Synthesis of PLLA-based block copolymers for improving melt strength and toughness of PLLA by in situ reactive blending

Accepted Manuscript Synthesis of PLLA-based block copolymers for improving melt strength and toughness of PLLA by in situ reactive blending Bao Zhang, Xinchao Bian, Sheng Xiang, Gao Li, Xuesi Chen PII:

S0141-3910(16)30357-3

DOI:

10.1016/j.polymdegradstab.2016.11.022

Reference:

PDST 8120

To appear in:

Polymer Degradation and Stability

Received Date: 12 September 2016 Revised Date:

18 November 2016

Accepted Date: 23 November 2016

Please cite this article as: Zhang B, Bian X, Xiang S, Li G, Chen X, Synthesis of PLLA-based block copolymers for improving melt strength and toughness of PLLA by in situ reactive blending, Polymer Degradation and Stability (2016), doi: 10.1016/j.polymdegradstab.2016.11.022. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

ACCEPTED MANUSCRIPT Synthesis of PLLA-based block copolymers for improving melt strength and toughness of PLLA by in situ reactive blending

RI PT

Bao Zhang, Xinchao Bian*, Sheng Xiang, Gao Li*, Xuesi Chen

（Key Laboratory of Polymer Ecomaterials, Changchun Institute of Applied Chemistry, Chinese Academy of Sciences, Changchun, 130022）

SC

*Corresponding author. Email: [email protected]; [email protected]

M AN U

ABSTRACT

Low melt strength and toughness of poly(L-lactide) (PLLA) limited its large-scale application. In this work, a facile method was proposed and demonstrated to a feasible route to solve these problems. A series of PLLA-based block copolymers

TE D

PLLA-block-poly(butylene succinate)-block-PLLA (PLLA-b-PBS-b-PLLA), and chain extender (PLLA-block-poly(glycidyl methacrylates))3 (PLLA-b-PGMA)3 were synthesized and used for the improvement of the melt strength and toughness of

EP

PLLA by in situ reactive blending. The structure and composition of the copolymers

AC C

were confirmed by nuclear magnetic resonance spectra, infrared spectra and gel permeation chromatography. The elongation at break increased from 4.2% for PLLA to 234% for the blend containing 40% block copolymer and 5% chain extender, and remarkably maintained their strength. Rheological analysis showed that the blends exhibited the strong strain-hardening behavior. Measurements of the linear viscoelastic properties of the melt blends suggested that the chain extender promoted the development of chain branching. DSC data showed that the crystallization was not perfect after addition of PLLA-b-PBS-b-PLLA and (PLLA-b-PGMA)3 compared with

ACCEPTED MANUSCRIPT neat PLLA. SEM measurements revealed the improved interface adhesion of the blends. The introduction of PLLA-b-PBS-b-PLLA and (PLLA-b-PGMA)3 imparted both high toughness and high melt strength to the PLLA.

RI PT

Keywords: Polylactide; Melt strength; Toughness; Reactive Blending; Compatibility

1. Introduction

The increasing concerns on environmental impact and sustainability of

SC

conventional petroleum-based polymers have motivated researchers to devote considerable efforts to develop thermoplastic polymers from renewable resources

M AN U

[1,2]. Among a few currently commercially available biobased polymers, Poly(L-lactide) (PLLA) is the most promising candidate and widely investigated [3-5]. PLLA is well-known biodegradable, biocompatible polyesters produced from annually renewable natural resources. With the good mechanical properties as thermal

TE D

stability, processability and biodegradability, PLLA has potential to replace petroleum-based plastic products as polypropylene from both environmental and economic perspectives [6-8]. However, the low melt strength and week toughness that

EP

substantially impedes its development in many potential applications, particularly for

AC C

packaging industries [9]. Many approaches have been taken to improve these properties

of

PLLA,

including

copolymerization,

polymer

blending,

and

branching/cross-linking reactions [10-13]. Copolymerization has been extensively investigated as a powerful means to obtain

polymer materials with improved properties compared with homopolymers [14]. Properties including tensile and reological performances of a copolymer can be tailored in versatile ways by manipulating the architectures of the molecule, sequence of monomers, and composition. Copolymerization of PLLA can be conducted either

ACCEPTED MANUSCRIPT through polycondensation of lactic acid with other monomers (or polymer segments) or the ring-opening copolymerization (ROC) of LA with other cyclic monomers. Because the latter synthesis route gives a more precise control of chemistry and higher molecular weight of copolymers, it is more widely used to improve the toughness of

RI PT

PLLA in the literatures [15].

Blending PLLA with immiscible or miscible polymer with higher melt viscosity is an effective and convenient method to improve the toughness of PLLA [16-21]. PLLA

polymer

blends

with

PLLA

have

been

SC

has been blended with various polymers to improve its toughness. Many studies on investigated

on

completely

poly(hydroxyalkanoate)s poly(butylene

(PHA)

M AN U

degradable/renewable resource polymers, including polycaprolactone (PCL) [16], [17],

adipate-co-terephthalate)

poly(butylene (PBAT)

succinate) [19] ,

(PBS)

[18],

poly(butylene

succinate-co-adipate) (PBSA) [20], and poly(butylene succinate-co-L-LA) (PBSL)

TE D

[21]. However, the simple melting blend of PLLA with these polymers usually led to a marginal improvement in toughness because of the low miscibility between the components of the blends.

