylbenzyl chloride, the mixture was warmed to room temperature
and stirred for 20 h. Then, 200 mL of distilled water was introduced to
the mixture slowly. The aqueous layer was separated through a
separatory funnel and extracted three times with diethyl ether.
The solvent was removed under reduced pressure. The obtained
crude product was dried with calcium hydride for 24 h and then distilled
(
2
2
2
8
nated PP (PP-t-p-MS).
In order to prepare PP copolymer with high catalyst activity,
high molecular weight, and narrow molecular weight distribu-
tion, we designed a new comonomer, p-(3-butenyl)toluene
under vacuum to yield the comonomer (p-BT) as a colorless liquid
(
p-BT), to be copolymerized with propylene using a metallocene
1
(20.9 g, yield: 95%). H NMR (CDCl , 400 MHz, ppm): δ 7.11 (m, 4H,
3
catalyst. Compared to p-MS, a longer distance between the double
bond and the bulky phenyl group in p-BT reduces the steric
hindrance and consequently facilitates the incorporation of p-BT
during copolymerization with propylene. In this work, we
describe a new method that can be used to prepare LCBPP with
well-defined molecular structure, i.e., known backbone molecular
weight, branch length, branch density, and narrow molecular
weight distribution of the LCBs. The poly(propylene-co-p-BT)
j-H), δ 5.87 (m, 1H, CꢀCHdC), δ 5.04 (m, 2H, CꢀCdCH ),
2
δ 2.70 (t, 2H, j-CH ꢀCdC), δ 2.34 (s, 3H,
j-CH ).
.3. Copolymerization of Propylene with p-BT. All the
2
), δ 2.39 (m, 2H, CH
2
3
2
polymerization runs were carried out in a 1 L stainless steel autoclave
equipped with a mechanical stirrer. The autoclave was dried in vacuo at
100 °C for 2 h and then cooled to 40 °C. The autoclave was then charged
with toluene, p-BT, MAO, and MBI catalyst solution. The total volume
of the feeding liquid was 300 mL. The polymerization was initiated by
introducing propylene gas. The autoclave was charged with propylene to
a total pressure of 0.65 MPa. The reaction vessel was stirred and
(
PPꢀBT) backbone was first synthesized using the catalyst system
of rac-Me Si(2-MeBenz[e]Ind) ZrCl (MBI)/MAO. Then the
2
2
2
methyl groups in the p-BT units were lithiated by sec-BuLi. The
4
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ARTICLE
Table 1. Comparison of the Experimental Results in the rac-Me Si(2-MeBenz[e]Ind) ZrCl (MBI)/MAO
2
2
2
a
Catalyzed Copolymerization of Propylene with p-(3-Butenyl)toluene
c
catalyst in feed
activity
p-BT in copolymer
b
6
ꢀ1 ꢀ1
d
d
e
e
sample
PP
[p-BT] (mol/L)
(μmol)
(10 g(mol Zr)
h
)
(mol %)
T
m
(°C) ΔH
m
(J/g)
w w n
M (kg/mol) PDI (M /M )
0
0.75
1.80
1.80
1.80
89.1
35.3
21.8
7.1
0
147.2
143.2
142.2
140.5
86.5
82.6
72.2
65.5
194
196
208
221
2.40
2.32
1.84
1.83
PPꢀBT1
PPꢀBT2
PPꢀBT3
0.023
0.046
0.092
0.20
0.39
0.73
a
Experimental conditions: solvent, toluene; total volume, 300 mL; propylene pressure, 0.65 MPa; polymerization temperature, 40 °C; Al/Zr (mol/mol) =
0 000; polymerization time, 90 min. p-BT denotes p-(3-butenyl)toluene. Calculated by H NMR spectra. Determined by DSC measurements.
b
c
1
d
1
e
Measured by SEC with light scattering detector.
maintained at 40 °C and 0.65 MPa for 90 min. The polymerization was
terminated by venting propylene gas and adding acidified ethanol
solution. The product was filtered, washed with a large amount of
ethanol, and then dried under vacuum at 60 °C for 24 h. The conditions
and results of the copolymerization of propylene with p-BT are
summarized in Table 1.
c
crystallization temperatures (T ) of polymers were determined by
differential scanning calorimetry (DSC) using a Perkin-Elmer DSC-7
instrument operating at a heating rate of 10 °C/min from 30 to 200 °C
under a nitrogen atmosphere. Absolute molecular weight, molecular
weight distribution, and intrinsic viscosity were obtained by size exclu-
sion chromatography (SEC) with triple detectors. The chromatographic
system consisted of a Polymer Laboratories PL 220 high-temperature
chromatograph equipped with a two-angle laser light scattering detector
(TALLS), a viscosity detector, and a differential refractive index
detector. Polymer solutions were prepared with amounts of about 20
mg of polymer in 10 mL of 1,2,4-trichlorobenzene containing a small
amount of antioxidant 2,6-di-tert-butyl-4-methylphenol (BHT), and were
eluted at 150 °C with a flow rate of 1 mL/min. Rheological measure-
ments were preformed on a rotation rheometer (Physica MCR-300) at
2
.4. Synthesis of PP-g-PB Graft Copolymers. The graft-from
2
6
polymerization was carried out according to the published literature.
In a typical experiment (sample 6 in Table 3), under an argon atmo-
sphere, 2.0 g of PPꢀBT2 containing 0.39 mol % of p-BT was suspended
in 60 mL of anhydrous cyclohexane in a dried Schlenk-type filtration
reactor with a magnetic stir bar. 2.9 mL of 1.3 mol/L sec-BuLi solution
and 1.2 mL of TMEDA were added to the reactor and then heated to
70 °C. After 4 h, the reaction was cooled to room temperature, and the
1
0
80 °C. The parallel plate with a diameter of 25 mm and a gap height of
resultant polymer was isolated by filtration and washed repeatedly with
hexane under an argon atmosphere until complete decoloration of the
filtrate.
.8 mm was used. The test samples were first treated with 0.2 wt %
Irganox B215 antioxidant and formed into disks with a diameter of
5 mm and a thickness of 1 mm by compression-molding at 180 °C
2
A prescribed amount of 1,3-butadiene was dissolved in 90 mL of
anhydrous cyclohexane and then transferred to the Schlenk-type filtra-
tion reactor containing lithiated PP-BT2 copolymer under an argon
atmosphere. The lithiated copolymer initiated the anionic polymeriza-
tion of 1,3-butadiene at 30 °C for a specific time (10ꢀ120 min). The
graft-from reaction was terminated by adding ethanol. The precipitated
polymers were filtered and then subjected to THF extraction to remove
soluble PB homopolymer. Then the product was dried in vacuo at 40 °C.
All the graft-from polymerizations were based on the same PPꢀBT2
backbone. The molar ratio of TMEDA and sec-BuLi used in all the
lithiation reactions of PP-BT2 copolymer was 2/1.
and 10 MPa. Then, the samples were quenched at room temperature.
The range of the frequency sweeps was from 0.01 to 100 Hz (0.0628
to 628 rad/s), and a strain of 1% was used, which was in the
linear viscoelastic regime for all samples. The rheometer oven was
purged with dry nitrogen to avoid degradation of samples during
measurements.