EP

Reactive blending was considered to be an effective method for improving the

AC C

miscibility of PLLA blends [22,23]. By this way, in situ compatibilization was achieved in PLLA blends [24-26]. More importantly, the branching/cross-linking reactions occurred among the blends with the increase of molecular weight, and the branching structure was an efficient approach to improve the melt rheological properties of PLLA blends. It had been widely reported that the branching structure had a significant effect on physical properties of polymers, which showed that even a small amount of branching component led to great changes on melt rheological, thermal and mechanical properties [23], as well as the biodegradability of polymers.

ACCEPTED MANUSCRIPT Various radical initiators had been introduced into PLLA/PBAT [24], PLLA/PCL [25], and PLLA/PBS [26] blends to form cross-linked and/or branched structures by heterogeneous and/or homogeneous radical coupling reactions. The copolymers formed between the two phases by heterogeneous reactions could improve the

RI PT

compatibility of the polymer blends and, therefore, to enhance the mechanical properties of the blends. Branching structures can be introduced, but it was eventually hard to get control on the molecular weight and topological structure of the final

SC

products due to the randomness of free radical branching and the high reactivity of peroxides.

M AN U

Incorporating functional groups onto the PLA backbone tended to more easily obtain branching with “controlled” structures. A many kinds of multifunctional monomers such as diisocyanates, acid anhydride, and epoxy were used effectively to produce branching of PLLA with high molecular weight and high melt strength

TE D

[27-29]. Although the diisocyanates increased the molecular weight of PLLA, they also had inherent toxicity issues [27]. Acid anhydride was one of the most widely used reactive compatibilizers due to its good chemical reactivity, low toxicity and low

EP

potential to polymerize itself [28]. In this process, the bridge could be formed between

AC C

grafted maleic anhydride (MAH) and polymer chains, and the interfacial adhesion was effectively improved. However, the residue MAH will damage the life stability of the PLLA products and the aging performance of the product will be reduced. The use of an epoxy-functionalized chain extender was able to enhance the

interfacial adhesion between the two components and form the branched structures, and the blends showed obvious strain hardening behavior. [29,30] Dong et al.[31] blended PLA/PBAT in the presence of epoxy-functionalized chain-extenders, and the compatibility between the PLLA and PBAT was significantly improved, resulting in

ACCEPTED MANUSCRIPT an enhanced toughness of the blends, e.g., the elongation at break was increased to 500% without any decrease in the tensile strength. Al-Itry et al.[32] used various amounts of chain-extenders agents for PLA/PBAT blends. The results revealed that the thermal stability was improved greatly, and the phase dispersion of the blends was

RI PT

improved, as the viscosity and storage modulus was also improved compared to the unmodified samples. The properties increased more pronounced as the concentration of chain-extenders increases.

SC

This study is aimed to establish a method to produce high melt strength and toughness of PLLA blending resin. Firstly, the high toughness PLLA-b-PBS-b-PLLA

M AN U

block copolymer was synthesized. The coexistence of the PLLA segment and flexible PBS segment in block copolymer was expected to improve compatibility and flexibility of the PLLA blends. Secondly, a chain extender (PLLA-b-PGMA)3 were synthesized and melted by blending with PLLA/PLLA-b-PBS-b-PLLA. The blending

TE D

resin not only had high enlongation at break, but also had high melt strength. Although much work has been done to improve the toughness of PLLA, the products was usually accompanied by a great loss in strength. Thus, how to greatly enhance

EP

toughness while minimizing the reductions in strength of the PLLA materials still

AC C

remained a challenge. This study illustrated a new approach to achieve PLLA blends with better toughness and melt strength, and the blends remarkably maintained high tensile strength.

2. Experimental section 2.1 Materials Tetra-n-butyl-titanate (Ti(OBu)4, Sigma-Aldrich), Stannous octoate (Sn(Oct)2, Sigma-Aldrich), and α-bromoisobutyryl bromide (Aldrich, 99%) were used without

ACCEPTED MANUSCRIPT further purification. Glycerol (Sigma-Aldrich) was used as initiators and dried under vacuum before use. L-lactide (LLA) was purchased from Purac and recrystallized in ethyl acetate for three times. Succinic acid (SA) (Beijing Chemical Co.), 1,4-butanediol (1,4-BD) (Beijing Chemical Co.), and 2, 2´-Bipyridine (bpy, Beijing

RI PT

Chemical Co.) were used without further purification. Glycidyl methacrylates (GMA, Sigma-Aldrich) were distilled with calcium hydride (CaH2) under vacuum before use. Copper(I) chloride (CuCl, Beijing Chemical Co.) was purified by precipitation from

SC

acetic acid to remove Cu2+, filtrated and washed with ethanol and then dried. Dichloromethane (Tianjin Chemical Co.) were dried with CaH2 and distilled.