3. RESULTS AND DISCUSSION
3.1. Copolymerization of Propylene with p-BT. The copoly-
2
.5. Synthesis of PP-g-EBR Graft Copolymers. The PP-g-EBR
merization reaction isdescribed in Scheme 1, and the experimental
results are summarized in Table 1. A series of copolymerizations
have been performed under variable p-BT comonomer feed
ratios, in the presence of metallocene catalyst MBI activated by
MAO, to prepare the copolymers with variable compositions.
With the increase of the comonomer feed, the incorporation of
p-BT in the copolymer increased while the catalyst activity
decreased. The reduction of catalyst activity should result from
the steric hindrance due to the large size of the pendant group of
p-BT linked to the central metal of the catalyst, which is similar to
graft copolymers were prepared by hydrogenation of the obtained
PP-g-PB copolymers. In a typical reaction, to a 250 mL round-bottom
flask equipped with a magnetic stir bar was added 1.5 g of PP-g-PB
copolymer (sample 6 in Table 3 containing 20.1 mol % of 1,
3-butadiene), 3.62 g of p-toluenesulfohydrazide, 3.7 mL (2.78 g) of
tripropylamine, and 100 mL of toluene under an argon atmosphere. The
mixture was heated to reflux and stirred for 6 h. The mixture was
then cooled and poured into a large amount of ethanol. The precipi-
tated polymer was stirred in ethanol overnight, then filtered, washed
with ethanol for several times, and dried at 60 °C for 24 h under
vacuum.
2
2,29
the case of the copolymerization of propylene with R-olefins.
1
13
Comparing with the experimental results for the copolymerization
2
.6. Characterization. All high-temperature H NMR and
C
of propylene with p-MS using a similar catalyst (rac-SiMe (2-
2
NMR spectra were recorded on a Bruker AV400 spectrometer at 125 °C
using o-dichlorobenzene-d (o-C D Cl ) as solvent. The measurements
Me-4-Ph-Ind) ZrCl /MAO) in the literature (when 0.11 mol/L
2
2
4
6
4
2
p-MS was added at 30 °C and under 0.51 MPa, the catalyst
on the linear dynamic mechanical properties were conducted on a
e
activity was less than 1/10 000 of that in propylene homo-
dynamic mechanical analyzer (DMA/SDTA861 ). The test samples
27
polymerization), one can observe that the negative influence
of p-BT on the catalyst activity is much less than that of p-MS
(for example, the catalyst activity of copolymerization of pro-
pylene with 0.092 mol/L p-BT was about 1/13 of that in
propylene homopolymerization). The distance between the
were first treated with 0.2 wt % Irganox B215 antioxidant and were
formed into rectangular bars with a length of 9 mm, a width of 4 mm, and
a thickness of 1 mm by compression-molding at 180 °C and 10 MPa.
DMA was carried out between ꢀ100 and 130 °C at a constant frequency
of 10 Hz and a heating rate of 2 °C/min. Melting temperatures (T ) and
m
4
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ARTICLE
Scheme 1. Synthesis Route for PP-g-EBR Copolymers
double bond and the bulky phenyl group in p-BT is longer than
that in p-MS, which reduces the steric hindrance, facilitates
the incorporation of p-BT during copolymerization, and con-
sequently favors the catalyst activity. The incorporation of p-BT
comonomer almost did not affect molecular weights, but it
decreased melting temperatures and melting enthalpies (see
Table 1).
½ðp-BTÞðp-BTÞꢁ þ 1=2½ðp-BTÞPꢁ
np-BT ¼
ð3Þ
can be
1
=2½ðp-BTÞPꢁ
In addition, the monomer reactivity ratios r and r
calculated from the C NMR spectra using the subsequent
P
p-BT
13
3
1
equations
1
A typical H NMR spectrum of PPꢀBT copolymer (sample
2½PPꢁ
rP ¼
ð4Þ
ð5Þ
PPꢀBT2 in Table 1) is shown in Figure 1a. In addition to the
three major peaks centered at 0.99, 1.39, and 1.72 ppm, which
correspond to the CH , CH , and CH protons of the PP
½
ðp-BTÞPꢁX
3
2
2
½ðp-BTÞðp-BTÞꢁX
backbone, respectively, two minor peaks at 2.28 and 2.67 ppm,
rp-BT ¼
corresponding to j-CH and j-CH of the p-BT units, respectively,
½ðp-BTÞPꢁ
3
2
can also be detected. The content of p-BT in the copolymer is
calculated by the ratio of the integrated peak area at 2.67 ppm to
those at 0.8ꢀ1.8 ppm and the number of protons each peak
represents. The equation is as follows:
where X is the concentration ratio of propylene and p-BT in the
feed. The dyad distributions, number-average sequence lengths,
and values of reactivity ratios for PPꢀBT2 are summarized in
Table 2. The number-average sequence lengths of P and p-BT in
PPꢀBT2 are 249 and 1, respectively.
1
=2A2:67
Cp-BT ¼
ð1Þ
1
=6A0:8ꢀ1:8
Figure 3 displays the melting parameters listed in Table 1 as a
function of the p-BT content of the copolymers. The T and
m
From the calculation based on eq 1, the contents of p-BT in
PPꢀBT1, PPꢀBT2, and PPꢀBT3 are 0.20, 0.39, and 0.73 mol %,
ΔH of PPꢀBT copolymer decrease linearly as the content of
m
p-BT in the copolymer increases. These behaviors result from the
rejection of the side groups of the comonomers from the crystal
cores, which reduces the length and the concentration of crystal-
13
respectively. Additionally, Figure 2 shows the C NMR spectrum
(
with an inset of the chemical shift assignments) of PPꢀBT2; no
signal of continuous p-BT sequence was observed with about
600 scans. The CH carbon resonances at 46.6 and 43.9 ppm are
lizable unit sequences and consequently leads to a reduction of
6
32ꢀ34
2
the size and amount of crystals.
The subsequent equation,
assigned to the dyads PP and (p-BT)P, respectively, where P
denotes propylene unit and p-BT denotes p-(3-butenyl)toluene
unit. Number-average sequence lengths can be determined from
35
developed by Monrabal et al. according to the classic Flory
equation, can be applied to describe the dependence of melting
temperature of random copolymer on its low comonomer
content:
30
the dyad distributions using the following relationships:
0
RðT Þ
m
2
½
PPꢁ þ 1=2½ðp-BTÞPꢁ
0
nP ¼
ð2Þ
T = T ꢀ
N2
ð6Þ
m
m
1
=2½ðp-BTÞPꢁ
ΔH
u
4
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Table 2. Dyad Distributions, Number-Average Sequence
Lengths, and Values of Reactivity Ratios for PPꢀBT2
Copolymer
(p-BT)(p-BT)
(p-BT)P
PP
n
p
n
p-BT
r
P
r
p-BT
h
h
0
0.008
0.992
249
1
3.08
ꢀ
1
Figure 1. H NMR spectra of (a) PPꢀBT2 and (b) a representative
PP-g-PB copolymer in o-C D Cl at 125 °C.