M AN U

N,N`-dimethylformamide (DMF) was dried with CaH2 and distilled under vacuum before use. Triethylamine (TEA, Beijing Chemical Co) was refluxed for 12 h in the presence of CaH2 and distilled. A semicrystalline grade PLLA was supplied by Zhe-jiang Hisun Biomaterials Co., Ltd (China). The number average molecular

TE D

weight(Mn) and polydispersity index were 8.21 104 g/mol and 1.32, respectively. All the reagents used in this study were analytical-grade. 2.2 Synthesis of poly(butylene succinate) (PBS)

EP

The PBS was prepared through a two-step reaction of esterification and

AC C

polycondensation in sequence (Scheme 1). The reactor was a 500 mL four-necked flask with a mechanical stirrer, thermometer, nitrogen inlet, and a condenser in a temperature controlled oil bath. The monomer succinic acid (118 g, 1 moL) and 1,4-butanediol(99 g, 1.1 moL) with designed mole ratio (1 : 1.1) were added into the flask. The reaction mixture was heated to 160 oC with stirring at a constant speed, and water resulted from the reaction was removed by distillation. When no more water was distilled out under the atmospheric pressure, the catalyst Ti(OBu)4 (0.68 g, 0.002 moL) was added, and the reaction was carried out under high vacuum. The

ACCEPTED MANUSCRIPT temperature was gradually raised to the final reaction temperature of 230oC and maintained at this condition for 3 h. The highly viscous product was cooled in the flask under nitrogen atmosphere, and dissolved in CHCl3. The solution in the flask

filtration and dried in vacuum at 50 oC overnight. 2.3 Synthesis of PLLA-b-PBS-b-PLLA block copolymer

RI PT

was poured into cold alcohol to precipitate the PBS. The product was isolated by

The procedure was exemplified for the sample PLLA-b-PBS(25%)-b-PLLA and was

SC

carried out in analogous way for all other PLLA-b-PBS-b-PLLA (Scheme 1). The monomer LLA (144 g, 1 mol) and initiator PBS (48g, 2.4 10-3 mol) was transferred

M AN U

into an oven-dried 500 ml reaction vial under an inert atmosphere of dry argon, and the vial was immediately sealed with a rubber septum. The reaction vial was sealed and degassed three times by freeze-pump-thaw cycles. The vial was then placed into a constant temperature (130 oC) oil bath with magnetic stirring. The Sn(Oct)2 (1 mL,

TE D

0.1 g/mL) were consecutively added via a gastight syringe under argon into the reaction vial after the LLA and PBS was melted homogeneously. All the procedures were carried out as rapidly as possible to avoid the entry of vapor. The reaction

EP

mixture was allowed to cool down to room temperature after 24 h. The solid product

AC C

was dissolved in CHCl3 and poured into cold alcohol to precipitate the PLLA-b-PBS-b-PLLA block copolymer. The product was isolated by filtration and dried in vacuum at 50 oC for 48 h. The yield was 94%. 2.4 Synthesis of 3-arm poly(L-lactide) (PLLA) The synthesis of 3-arm PLLA was according to our previous reports.[33,34] The monomer LLA (144 g, 1 mol) was transferred into an oven-dried 500 ml reaction vial under an inert atmosphere of dry argon, and the vial was immediately sealed with a rubber septum. The initiator glycerol (1.84 g, 0.015 mol) and Sn(Oct)2 (1.44 mL, 0.1

ACCEPTED MANUSCRIPT g/mL) were consecutively added via a gastight syringe under argon into the reaction vial. The vial was then placed into a constant temperature (120 oC) oil bath with magnetic stirring. All the procedures were carried out as rapidly as possible to avoid the entry of vapor. The reaction mixture was allowed to cool down to room

RI PT

temperature after 24 h. The solid product was dissolved in CHCl3 and poured into cold alcohol to precipitate the multi-arm PLLA. The product was isolated by filtration and dried in vacuum at 60 oC for 48 h. The yield was 97%.

SC

2.5 Synthesis of macroinitiator

3-arm PLLA (50 g, 5.3 mmol) was dissolved in 200 mL of dry dichloromethane and

M AN U

then cooled in an ice bath (0 oC). To this solution was added 10 mL (100 mmol) of triethylamine. After 5 min of stirring, 10 mL of a dichloromethane containing α-bromoisobutyryl bromide (4.6 g, 20 mmol) was added dropwise to the solution over a period of 1 h. The reaction was carried out at 0 oC for 2 h and then at room

TE D

temperature for 22 h under stirring continuously. The solution was filtrated to remove the quaternary ammonium halide (CH3CH2)3NH+Br-. The filtrate was concentrated and then precipitated in alcohol. The yield was 95%.

EP

2.6 Synthesis of the 3-arm star-block copolymer (PLLA-b-PGMA)3

AC C

A dry flask equipped with a magnetic stirrer was charged with CuCl (0.9 g, 0.027 mol), bpy (4.2 g, 0.08 mol), and macroinitiator 3-arm PLLA (25 g, 0.0025 mol). The reaction vial was sealed and degassed three times by freeze-pump-thaw cycles. Pre-degassed solvent DMF (100.0 mL) and monomer GMA (28 g, 0.2 mol) by argon were introduced into the flask via an Ar-washed syringe. Subsequently, the reaction flask was immersed in a constant temperature (80 oC) oil bath for the 12 h. The reaction was rapidly terminated in an ice bath. The catalyst was removed by passage of the polymer solution through an aluminum oxide column. The crude polymer

ACCEPTED MANUSCRIPT solution was precipitated in alcohol, and then dried under vacuum overnight. The yield was 50.4%. 2.7 Sample Preparation PLLA, PLLA-b-PBS-b-PLLA, and (PLLA-b-PGMA)3 were dried at 60 oC under

RI PT

vacuum for 12 h before use. All the samples were prepared by melt mixing using a Torque Rheometer (XSS-300) at the rotary speed of 32 rpm for 7 min at 180 oC, then followed by compression molding under 10 MPa pressure for 3 min at 180 oC to

room

temperature

in

another

compression

SC

obtain the sheet samples with the thickness of 1 mm. The samples were then cooled to molding

machine

for

further

mixed before melt blending. 2.8 Characterization

M AN U

characterizations. All compositions made were on weight/weight basis, and manually

Nuclear magnetic resonance (NMR) spectra were recorded on a Bruker AV 400 NMR

TE D

spectrometer in chloroform-d (CDCl3).