6
4
2
Figure 3. Dependences of T
their p-BT contents.
m
and ΔH
m
of the PPꢀBT copolymers on
the PPꢀBT2 backbone is not soluble in cyclohexane, most of the
p-CH groups had to be located in the amorphous phase and
3
26
were readily accessible for the lithiation reaction. Therefore, in
most parts, the lithiation reaction and the subsequent graft-from
polymerization of 1,3-butadiene were not hindered by the
heterogeneous reaction conditions due to the swelling of the
amorphous phase. The branch density of PP-g-PB copolymer can
be controlled by regulating the efficiency of lithiation reaction of
PPꢀBT2 backbones. Because the graft-from polymerization of
1
,3-butadiene was a living anionic polymerization, it is reasonable
to assume that each benzylic anion produced one PB branch and
26
each branch had a similar length. In order to obtain PP-g-PB
copolymers with variable branch lengths, the graft-from poly-
merizations of 1,3-butadiene were carried out at different
concentrations and with different polymerization times due
to the living character of anionic polymerization. Finally,
a hydrogenation reaction was done to produce PP-g-EBR
copolymers.
1
3
Figure 2. Aliphatic region of C NMR spectrum of PPꢀBT2 in
o-C Cl at 125 °C.
6
D
4
2
0
where T is the melting temperature of homopolymer, T is
the equilibrium melting temperature of the copolymer, R is the
m
m
1
Figure 1b shows a representative H NMR spectrum of
gas constant, N is the molar fraction of comonomer incorpo-
the PP-g-PB copolymer (sample 6 in Table 3). In addition to
three major chemical shifts corresponding to the PP backbone,
there are several new chemical shifts corresponding to the PB
side chains. The coexistence of 1,2- and 1,4-microstructures are
demonstrated in chemical shifts at 5.01 ppm corresponding to
2
rated, and ΔH is the fusion heat per polymer repeating unit.
u
Therefore, a linear dependence of the observed Tm of the
obtained PPꢀBT copolymer on its p-BT content, consistent
with eq 6, implies the random distribution of p-BT in the
copolymer.
two external olefinic protons (dCH ), at 5.65 ppm correspond-
2
3
.2. Synthesis of PP-g-EBR Graft Copolymers. Scheme 1
ing to internal protons (ꢀCHd) in the 1,2-structure, and at
5.47 ppm corresponding to internal protons (ꢀCHd) in the
1,4-structure. The signals between 1.95 and 2.45 ppm corres-
pond to the allylic (ꢀCHꢀCdC) protons in the PB side chains
generated by the 1,2-insertion of 1,3-butadiene. The chemical
also illustrates the reaction process for the synthesis of PP-g-EBR
copolymers. The graft copolymers were obtained through three
further steps. In the first step, the above-obtained PPꢀBT2
copolymer was lithiated by sec-BuLi/TMEDA to obtain stable
benzylic anions. Then the obtained lithiated polymer was washed
with hexane for several times. Because of the insolubility of the
PPꢀBT2 copolymer in hexane at room temperature, the excess
sec-BuLi could be removed completely from the lithiated polymer.
In the second step, the benzylic anions effectively initiated
anionic polymerization of 1,3-butadiene in cyclohexane with a
living manner at 30 °C. All the PP-g-PB graft copolymers were
synthesized based on the same PPꢀBT2 backbone. Although
shift of aliphatic protons (ꢀCH ꢀCꢀCdC), generated by the
2
1,2-insertion of 1,3-butadiene and located at about 1.35 ppm, is
immerged in that of PP backbone, but its integrated peak area
is equal to that of 5.01 ppm. The content of PB side chains in
PP-g-PB graft copolymer and the molar ratio of 1,2 and 1,4
microstructures of PB side chains can be determined by
the relative integrated peak areas of 0.8 to 1.8, 5.01, and
5.47 ppm and the number of protons each peak represents.
4
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Table 3. Polymerization Conditions and Main Results of the H NMR Characterization of PP-g-PB Copolymers
ARTICLE
1
a
f
reaction conditions
structure of PB (mol %)
ꢀ
þ
b
c
d
e
g
h
i
sample PP Li (g)
BD (g) yield (g) BD in branched polymer (mol %)
1,2
1,4
GE (%) branch density branch length
1
2
3
4
5
6
7
8
9
1
1
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
10.4
8.3
7.0
9.8
6.9
8.6
6.8
9.6
6.3
9.6
6.8
2.05
2.02
2.15
2.30
2.10
2.62
2.75
3.20
2.64
3.55
2.40
2.5
0.9
77.7
77.9
79.7
77.3
76.6
78.8
77.4
79.2
77.8
77.0
76.9
22.3
22.1
20.3
22.7
23.4
21.2
22.6
20.8
22.2
23.0
23.1
25.9
6.0
5.07
1.17
4.49
7.46
1.78
1450
2180
6.1
22.9
38.1
9.1
4030
11.3
3.6
4760
5880
20.1
23.7
32.5
20.6
38.0
13.8
60.8
38.4
59.2
20.6
31.5
6.1
11.9
7.51
11.6
5930
11600
11700
18000
27800
37800
4.04
6.17
1.19
0
1
a
b
c
Solvent: cyclohexane, 90 mL; reaction temperature: 30 °C. Starting PPꢀBT2 copolymer (containing 0.39 mol % of p-BT). BD denotes 1,
d
e
1
f
3
-butadiene. Weight of the resulting PP-g-PB copolymer. Mol % of 1,3-butadiene in PP-g-PB, calculated by H NMR spectra. Mol % of 1,2 and 1,4
1 g 1 h
structure in PB branch, calculated by H NMR spectra. Grafting efficiency, calculated by H NMR spectra. Average number of branches per 10 000
1
i
1
carbons in the PP backbones, calculated by H NMR spectra. Average molecular weight of the EBR branches, calculated by H NMR spectra.
The equations are as follows:
1
=2A5:01 þ 1=2A
5:47
BD mol % ¼ 1=2A5:01
þ 1=2A5:47 þ 1=6ðA
ꢀ A5:01Þ
0:8ꢀ1:8
ð7Þ
1
=2A5:01
1
; 2-structure of PB mol % ¼ 1
ð8Þ
=2A5:01 þ 1=2A5:47
BD in eq 7 denotes 1,3-butadiene. The percentage of
,2-addition of 1,3-butadiene in all the graft-from polymer-
1
izations was about 80 (see Table 3). The reason for this
phenomenon must have been the presence of the polar
chelating agent TMEDA at the Li active site. From
þ
26
1
the comparison of H NMR spectra of the PP-g-PB copoly-
mer and the backbone PPꢀBT2 copolymer (see Figure 1),
1
Figure 4. H NMR spectrum of a representative PP-g-EBR copolymer
the increase of the relative intensity of the peak at 2.67
in o-C D Cl at 125 °C.
6
4
2
ppm corresponding to j-CH can be observed due to the
2
partial conversion of j-CH to j-CH in p-BT units during the
3
2
was observed due to the disappearance of chemical shifts at 5.01,
.47, and 5.65 ppm and the increased intensity of chemical shifts
between 0.8 and 1.8 ppm.