Infrared spectra were recorded on a Bruker Vertex 70 FTIR spectrophotometer, and 50 scans were collected with a spectral resolution of 1 cm 1. The chloroform solutions −

EP

containing the samples were cast onto a KBr disk, allowed to evaporate at room

AC C

temperature, and then dried under a vacuum at 50 °C for 24 h. Number- and weight-average molecular weights (Mn, Mw) and molecular weight distributions (polydispersity index, D=Mw/Mn) were determined by gel permeation chromatography (GPC), using a series of linear Styragel columns (HT2 and HT4) and a Waters 410 HPLC pump, with an Waters 2414 RI Detector. The eluent was chloroform at a flow rate of 1.0 mL min-1 at 25°C. Conventional calibrations were performed using monodispersed polystyrene standards. Differential scanning calorimetry (DSC) was carried out on a TA Instrument

ACCEPTED MANUSCRIPT DSC-Q10 to study the thermal properties of the polymers at a heating and cooling rate of 10 oC/min under a nitrogen flow of 200 mL/min. A polymer (about 3.0 mg) was loaded in a cell, and the heat exchange was recorded during the heating and cooling cycles.

RI PT

Rheological studies of neat PLLA and its blends were conducted with an ARES parallel plate rheometer (Anton-Paar,Physica MCR301) under a nitrogen atmosphere to avoid thermal oxidative degradation. A dynamic frequency sweep test was

SC

performed to determine the viscoelastic properties of blends. The plates had 25 mm diameter. The gap between plates was 1 mm during testing. The test was conducted in

M AN U

the frequency range of 0.1 to 100 rad/s at a strain rate of 1% and at a testing temperature of 180oC.

Uniaxial elongational viscosity measurements were carried out on an ARES rheometer (Anton-Paar,Physica MCR301) at 178.5 °C with the extensional viscosity

TE D

fixture (EVF) at constant strain rates of 0.3 s−1. Pre-elongation was for 6 s applied before the measurements to ensure that no slipping occurred between the sample and the fixture. The sample sheets were cut into pieces with a width of 5 mm and a length

EP

of 20 mm for the uniaxial elongational viscosity measurements.

AC C

The tensile properties were measured using an Instron 1211 machine (Instron Co., UK) at a crosshead speed of 5 mm/min. The dumbbell-shaped samples for tensile testing were cut from the compression molded sheets, and the dimension of the samples was 20 4 1 mm3. An average value of five replicated measurements was taken for each sample. The field-emission scanning electron microscope (FE-SEM, XL30) was used to observe the morphologies of the fractured surface of the PLLA blends. The PLLA blends were quenched into liquid nitrogen and then fractured. The fractured surfaces

ACCEPTED MANUSCRIPT were sputtered with gold before observation and an accelerating voltage of 15 kV was used to produce the SEM photographs. H O (CH2)4OOC(CH2)2CO

HO(CH2)4OH + HOOC(CH2)2COOH

O(CH2)4 n

O

OH

RI PT

O

Sn(Oct)2

O

O

CH3 C O

p

O(CH2)4OOC(CH2)2CO

O(CH2)4O n

C

O

CH O H p

SC

H O CH

CH3

3. Results and discussion

M AN U

Scheme 1. Synthesis of PLLA-b-PBS-b-PLLA.

3.1 Synthesis of PLLA-b-PBS-b-PLLA block copolymer

PBS was used as the initiator for the PLLA-b-PBS-b-PLLA block copolymer and was firstly synthesized, as shown in Scheme 1. The 1H NMR was used to study the

TE D

chemical structure and composition of PBS (Fig.1(A)). The peak appearing at 2.63 ppm (Proton c) corresponded to the methylene protons in SA units. The peaks at 1.71 ppm (Proton a) and 4.12 ppm (Proton b) originated from the central and terminal

EP

protons of 1,4-BD segments, respectively. Moreover, the ratio of integrated areas of

AC C

peaks a, b, and c was close to 1:1:1, corresponding to those of BS unit in equimolar ratio. The FTIR spectrum of PBS was shown in Fig. 2. The intensive absorption bands appeared in the wave number region of 1726 cm-1 and 1170 cm-1 assigned to the ester carbonyl group of the PBS main chains. The characteristic bands at 2956 cm-1 were assigned to the stretching of methylene group in the PBS blocks; the weak absorption bands at the wave numbers of approximately 3558 cm-1 ascribed to the hydroxyl group in the end of the PBS. The appearance of these characteristic peaks could be taken as indicative of successful PBS synthesis. At the same time, the GPC analysis of

ACCEPTED MANUSCRIPT PBS showed a unimodal and symmetrical trace (Fig. 3). The molecular weight is 17000 g/mol determined by GPC. All above data indicate that biodegradable PBS polyesters were synthesized successfully. a

c

b

C H 2 C H 2 C H 2 C H 2 O O C (C H 2 ) 2 C O O

(C H 2 ) 4 n

OH

RI PT

HO

a` CH3 H

O

CH

b`

C O

b

c

p

(C H 2 ) 4O n

C

O

CH3 CH

O

p

H

M AN U

(B )

a

O C H 2 C H 2 C H 2 C H 2 O O C (C H 2 ) 2 C O O

SC

(A )

Fig. 1. 1H NMR spectra of PBS (A), and PLLA-b-PBS-b-PLLA (PBS-25%) (B) were recorded at

AC C

EP

TE D

room temperature in CDCl3.