The branch length (defined as the average molecular weight of
the corresponding EBR branches) of the corresponding PP-g-
EBR copolymer can be determined using the following equation:
lithiation reaction. As a consequence, the branch density (defined
as the average number of branches per 10 000 carbons in the PP
backbone) and the grafting efficiency can be calculated by the
subsequent equations:
5
ꢀ
1
=2A2:67
branch density ¼
ðPP-g-PBÞ
1
=6ðA0:8ꢀ1:8 ꢀ A5:01Þ
=2A2:67
1
=2A5:01 þ 1=2A
5:47
ꢁ
EBR branch length ¼ 1
ðPP-g-PBÞ
1
10000
=6ðA0
ꢀ A5:01Þ
:8ꢀ1:8
ꢀ 1
=6A0:8ꢀ1:8
ðPPꢀBT2Þ ꢂ
2
ð9Þ
ꢀ
ꢁ
2
=
ðbranch densityÞ ꢂ 1
ꢂ 56
ꢀ
ꢁ
0000
2
grafting efficiency ¼ ðbranch densityÞ ꢂ
ð11Þ
1
0000
ꢀ
ꢁ
1
=2A2:67
All the calculated results based on the eqs 7ꢀ11 are collected in
Table 3. Upon inspecting Table 3, one can note that the PP-g-EBR
copolymers containing the same backbones differ in average
length and in average density of the EBR branches. In particular,
by comparing the samples 5 and 6, one can observe that their
branch lengths are nearly the same, but their branch densities are
quite different. The same can be noted for the samples 7 and 8.
=
ðPPꢀBT2Þ ꢂ 100%
1
=6A0:8ꢀ1:8
ð10Þ
1
Figure 4 shows a typical H NMR spectrum of the PP-g-EBR
copolymer (sample PP208-11.9g-EBR5.93 in Table 4). A com-
plete hydrogenation of PB side chains in PP-g-PB copolymers
4
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Table 4. Summary of Characterization Results by SEC-TALLS, DSC, and Rheology Measurements for PPꢀBT2 and PP-g-EBR
Copolymers
a
a
b
b
m
b
m
c
c
3
d
sample
M
w
(kg/mol) PDI (M
w
/M
n
)
EBR content (wt %)
T
c
(°C)
T
(°C) ΔH
(J/g)
X
(%) η*0.01 Hz (10 Pa s)
n
3
PPꢀBT2
208
n.m.
n.m.
252
333
n.m.
289
448
n.m.
367
534
230
1.84
0.0
3.3
92.5
142.2
142.5
142.5
142.4
141.5
142.9
142.4
141.4
140.7
141.7
139.4
72.2
74.6
74.2
68.5
64.7
72.9
59.5
53.2
45.9
57.9
42.4
65.2
100
107
105
103
105
106
108
102
100
106
99
3.1
13.6
33.9
47.5
80
0.853
0.733
0.643
0.608
0.530
0.527
0.428
0.417
0.398
0.403
0.391
0.501
e
PP208-5.07g-EBR1.45
PP208-1.17g-EBR2.18
PP208-4.49g-EBR4.03
PP208-7.46g-EBR4.76
PP208-1.78g-EBR5.88
PP208-11.9g-EBR5.93
PP208-7.51g-EBR11.6
PP208-11.6g-EBR11.7
PP208-4.04g-EBR18.0
PP208-6.17g-EBR27.8
PP208-1.19g-EBR37.8
n.m.
n.m.
1.97
2.76
n.m.
2.35
2.44
n.m.
2.37
2.64
1.88
106.5
106.1
104.6
104.8
106.8
103.6
102.1
101.6
103.5
97.1
1.2
7.9
14.5
4.7
100
24.1
27.9
36.4
24.4
40.9
17.6
187
201
262
228
402
107
104.5
142.0
c
109
a
b
Measured by SEC with light scattering detector. Determined by DSC measurements. Relative crystallinity of PP segment (%) = [ΔHPP-g-EBR
ΔHPPꢀBT2 ꢂ weight fraction of PP in the PP-g-EBR copolymer)] ꢂ 100. Power-law exponent. n.m. = not measured.
/
d
e
(
Figure 5. Molecular weight distributions of linear PPꢀBT2 and several
PP-g-EBRs determined by light scattering detector.
EF iBg Ru sr .e 6. MarkꢀHouwink plots of linear PPꢀBT2 and several PP-g-
On the other hand, the couples of samples 2 and 11, samples 4 and
Figure 6 shows the MarkꢀHouwink plots of the linear
PPꢀBT2 and several PP-g-EBR samples. The negative deviation
of the slope from linear behavior in all PP-g-EBR plots starts at
low molecular weight and gradually increases with the increase of
molecular weight, which implies a uniform branch distribution
along the PP backbone. As the branch length and the branch
density of PP-g-EBRs increase, the negative deviation of [η]
becomes more pronounced. These results demonstrate that the
higher the LCB level in a branched polymer, the larger the
negative deviation of [η].
7
, and that of samples 6 and 8 are almost identical in the branch
densities, but their EBR branch lengths are distinctly different.
.3. Characterization of PP-g-EBR by SEC with Triple
3
Detectors. Triple detection SEC can determine the absolute
molecular weight, radius of gyration (R ) and intrinsic viscosity
g
(
[η]) simultaneously. Therefore, branched polymers can be
characterized with this technique by direct application of the
36,37
ZimmꢀStockmayer approach.
A branched chain has lower
R and [η] than its linear counterpart with the same molecular
g
weight. The double-logarithmic plot of [η] versus M , generally
3.4. Compatibility of PP Backbones with EBR Branches in
PP-g-EBRs. Different from a normal LCB polymer, the chemical
structures of PP backbone and EBR branch are different. The
compatibility between both segments can be investigated by
w
referred to as the MarkꢀHouwink plot, can be used to qualita-
tively describe the LCB distribution across the molecular weight
distribution.
0
Figure 5 shows molecular weight distribution curves of linear
PPꢀBT2 and PP-g-EBR samples (determined by light scattering
detector). Compared to the linear copolymer PPꢀBT2, molecular
weights of PP-g-EBR samples were larger due to the introduction
of LCBs onto PPꢀBT2 backbones, and molecular weight dis-
tributions were broadened slightly. The polydispersity indices,
M /M , of all the resulting PP-g-EBRs remained below 3. The
means of DMA. The dependences of storage modulus (E ), loss
0
0
modulus (E ), and loss tangent (tan δ) on temperature for
PPꢀBT2 and PP208-1.19g-EBR37.8 copolymers, measured by
DMA, are shown in Figure 7a,b. Tan δ of PPꢀBT2 showed an
apparent peak at around 8 °C corresponding to glass transition
temperature (T ) of iPP in amorphous phase (Figure 7a). In
g
order to investigate the compatibility between PPꢀBT2 back-
bones and EBR branches in PP-g-EBRs, the corresponding
PPꢀBT2/EBR37.5 blend with the same EBR weight fraction
as that of PP208-1.19g-EBR37.8 copolymer was prepared by
solution mixing in toluene under an argon atmosphere. Herein
the EBR37.5 (M = 37.5 kg/mol, M /M = 1.07) copolymer
w
n
absolute molecular weights and molecular weight distributions
determined by light scattering are listed in Table 4 (The
nomenclature denotes the topology of the graft polymer which
is described as PPx-dg-EBRy: The symbol x denotes the M of
w
the backbone (kg/mol), d denotes the EBR branch density, and y
n
w
n
denotes the EBR branch length (kg/mol).)
was an ethylene/1-butene copolymer containing 62.2 mol % of
4
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Figure 8. (a) Complex viscosity and (b) storage modulus vs angular
frequency for the linear PPꢀBT2, PP-g-EBR (PP208-1.19g-EBR37.8),
and PPꢀBT2/EBR37.5 blend at 180 °C.
0
modulus (G ) as defined in eq 12 were investigated.
"
#
00ꢃ 1=2
2
ꢂ ꢃ
ꢂ
0
2
G
G
ꢃ
η ¼
þ
ð12Þ
ω
ω
3
.5.1. Comparison of Rheological Behaviors of PP-g-EBR and
PPꢀBT2/EBR Blend with the Same Composition. Figure 8 shows
0
0
Figure 7. Dependences of the mechanical storage modulus (E ), loss
the η* versus angular frequency (ω) and G versus ω at 180 °C
0
0
modulus (E ), and loss tangent (tan δ) on temperature for PPꢀBT2 (a)
for the linear PPꢀBT2, PP208-1.19g-EBR37.8, and PPꢀBT2/
EBR37.5 blend. Although the molecular weight and molecular
structure of EBR37.5 are almost identical with the branches of
PP208-1.19g-EBR37.8 copolymer, the rheological behavior of
PP208-1.19g-EBR37.8 is totally different from that of PPꢀBT2/
EBR37.5 blend. For linear PPꢀBT2, the η* becomes indepen-
dent of the frequency at low shear frequencies, in accordance
with the characteristic behavior of a viscoelastic material in the
Newtonian region. Compared to PPꢀBT2, the addition of
and PP208-1.19g-EBR37.8 (b). (c) Comparison of dependence of loss
0
0
modulus (E ) on temperature for the two samples and the correspond-
ing PPꢀBT2/EBR37.5 blend.
1
-butene, which was synthesized through two steps: first,
the anionic polymerization of 1,3-butadiene was initiated by
sec-BuLi/TMEDA (molar ratio: 1/5); then, the obtained PB
the content of 1,2-structure in the PB is 76.7 mol %) was
hydrogenated by p-toluenesulfohydrazide and tripropylamine.
Since the EBR37.5 copolymer is too viscous to be formed into a
permanent shape, its T could not be characterized by DMA. The
(
0
EBR37.5 into PPꢀBT2 leads to a slight reduction of η* and G ,
but the changing of η* exhibits a similar trend. However, the
introduction of EBR LCBs onto the PPꢀBT2 backbone signifi-
g
0
cantly enhances the η* and G at low shear frequencies and leads to
T of EBR37.5 copolymer was about ꢀ46 °C, determined by
g
more prominent shear-thinning phenomenon. The above results
show that the presence of EBR LCBs on the PP backbone strongly
changes the rheological behavior. The influences of branching
parameterson the rheology weredetailedly studied in the following
sections.
DSC measurement. Furthermore, DSC results confirmed that
the EBR37.5 copolymer was fully amorphous and rubbery. As
seen in Figure 7c, the only single T of PPꢀBT2/EBR37.5 blend
g
(
around ꢀ15 °C) was located between the T s of PPꢀBT2 and
g
EBR37.5 copolymers. The T of PP208-1.19g-EBR37.8 was the
g
3
.5.2. Effect of Branch Length on the Rheology. Figures 9ꢀ11
same as that of the PPꢀBT2/EBR37.5 blend. It should be noticed
that the molecular weight and molecular structure of EBR37.5
are almost identical with the branches of PP208-1.19g-EBR37.8
copolymer. These results suggest that the EBR copolymer
containing 62.2 mol % of 1-butene is miscible with the amorphous
0
0
show the η* versus ω, G versus ω, and G versus loss modulus
0
0
(
G ) at 180 °C for the linear PPꢀBT2 and several PP-g-EBR
samples gathered in groups of nearly identical branch densities
around 1.2, 7.5, and 11.6, respectively. We attempt to fit the
39
38
viscosity data by the Cross equation written as follows:
part of iPP, consistent with the results reported by Yamaguchi et al.
Therefore, the amorphous parts of PPꢀBT2 backbones are mis-
η0
ꢃ
η ðωÞ ¼
ð13Þ
R
cible with EBR branches in the PP-g-EBRs studied in this work.
1
þ ðλωÞ
3
.5. Effects of Branching Parameters on the Rheology.
The influences of branching parameters (branch length and
branch density) on the complex viscosity (η*) and storage
where η is the zero-shear viscosity, λ is a relaxation time which
inversely accounts for the onset of shear-thinning region, and R is
0
4
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Figure 9. (a) Complex viscosity vs angular frequency, (b) storage
modulus vs angular frequency, and (c) storage modulus vs loss modulus
for the linear PPꢀBT2 and PP-g-EBR samples (PP208-1.17g-EBR2.18
and PP208-1.19g-EBR37.8) at 180 °C.
Figure 10. (a) Complex viscosity vs angular frequency, (b) storage
modulus vs angular frequency, and (c) storage modulus vs loss modulus
for the linear PPꢀBT2 and PP-g-EBR samples (PP208-7.46g-EBR4.76
and PP208-7.51g-EBR11.6) at 180 °C.
a shear-thinning index. Nevertheless, for all the PP-g-EBR
samples, a continuous shear-thinning region can be observed in
almost all the frequency range measured, without any symptom
of leveling off at low shear frequencies due to their remarkable
different branch lengths, the value of η increases from 3.39 ꢂ
0
4
5
1
1
0 Pa s (PP208-1.17g-EBR2.18) to 1.07 ꢂ 10 Pa s (PP208-
3
3
.19g-EBR37.8). This shows that the increase of branch length
enhances the η of PP-g-EBR sample. It can also be seen that
long relaxation times. Therefore, the value of η can hardly be
0
0
increasing the branch length remarkably strengthen the shear-
thinning behavior. The n of PP-g-EBR sample decreases with the
increase of branch length. Similar results can be observed in
Figures 10a and 11a. Samples PP208-7.46g-EBR4.76 and PP208-
related to the Newtonian viscosity of the general linear visco-
elastic model. In order to correlate the rheological properties
with topological structure of PP-g-EBR, we approximated η with
0
the η* obtained at the frequency of 0.01 Hz. This parameter is
listed in Table 4. Moreover, the viscosity data at high shear
frequencies were fitted by eq 14 to quantify the shear-thinning
behavior:
7
.51g-EBR11.6 possess a similar branch density, but their branch
lengths are 4.76 and 11.6 kg/mol, respectively. The η value of
0
PP208-7.51g-EBR11.6 is as 2.5 times that of PP208-7.46g-
EBR4.76 sample. Furthermore, the η value of PP208-11.6g-
0
n ꢀ 1
ꢃ
EBR11.7 is also higher than that of the PP208-11.9g-EBR5.93
η ¼ mω
ð14Þ
sample.
0
where m is the consistency and n is the power-law exponent,
which indicates the degree of non-Newtonian behavior. The
more pronounced the shear-thinning phenomenon, the smaller
the value of n. The power-law exponents for all the samples at
Besides η , the G is also very sensitive to the presence of
0
0
LCBs. Figures 9bꢀ11b show the G versus ω for the linear
PPꢀBT2 and several PP-g-EBR samples at 180 °C. The linear
PPꢀBT2 showed a representative behavior of a viscoelastic
0
1
80 °C are also included in Table 4.
material in the terminal region. At low shear frequencies, G of
2
As shown in Figure 9a, the η of PP208-1.17g-EBR2.18 sample
PPꢀBT2 is proportional to ω . PP-g-EBR samples exhibited
0
0
is much higher than that of PPꢀBT2, meaning that the EBR
branch with the length of 2.18 kg/mol is long enough to enhance
η0. Comparing PP208-1.17g-EBR2.18 with PP208-1.19g-
EBR37.8, both having nearly the same branch density but
much higher G at low shear frequencies, and the terminal slopes
0
of G were all less than 2 ascribed to their longer relaxation times.