Fig. 2. IR spectrum of PBS.

The triblock copolymer PLLA-b-PBS-b-PLLA was synthesized by ROP of LLA in bulk at 130 oC in the presence of PBS. The general synthetic route used for the preparation of the PLLA-b-PBS-b-PLLA was shown in Scheme 1. Figure 1B showed

ACCEPTED MANUSCRIPT a typical 1H NMR spectrum of the PLLA-b-PBS-b-PLLA (PBS-25%). It showed that in addition to the dominant PBS signals a-c, the doublet signal a` at 5.2 ppm was assigned to the methine protons of PLLA main chain, and the quartet signal at 1.4 ppm were associated with the methyl protons of PLLA main chain. In addition, the

formation

of

additional

PLLA

block.

The

RI PT

molecular weight was increased after the polymerization (Fig.3), attesting the compositions

PLLA-b-PBS-b-PLLA copolymer were provided in Table 1.

of

PBS

and

SC

Table 1 Results for PBS, PLLA-b-PBS-b-PLLA block copolymers. Monomerc) mPBS/mLA

PBS

－

PBS-5%

5/95

PBS-15%

15/85

PBS-25%

25/75

PBS-35%

35/65

conversion

[BS]/[LA]b)

(g/mol)

Mw/Mna)

M AN U

Sampled)

Mn,GPCa)

－

－

17000

1.7

90%

1/44

83000

1.5

92%

1/13.5

71000

1.4

94%

1/7.5

62000

1.5

96%

1/4.7

50000

1.6

TE D

a) Determined by GPC measurements.

b) Degree of polymerization of BS/LA calculated from the 1H NMR spectra. c) The conversion was determined gravimetrically. d) PBS-5% represent the copolymer PLLA-b-PBS-b-PLLA in which the contents of PBS

AC C

EP

was 5%. The others were explained as above.

Fig. 3. GPC traces of PBS and PLLA-b-PBS-b-PLLA. PBS (Mn=17000, PDI=1.7), PLLA-b-PBS-b-PLLA (Mn=62000，PDI=1.5).

ACCEPTED MANUSCRIPT 3.2 Synthesis of the 3-arm block copolymer (PLLA-b-PGMA)3 3-arm PLLA was synthesized by the ROP of LLA (Scheme 2). The structure of 3-arm PLLA was confirmed by 1H NMR spectrum (Fig. 4(A)). The doublet signal a at 5.2 ppm corresponded to the methine protons of PLLA main chain, and the quartet

RI PT

signal b at 1.6 ppm corresponded to the methyl protons of PLLA main chain. The terminal methine adjacent to the OH groups led to resonances at 4.2 ppm, and the terminal methyl protons adjacent to the OH groups led to resonances at 1.4 ppm. 13

C NMR spectrum (Fig. 5), the signal e at 67.1 ppm was assigned to the

SC

From the

methylene carbons of initiator, and the methine protons d in the initiator was

M AN U

overlapped with the methine protons b in the 3-arm PLLA main chain and appeared at 70 ppm. The signal c at 171.2 ppm represented the carboxyl carbons of 3-arm PLLA main chain, and the signal a at 16.6 ppm was corresponded to methyl carbons of 3-arm PLLA. The coexistence of the characteristic peak of initiator and PLLA further

TE D

confirmed that glycerol initiated ROP of LLA successfully. O

OH H

O

O

H

n

O

+

O

O

O

CH3

O

O

AC C

n

C

CH

a

O

n

C

n

c CH2 C m k O C e

O

c

CH

O

a

n

H

O Br Br b

O

GMA CuCl/bpy

O

C O

0℃ CH2Cl2

d CH3

f O

CH3 O

O

b

b

O

Sn(Oct)2 120 oC

EP

O

O

CH3 O

n

C O

CH

a

O O

n

C

Br k

g

Scheme 2. Synthesis of 3-arm block copolymer (PLLA-b-PGMA)3.

The 3-arm PLLA macroinitiator (Br)3-PLLA for ATRP was synthesized by an esterification reaction between terminal OH group of the 3-arm PLLA with α-bromoisobutyryl bromide. In order to avoid cleavage of the polymer chain, the reaction was carried out at 0 oC in dried CH2Cl2 in the presence of triethylamine (TEA). The catalyst TEA absorbed HBr from the solution to generate a precipitate of

ACCEPTED MANUSCRIPT -

quaternary ammonium halide (CH3CH2)3NH+Br , which benefited the esterification. The chemical structure of resulting macroinitiator was confirmed by 1H NMR analysis (Fig. 4(B)). The new signals k appearing at 1.95 ppm were assigned to the methyl protons adjacent to the active bromide, which indicated that the α-bromoester groups

RI PT

was attached to the PLLA chain end. The signal c around 2.65 ppm disappeared completely after the esterification, which indicated the complete substitution of the end hydroxyl groups. Besides, a narrow symmetrical GPC signal for macroinitiator

SC

was also observed in Fig. 6(b).

The ATRP of GMA from macroinitiator (Br)3-PLLA was carried out in DMF at 80 C with CuCl/bpy as the catalyst system (Scheme 2). From the 1H NMR spectrum of

M AN U

o

the 3-arm block copolymer (Fig. 4(C)), the signals at 4.35 ppm and 3.78 ppm were due to the splitting of methylene protons in the -CH2OCO- group of the GMA unit by the methine proton of the epoxy group. The peaks at 3.22 ppm and 2.62, 2.82 ppm

TE D

were assigned to the methine proton and methylene protons of the epoxy group, respectively. The chemical shifts at 0.9-1.4 ppm were ascribed to the α-CH3 protons of the PGMA block. The appearance of these new characteristic peaks could be taken

EP

as indicative of successful ATRP initiation. The evolution of the GPC curves was

AC C

shown in Fig. 6(c). No obvious shoulders and tailings were observed, indicating a high

initiation

efficiency

and

negligible

radical-radical

coupling

in

the

polymerization. The GPC curves shifted left with the monomer conversion, indicating a gradual increase in the molecular weight. The unimodal and symmetrical shape on the GPC plot of the block copolymer suggested the absence of a homopolymer composed of either PLLA or PGMA and the complete initiation of the macroinitiator during the ATRP process.

ACCEPTED MANUSCRIPT

(A) b

O

b`

CH3

O

O

C

CH3

CH

O

a

n

a`

O

b

O

CH3

O

O

C

CH

a

O

O

n

C

Br

k

M AN U

O

SC

(B)

TE D

(C)

b

O

O

CH3

O

C

AC C

EP

O

Fig. 4.

c

OH

CH

RI PT

O

C

CH

a

d CH3

O O

n

C

c CH2 C m k O C O f O

1

e g

H NMR spectra of 3-arm PLLA (A) macroinitiator (B) and block copolymer

(PLLA-b-PGMA)3 (C) was recorded at room temperature in CDCl3.

ACCEPTED MANUSCRIPT

a

O O

d

e

c

C

CH

O

b

O

n

H

SC

RI PT

e

O

CH3

TE D

M AN U

Fig. 5. 13C NMR spectrum of 3-arm PLLA was recorded at room temperature in CDCl3.

EP

Fig. 6. GPC traces of 3-arm PLLA (a), macroinitiator (b), and block copolymer (PLLA-b-PGMA)3 (c). (a) Mn=9200，PDI=1.51; (b) Mn=9500，PDI=1.42; (c) Mn=15000，PDI=1.38.

AC C

3.3 Mechanical and rheological properties The composition of block copolymers had great influence on their mechanical

properties. Tensile testing can give many parameters of mechanical properties such as tensile strength, elongation at break, Young’s modulus and toughness. The stress-strain curves of the PLLA-b-PBS-b-PLLA block copolymers with different compositions were shown in Fig.7. No yielding happened for neat PLLA during the course of stretching, and the elongation at break was only 4.2%. In contrast, all the block copolymers yielded better elongation at break except PBS-5%. When the

ACCEPTED MANUSCRIPT weight fraction of PBS reached 25% (PBS-25%), the strain at the break increased to 380%, which was much higher than that of neat PLLA. Such results were greatly consistent with our expectation to improve toughness of PLLA with flexible PBS in the block copolymers. In addition to the elongation at break, both of the tensile

RI PT

strength and Young’s modulus were reduced after the introduction of PBS block to the PLLA block, and the values were decreased with the increase of PBS weight fraction. However, it was worth noting that when the weight fraction of PBS increased to 35%

SC

(PBS-35%), the elongation at break increased slightly in compare with PBS-25%, and the tensile strength was further decreased. So in the following study, PBS-25% was

TE D

M AN U

selected for the modification of PLLA.

EP

Fig. 7. Stress-strain curves obtained at a cross-head speed of 20 mm/min for the PLLA-b-PBS-b-PLLA with different PBS contents: (a) neat PLLA; (b) 5%; (c) 15%; (d) 25%; (e) 35%. The right curves gave details of stress-strain of the blends in the neighborhood of yield

AC C

points.

The tensile measurement of the PLLA/PBS-25% blends with different ratios was

first studied to observe the effect of PBS-25% on those mechanical parameters of PLLA. The stress-strain curves of the PLLA/PBS-25% blends were shown in Fig.8, and the details of the tensile properties were given in Table 2. From Table 2, it can be found that the fracture behavior of the specimens in the tensile tests changed from brittleness of neat PLLA to ductile fracture of PLLA/PBS-25% blends.

RI PT

ACCEPTED MANUSCRIPT

SC

Fig. 8. Stress-strain curves obtained at a cross-head speed of 20 mm/min for the PLLA/PBS-25% blends with different PBS-25% contents.

PLLA/PBS-25%

Elongation at break (%)

M AN U

Table 2. Mechanical properties of PLLA/PBS-25% for 5 times. Tensile strength

Tensile

Toughness

modulus (MPa)

(MPa)

72.7±1.5

2110±300

192±10

67.0±2.0

1980±160

303±90

54.8±2.0

1760±200

524±220

49.7±3.2

1630±350

903±303

at yield (MPa)

4.2±0.5

90/10

7.8±2.5

80/20

14.0±6.0

70/30

29.4±10

60/40

81.4±20

47.3±3.5

1420±300

1974±480

94.2±30

41.8±4.1

1150±280

2270±680

50/50

TE D

100/0

EP

The toughness exhibited significant increase after addition of PBS-25%, and increased with the PBS-25% contents. The elongation at break of PLLA blends was

AC C

improved dramatically with the increase of PBS-25% content, from 4.2% for neat PLLA to 29% for the blend containing 30% PBS-25%, and achieved a sharp increase when the weight fraction of PBS-25% reached 40%. However, with further increase of PBS-25% contents to 50%, the elongation at break was not improved obviously, but the stress strength was lowered greatly. Considering the loss in strength accompanied by the block copolymer, the appropriate contents of copolymer should be 40% in the PLLA blends. The great improvement of strain at the break of the blend was ascribed to the intrinsic high flexibility of the PBS-25%, and the elongation at

ACCEPTED MANUSCRIPT

M AN U

SC

RI PT

break did not improve greatly when PLLA blends had no enough flexible elements.