0
The G of PP-g-EBR samples with the same branch density
increased as the branch length increased.
4
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Figure 11. (a) Complex viscosity vs angular frequency, (b) storage
modulus vs angular frequency, and (c) storage modulus vs loss modulus
for the linear PPꢀBT2 and PP-g-EBR samples (PP208-11.9g-EBR5.93
and PP208-11.6g-EBR11.7) at 180 °C.
Figure 12. (a) Complex viscosity vs angular frequency, (b) storage
modulus vs angular frequency, and (c) storage modulus vs loss modulus
for the linear PPꢀBT2 and PP-g-EBR samples (PP208-1.78g-EBR5.88
and PP208-11.9g-EBR5.93) at 180 °C.
The nonterminal behaviors of the PP-g-EBR samples can also
for PP208-1.78g-EBR5.88), more distinct shear-thinning beha-
vior, higher storage modulus at low shear frequencies, and more
pronounced upshift from linear PPꢀBT2 in the Han plot than
the PP208-1.78g-EBR5.88 sample. Similar results can also be
obtained from the comparison between PP208-11.6g-EBR11.7
and PP208-7.51g-EBR11.6. Therefore, the rheological properties
of PP-g-EBR samples having the same branch length can be
adjusted through changing the branch density.
0
00
be illustrated in the Han plot (log G ꢀ log G , Figures 9cꢀ11c).
The Han plot has been proven to be a useful tool to investigate
the effects of LCB and polydispersity on rheological
5
,8,23,40,41
properties,
and this plot is independent of molecular
weight, only a very weak function of temperature, for high
molecular weight polymer melts with narrow molecular weight
distributions, just like the case of PP-g-EBR copolymers. As
shown in Figures 9cꢀ11c, the Han plot of PPꢀBT2 is almost a
3
.5.4. Comparison of the Effects of Branch Length and
0
00
straight line, and its formula is log G = ꢀ2.79 þ 1.59 log G ,
implying that the structure of PPꢀBT2 is linear. However, an
upshift from the linear sample PPꢀBT2 is observed for PP-g-
EBR samples at low shear frequencies, indicating that a long
relaxation mechanism occurred in these samples. Furthermore,
as the branch length of PP-g-EBR samples with almost the same
branch density increased, the up-deviation became more promi-
nent, implying that the contribution of LCB to the elasticity is
more than that to the viscosity in the PP-g-EBR samples.
Branch Density on the Rheology. From the aforementioned
experimental results, a conclusion can be drawn that increasing
the branch length and the branch density of PP-g-EBR samples
both lead to a considerable change of rheological properties.
However, the comparison of the effects of the two branching
parameters on the rheology was seldom made, which was
ascribed to the difficulty of preparing LCBPP samples with
well-defined molecular structures. Figure 14 shows the η* as a
function of ω at 180 °C for linear PPꢀBT2 and several PP-g-EBR
samples with various branch densities and branch lengths. The
branch density of PP208-1.78g-EBR5.88 is about 40% of that of
PP208-4.49g-EBR4.03, and the branch length of the former is
only as 1.46 times as that of the latter, whereas the former exhibits
3
.5.3. Effect of Branch Density on the Rheology. Figures 12
0
0
00
and 13 show the η* versus ω, G versus ω, and G versus G at
1
80 °C for the linear PPꢀBT2 and several PP-g-EBR samples
collected in groups of nearly the same branch lengths around 5.9
and 11.6 kg/mol, respectively. The sample PP208-11.9g-
a higher value of η . A similar phenomenon can also be obser-
0
EBR5.93 with higher branch density exhibited a higher η
ved by comparing PP208-1.17g-EBR2.18 with PP208-5.07g-
EBR1.45. These experimental results reveal that branch length
0
5
5
(
1.87 ꢂ 10 Pa s for PP208-11.9g-EBR5.93 and 1.0 ꢂ 10 Pa s
3
3
4
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Figure 14. Complex viscosity vs angular frequency for the linear PPꢀ
BT2 and several PP-g-EBR samples with different branch densities and
branch lengths at 180 °C.
Figure 13. (a) Complex viscosity vs angular frequency, (b) storage
modulus vs angular frequency, and (c) storage modulus vs loss modulus
for the linear PPꢀBT2 and PP-g-EBR samples (PP208-7.51g-EBR11.6
and PP208-11.6g-EBR11.7) at 180 °C.
contributes more to the increase of η than branch density does
Figure 15. Zero-shear viscosity (η
weight (Mw,LS) for PP-g-EBR copolymers at 180 °C. The arrows imply
that the actual η is higher than the plotted value η*0.01 Hz
0
) vs weight-average molecular
0
in PP-g-EBRs. Further evidence can be obtained from Figure 14b,
in which the three PP-g-EBR samples (PP208-11.9g-EBR5.93,
PP208-7.51g-EBR11.6, and PP208-4.04g-EBR18.0) possess a
similar EBR content (about 20 mol %) but differ in the branch
densities and branch lengths. From the comparison among the
0
.
three samples, one can notice that the sample with longer branch
5
length exhibited a larger value of η (2.28 ꢂ 10 Pa s for PP208-
0
3
5
4
.04g-EBR18.0, 2.01 ꢂ 10 Pa s for PP208-7.51g-EBR11.6, and
3
5
1
.87 ꢂ 10 Pa s for PP208-11.9g-EBR5.93). Therefore, we can
3
conclude that the branch length has more significant influence on
rheological properties of the PP-g-EBR samples compared to the
branch density.
The η as a function of M is a very sensitive indicator to
0
w
4
2,43
detect even low level of LCB.
For linear polymer, the η is
0
independent of its molecular weight distribution, which is an
important precondition for the interpretation of the influence of
44
LCB structure on the complex viscosity in the linear regime.
The dependence of η on M for PP-g-EBR samples at 180 °C is
0
w
Figure 16. vGP plots of the linear PPꢀBT2 and several PP-g-EBR
samples at 180 °C (vGP plots for the other PP-g-EBRs were omitted for
clarity).
shown in Figure 15. The relationship for the linear PP can be
11
described by log η = ꢀ15.4 þ 3.5 log M at 180 °C.
A
0
w
considerably increased η compared to the linear PP with the
0
same M is observed for all the PP-g-EBR samples due to the
w
formation of additional entanglements by their comblike LCBs.
structure, having the advantage of without demanding normali-
0
0
45,46
3
.5.5. Van GurpꢀPalmen Analysis. Loss angle, δ (tan δ = G /
zation with respect to the molecular weight.
The δ plotted as
0
G ), is another versatile tool for getting insight into the molecular
a function of the absolute value of complex modulus (|G*|) is the
4
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Macromolecules
ARTICLE
copolymer, suggesting that crystallization rate enhances and
the supercooling degree needed for crystallization decreases
due to the introduction of EBR LCBs onto the PP backbones.
It is very interesting to notice that the crystallization temperature
increases with the enhancement of branch length and branch
density when the PP-g-EBR samples contain a low level of LCB,
among which PP208-1.78g-EBR5.88 exhibits the highest crystal-
lization temperature. However, further elevating the LCB level
result in the reduction of crystallization temperature. The crystal-
lization process from melt state during DSC measurements was a
phase transition process without shear field. According to the
rheological data, the η of PP-g-EBRs was much higher than that
0
of PPꢀBT2 due to the formation of additional entanglements.