Fig. 9. Frequency dependence of A complex viscosity (η*), B storage modulus (Gʹ ) and C loss

(d) 60/40/2;(e) 60/40/5.

TE D

modulus (Gʺ) of PLLA/PBS-25%/(PLLA-b-PGMA)3 blends: (a)100/0/0; (b)60/40/0; (c)60/40/1;

The rheological behavior of the PLLA/PBS-25% (60/40) blend was measured at the

EP

processing temperature, and the complex viscosity (η*), storage modulus (Gʹ), and

AC C

loss modulus (Gʺ) as a function of frequency for a typical response were shown in Fig. 9A(b), Fig. 9B(b), Fig. 9C(b). The η* of linear PLLA decreased with increasing frequency due to the pronounced shear shinning at high frequencies (Fig. 9A(a)), indicating a non-Newtonian behavior and pseudoplastic characteristics over the entire testing frequency range. The incorporation of PBS-25% into PLLA matrix had no positive influence on the PLLA rheological behavior, and the complex viscosity (Fig. 9A(b)) was decreased. Gʹ(Fig. 9B(b)) and Gʺ(Fig. 9C(b)) of PLLA/PBS-25% blend were also decreased compared to linear PLLA, which can be ascribed to the lower

ACCEPTED MANUSCRIPT viscosity of the PBS-25% and lack of branched structure. Many attempts had been made to improve the melt strength of PLLA, and it was proved that modification of PLLA to get branching structures with high molecular weight was an efficient approach to achieve this purpose. The use of an

RI PT

epoxy-functionalized chain extender was able to form the branched structures. Based on above analysis, (PLLA-b-PGMA)3 was selected as the chain extender for the PLLA/PBS-25% (60/40) system.

shear

flow.

In

this

study,

the

SC

The melt strength of polymers was often reported as the elastic behavior under rheological

behavior

of

the

M AN U

PLLA(60)/PBS-25%(40)/(PLLA-b-PGMA)3 blend was measured at a temperature of 180 oC to investigate the influence of (PLLA-b-PGMA)3 on the melt strength of the blends. The complex viscosity (η*), storage modulus (Gʹ), and loss modulus (Gʺ) as a function of frequency for a typical response were shown in Fig. 9. It was clearly

TE D

observed that the addition of (PLLA-b-PGMA)3 had a great influence on the rheological behavior of PLLA blends. The incorporation of a small quantity of (PLLA-b-PGMA)3 into PLLA/PBS-25% increased the η* compared to neat PLLA

EP

and PLLA/PBS-25% blend, especially at low frequency, indicating the presence of

AC C

chain extension and/or a branching reaction phenomenon. The longer chains created more entanglements, thus giving rise to a higher viscosity. Also, the η* of the PLLA blends was increased quickly at first and then slowly with the increasing of (PLLA-b-PGMA)3

content,

suggesting

that

the

addition

of

excessive

(PLLA-b-PGMA)3 was not necessary for long-chain branching reaction. For the Gʹ, there was a distinct formation of a rubbery plateau of highest length in the case of the PLLA(60)/PBS-25%(40)/(PLLA-b-PGMA)3 blends unlike that of neat PLLA, and the PLLA/PBS-25% (60/40) blend in low frequency region. The formation

ACCEPTED MANUSCRIPT of a rubbery plateau was determined by the presence of branching and entanglement in a polymer. From the low frequency region of all curves, it was found that the PLLA(60)/PBS-25%(40)/(PLLA-b-PGMA)3 blends had highest value of plateau modulus. Such an increase in plateau modulus was caused by the presence of

RI PT

branching or entangled networks in a polymer melt. On the other hand, the in-situ reactive blending of PBS-25% with (PLLA-b-PGMA)3 in the PLLA melt caused the formation of branching structure, which had the tendency to form entanglements with

SC

the PLLA chain. The terminal slope of the curve approached zero for the PLLA(60)/PBS-25%(40)/(PLLA-b-PGMA)3 blends, reflecting its pseudo-solid-like

M AN U

behavior. The relaxation and motion of the PLLA chain, which was a major fraction in the blend, would have a prominent effect on rheological properties in low shear regime.

The

longer

relaxation

of

PLLA(60)/PBS-25%(40)/(PLLA-b-PGMA)3

the

blends

PLLA caused

chain the

in

the

frequency

TE D

independence of storage modulus in lower shear regime. This behavior also strongly indicated toward formation of a physical network of the PLLA chains with PBS-25%/(PLLA-b-PGMA)3 blends during their in-situ reactive blending in the

structure in

AC C

branched

EP

PLLA matrix. The increase of Gʺ also supported the truly existence of the long chain the PLLA matrix.