The entangled state of polymer chains influences the crystal-
lization process since the rearrangement of chains is necessary to
chain folding and perfection. The high viscosity is unfavorable to
disentanglement of the chains in polymer melt during crystal-
lization. The above results indicate that, on one hand, the
amorphous EBR LCBs perform the function of heterogeneous
nucleation to facilitate the crystallization of PP backbone in the
PP-g-EBR; on the other hand, the mobility and reptation ability
of PP chains are restrained by the LCBs, which hinders the
crystallization process.
Figure 17. DSC cooling curves (a) and heating curves (b) of linear
PPꢀBT2 and several PP-g-EBRs (DSC curves for the other PP-g-EBRs
were omitted for clarity).
In addition, the melting temperatures of all the samples listed
in Table 4 are around 142 °C. The similarity of melting
temperatures between PP-g-EBR copolymers and PPꢀBT2
copolymer suggests that the thickness of lamellae is almost
independent of the content of EBR branches. Though ΔHm
values of the resulting PP-g-EBRs gradually decreased with
increasing the content of EBR branches, the relative crystallinity
4
7
so-called vGP plot proposed by Van Gurp and Palmen. Trinkle
et al. applied vGP plot to characterize polydispersity of linear
4
8
polymers and classify long chain branched polymers by their
4
9
topology.
Figure 16 shows the vGP plots of the linear PPꢀBT2 and
several PP-g-EBR samples. It is known that materials are almost
completely viscous when the δ terminal value is close to 90° and
almost completely elastic when the δ terminal value is close to 0°.
The δ terminal value of linear PPꢀBT2 is close to 90°,
demonstrating that this sample is nearly totally viscous. The
sample PP208-5.07g-EBR1.45 possessed relatively high branch
density but showed similar behavior to that of PPꢀBT2, as its
EBR branches were not long enough to dramatically influence
the rheological properties. However, other PP-g-EBR samples
exhibited remarkably reduced terminal values of δ compared to
linear PPꢀBT2. Hence, the elasticity of PP-g-EBR sample is
much higher than that of linear PPꢀBT2, and the elasticity
enhances with the increase of branch length and branch density.
Furthermore, we can observe that the increase of branch density
and branch length also changes the shape of vGP plot due to the
added long-time relaxation mechanism in the viscoelastic beha-
vior. For the samples of PP208-7.46g-EBR4.76, PP208-7.51g-
EBR11.6, PP208-4.04g-EBR18.0, and PP208-11.6g-EBR11.7
with relatively high branch densities and branch lengths, a plateau
of their vGP plots started to appear and tended to become more
prominent with the increase of branch density and branch
length. This interesting phenomenon is also noticed in other
(calculated by the ΔH of the PP-g-EBR sample divided by the
m
ΔH of the PPꢀBT2 copolymer and the weight fraction of PP in
m
53
the PP-g-EBR sample), which can be used to estimate the
crystallinity of the PP part in the PP-g-EBR samples compared
to PPꢀBT2 copolymer, is nearly constant. This result demon-
strates that EBR branches grafted onto the PP backbone have no
enormous influence on the crystallinity of PP backbone at least
within 40.9 wt % of EBR content, which is in good accordance
3
8
with the results previously reported by Yamaguchi et al.
4. CONCLUSIONS
A series of PP-g-EBR graft copolymers with controlled branch
densities and branch lengths were prepared. Structureꢀproperty
relationships of PP-g-EBR graft copolymers were systematically
investigated. The resultant LCBPP consisted of an iPP backbone
and EBR LCBs. The branch density of PP-g-EBR copolymer can
be controlled by tuning the efficiency of lithiation reaction of
PPꢀBT2 backbones. On the other hand, the living 1,3-butadiene
polymerization step allows the regulation of branch length
through controlling the feed amount of 1,3-butadiene and the
reaction time.
4
1,43,50,51
LCB polyolefins.
The presence of a plateau in vGP
plot showed that these PP-g-EBR samples exhibited a physical
The DMA measurements demonstrated that the amorphous
EBR LCBs were miscible with the PPꢀBT2 backbones. Char-
acterization using SEC with triple detectors indicated that EBR
LCBs were distributed uniformly along the PP chains. The
increase of branch length and branch density both reduced the
slope of the MarkꢀHouwink plot in the high molecular weight
range. Moreover, increasing the branch length and branch
5
2
gel-like behavior.
.6. Effect of EBR Branches on Crystallization of PP Back-
3
bones in PP-g-EBRs. The influence of LCB on melt behavior
was also reflected in the transition from melt state to solid state,
i.e., crystallization process. The nonisothermal crystallization
process was analyzed by means of DSC. Figure 17 shows the
DSC curves of linear PPꢀBT2 and PP-g-EBR samples, and the
results are summarized in Table 4. The crystallization tempera-
tures of PP-g-EBRs are higher than that of linear PPꢀBT2
density both lead to increased η , more pronounced shear-
0
thinning behavior, elevated value of storage modulus at low
shear frequencies, and more significant upshift deviation from
4
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Macromolecules
ARTICLE
linear behavior in the Han plot. Branch length contributes more
to the improvement of rheological properties than branch
density does. The vGP analysis showed that PP-g-EBRs exhibited
lower loss angle compared to the linear counterpart, indicating
the enhancement of elasticity. The shape of vGP plot also
changed with the increase of branch length and branch density.
A plateau began to appear when the branching level exceeded a
certain degree, which became more prominent with the further
increase of branch length and branch density. DSC results
showed that crystallization temperatures of PP-g-EBRs were
higher than that of linear PPꢀBT2. Interestingly, the crystal-
lization temperature of PP-g-EBR increased with the enhance-
ment of branch length and branch density; however, when the
LCB level exceeded a certain degree, further elevating the LCB
level would reduce the crystallization temperature. This result
indicates that, on one hand, the amorphous EBR LCBs can act as
heterogeneous nucleating agent to accelerate the crystallization
of PP; on the other hand, the mobility and reptation ability of PP
chains are restrained bythe LCBs, which hindersthe crystallization
process. Future work will be focused on the relationships between
molecular structure and foaming behavior of PP-g-EBRs.
(17) Kolodka, E.; Wang, W. J.; Zhu, S. P.; Hamielec, A. E. Macro-
molecules 2002, 35, 10062–10070.
(
18) Cherian, A. E.; Lobkovsky, E. B.; Coates, G. W. Macromolecules
005, 38, 6259–6268.
19) Naga, N.; Shiono, T.; Ikeda, T. Macromolecules 1999, 32
348–1355.
20) Kim, I.; Shin, Y. S.; Lee, J. K. J. Polym. Sci., Part A: Polym. Chem.
000, 38, 1590–1598.
21) Agarwal, P. K.; Somani, R. H.; Weng, W. Q.; Mehta, A.; Yang, L.;
2
1
2
(
(
(
Ran, S. F.; Liu, L. Z.; Hsiao, B. S. Macromolecules 2003, 36, 5226–5235.
(22) Paavola, S.; Saarinen, T.; Lofgren, B.; Pitkanen, P. Polymer 2004
45, 2099–2110.
(23) Ye, Z. B.; AlObaidi, F.; Zhu, S. P. Ind. Eng. Chem. Res. 2004,
43, 2860–2870.
(24) Lu, B.; Chung, T. C. Macromolecules 1999, 32, 8678–8680.
(25) Langston, J. A.; Colby, R. H.; Chung, T. C. M.; Shimizu, F.;
Suzuki, T.; Aoki, M. Macromolecules 2007, 40, 2712–2720.
26) Lu, H. L.; Chung, T. C. J. Polym. Sci., Part A: Polym. Chem. 1999,
7, 4176–4183.
27) Lu, H. L.; Hong, S.; Chung, T. C. J. Polym. Sci., Part A: Polym.
(
3
(
Chem. 1999, 37, 2795–2802.
(28) Chung, T. C.; Dong, J. Y. J. Am. Chem. Soc. 2001, 123, 4871–4876.
(29) Wu, C. H.; Li, H. Y.; Feng, Y. Q.; Hu, Y. L. Polym. Int. 2007,
56, 711–717.
(
(
30) Randall, J. C. Macromolecules 1978, 11, 592–597.
31) Uozumi, T.; Soga, K. Makromol. Chem. 1992, 193, 823–831.
’
AUTHOR INFORMATION
Corresponding Author
(32) Alamo, R. G.; Mandelkern, L. Macromolecules 1991, 24
6480–6493.
*Tel þ86 (0) 431 85262004; Fax þ86 (0) 431 85262827; e-mail
(33) Simanke, A. G.; Galland, G. B.; Freitas, L.; da Jornada, J. A. H.;
ttang@ciac.jl.cn.
Quijada, R.; Mauler, R. S. Polymer 1999, 40, 5489–5495.
(34) Villar, M. A.; Failla, M. D.; Quijada, R.; Mauler, R. S.; Valles,
E. M.; Galland, G. B.; Quinzani, L. M. Polymer 2001, 42, 9269–9279.
(35) Monrabal, B.; Blanco, J.; Nieto, J.; Soares, J. B. P. J. Polym. Sci.,
Part A: Polym. Chem. 1999, 37, 89–93.
’
ACKNOWLEDGMENT
We thank the financial support from the National Natural
(36) Zimm, B. H.; Stockmayer, W. H. J. Chem. Phys. 1949, 17
Science Foundation of China for the projects (20734006,
0804045, 50873099, and 50921062) and the Ministry of Science
and Technology of China for Project No. 2008AA03Z508.
1301–1314.
2
(
(
37) Zimm, B. H.; Kilb, R. W. J. Polym. Sci. 1959, 37, 19–42.
38) Yamaguchi, M.; Miyata, H.; Nitta, K. H. J. Appl. Polym. Sci.
1
996, 62, 87–97.
(
(
39) Cross, M. M. J. Colloid Sci. 1965, 20, 417–437.
40) Vega, J. F.; Santamaria, A.; Munoz-Escalona, A.; Lafuente, P.
’
REFERENCES
(1) Small, P. A. Adv. Polym. Sci. 1975, 18, 1–31.
Macromolecules 1998, 31, 3639–3647.
(2) Santamaria, A. Mater. Chem. Phys. 1985, 12, 1–28.
(3) Gahleitner, M. Prog. Polym. Sci. 2001, 26, 895–944.
(4) Graebling, D. Macromolecules 2002, 35, 4602–4610.
(5) Tian, J. H.; Yu, W.; Zhou, C. X. Polymer 2006, 47, 7962–7969.
(6) Parent, J. S.; Bodsworth, A.; Sengupta, S. S.; Kontopoulou, M.;
(41) Wang, W. J.; Ye, Z. B.; Fan, H.; Li, B. G.; Zhu, S. P. Polymer
2004, 45, 5497–5504.
(42) Janzen, J.; Colby, R. H. J. Mol. Struct. 1999, 485ꢀ486, 569–584.
(43) Malmberg, A.; Gabriel, C.; Steffl, T.; Munstedt, H.; Lofgren, B.
Macromolecules 2002, 35, 1038–1048.
Chaudhary, B. I.; Poche, D.; Cousteaux, S. Polymer 2009, 50, 85–94.
7) El Mabrouk, K.; Parent, J. S.; Chaudhary, B. I.; Cong, R. J.
Polymer 2009, 50, 5390–5397.
8) Li, S. Z.; Xiao, M. M.; Wei, D. F.; Xiao, H. N.; Hu, F. Z.; Zheng,
A. N. Polymer 2009, 50, 6121–6128.
9) Zhang, Z. J.; Xing, H. P.; Qiu, J.; Jiang, Z. W.; Yu, H. O.; Du,
X. H.; Wang, Y. H.; Ma, L.; Tang, T. Polymer 2010, 51, 1593–1598.
10) Yoshii, F.; Makuuchi, K.; Kikukawa, S.; Tanaka, T.; Saitoh, J.;
Koyama, K. J. Appl. Polym. Sci. 1996, 60, 617–623.
11) Auhl, D.; Stange, J.; Munstedt, H.; Krause, B.; Voigt, D.; Lederer,
A.; Lappan, U.; Lunkwitz, K. Macromolecules 2004, 37, 9465–9472.
12) Krause, B.; Stephan, M.; Volkland, S.; Voigt, D.; Haussler, L.;
Dorschner, H. J. Appl. Polym. Sci. 2006, 99, 260–265.
13) Weng, W. Q.; Hu, W. G.; Dekmezian, A. H.; Ruff, C. J.
Macromolecules 2002, 35, 3838–3843.
14) Ye, Z. B.; Zhu, S. P. J. Polym. Sci., Part A: Polym. Chem. 2003, 41
152–1159.
15) Shiono, T.; Azad, S. M.; Ikeda, T. Macromolecules 1999,
2, 5723–5727.
16) Weng, W. Q.; Markel, E. J.; Dekmezian, A. H. Macromol. Rapid
Commun. 2001, 22, 1488–1492.
(44) Gabriel, C.; Munstedt, H. Rheol. Acta 2002, 41, 232–244.
(45) Lohse, D. J.; Milner, S. T.; Fetters, L. J.; Xenidou, M.;
Hadjichristidis, N.; Mendelson, R. A.; Garcia-Franco, C. A.; Lyon,
M. K. Macromolecules 2002, 35, 3066–3075.
(46) Stadler, F. J.; Piel, C.; Klimke, K.; Kaschta, J.; Parkinson, M.;
Wilhelm, M.; Kaminsky, W.; Munstedt, H. Macromolecules 2006,
39, 1474–1482.
(47) Van Gurp, M.; Palmen, J. Rheol. Bull. 1998, 67, 5–8.
(48) Trinkle, S.; Friedrich, C. Rheol. Acta 2001, 40, 322–328.
(49) Trinkle, S.; Walter, P.; Friedrich, C. Rheol. Acta 2002, 41
103–113.
(50) Wood-Adams, P. M.; Dealy, J. M.; deGroot, A. W.; Redwine,
O. D. Macromolecules 2000, 33, 7489–7499.
(51) Robertson, C. G.; Garcia-Franco, C. A.; Srinivas, S. J. Polym. Sci.,
Part B: Polym. Phys. 2004, 42, 1671–1684.
(52) Garcia-Franco, C. A.; Srinivas, S.; Lohse, D. J.; Brant, P.
(
(
(
(
(
(
(
(
1
Macromolecules 2001, 34, 3115–3117.
(53) Chung, T. C.; Rhubright, D. Macromolecules 1994, 27, 1313–1319.
(
3
(
4
179
dx.doi.org/10.1021/ma200604y |Macromolecules 2011, 44, 4167–4179