In

addition,

the influence

of

(PLLA-b-PGMA)3 contents on Gʹ and Gʺ of PLLA/PBS-25% (60/40) was in agreement

with

the

influence

PLLA/PBS-25% (60/40) blend.

of

(PLLA-b-PGMA)3

contents

on

η*

of

RI PT

ACCEPTED MANUSCRIPT

Fig. 10. Elongational viscosity of the PLLA/PBS-25%/(PLLA-b-PGMA)3 blends at strain rate of 3

SC

s-1: (a) neat PLLA; (b) 60/40/0; (c) 60/40/1; (d) 60/40/2; (e) 60/40/5.

M AN U

As compared to the shear rheology, elongational viscosity was more sensitive to the branched molecular structure. The elongational viscosity was recorded as a function of specimen strain, a sudden increase in viscosity at a certain strain was called strain-hardening

behavior,

which

had

been

reported

for

many

branched

polymers.[35,36] Such a strain hardening was important for polymer processing

TE D

where high melt strength was required, such as film blowing.[37] The elongational viscosities of PLLA and PLLA blends as a function of time at elongational rates of 0.3

EP

s−1 were shown in Fig. 10. It was clear that the elongational viscosities of neat PLLA, and PLLA/PBS-25% (60/40) blend increased with time at the start of stretching, and

AC C

then decreased, which exhibited obvious strain softening phenomena. It meant that neat PLLA was not suitable for blown-filming because of its low melt strength. In contrast, the elongational viscosity of PLLA(60)/PBS-25%(40)/(PLLA-b-PGMA)3 blends were markedly improved, and exhibited strain-hardening behavior. The linear PLLA and PLLA/PBS-25% (60/40) blend, where the polymer chains were free from any branch points, maintained a linear response over time, and the chains were not prevented from slipping over each other. In chain branched polymers, the chain slipping was perturbed by the presence of side branches entangled with the main

ACCEPTED MANUSCRIPT chains, causing strain hardening. In this case, the stretching force or stress reached a pseudosteady state before increasing again due to the resistance caused by the branches entanglement with the main chains of the PLLA. In this work, the branched structure was formed by the reaction between PBS-25% and (PLLA-b-PGMA)3, and

RI PT

the macromolecular chain disentanglement rate was too low in relation with the

M AN U

SC

deformation rate.

of 20 mm/min.

TE D

Fig. 11. Stress-strain curves of PLLA/PBS-25%/(PLLA-b-PGMA)3 obtained at a cross-head speed

Table 3. Mechanical properties of PLLA/PBS-25%/(PLLA-b-PGMA)3 for 5 times. Elongation at

Tensile strength

Tensile modulus

Toughness

/(PLLA-b-PGMA)3

break (%)

at yield (MPa)

（MPa）

(MPa)

6.7±3.2

64.6±2.5

1980±150

266±120

117±10.0

49.7±2.0

1360±100

2587±220

60/40/2

195±10.0

50.6±1.5

1240±80

5002±245

60/40/5

234±12.0

47.8±1.5

1050±90

6002±280

100/0/5

AC C

60/40/1

EP

PLLA/PBS-25%

Stress-strain curves of PLLA(60)/PBS-25%(40)/(PLLA-b-PGMA)3 blends with

different contents of (PLLA-b-PGMA)3 were shown in Fig. 11, and the details of the tensile properties were given in Table 3. The toughness increased from 192 MPa for the neat PLLA to 5002 MPa for PLLA(60)/PBS-25%(40)/(PLLA-b-PGMA)3(2), and further increased with the (PLLA-b-PGMA)3. The addition of (PLLA-b-PGMA)3

ACCEPTED MANUSCRIPT improved interfacial adhesion of the PLLA blends, which contributed to the improvement of toughness. The elongation at break increased to 117% even a small fraction (1%) of (PLLA-b-PGMA)3 was added, which was much higher than those of PLLA/PBS-25% (60/40) blend (81%). When the (PLLA-b-PGMA)3 contents

RI PT

increased from 1% to 2% and 5%, the elongation at break increased from 117% to 195% and 234%, respectively. The PLLA(60)/PBS-25%(40)/(PLLA-b-PGMA)3 blends showed initial strain softening after yielding and then underwent considerable

SC

cold drawing. The stress-strain curve after the yield point showed a combination of strain softening and cold drawing. In this region, there was competition between

M AN U

PLLA chain orientation and crack formation. Hence, there was a drop in stress with increasing strain. After 20% of strain, only cold drawing dominated at a constant stress. This suggested that a large energy dissipation occurred in the presence of branched structure in the PLLA(60)/PBS-25%(40)/(PLLA-b-PGMA)3 blend, and it

TE D

was effective in improving the toughness of PLLA by the reactive blending. The tensile strength was increased firstly to 50.6 MPa and then decreased to 47.8 MPa. Therefore, a given quantity of (PLLA-b-PGMA)3 molecules was needed to participate

EP

in the reactions. Beyond the required number of (PLLA-b-PGMA)3 molecules, the

AC C

excess (PLLA-b-PGMA)3 would only plasticize the polymer blend, thus lowering the strength of the blends. It was worth noting that the tensile strength of PLLA/PBS-25%/(PLLA-b-PGMA)3 in our study was still greater than that of general petroleum based package polymer, such as PE (40 MPa) and PP (32 MPa). The result was very significant in obtaining a biobased material with remarkable toughness balance. A control sample of the PLLA/(PLLA-b-PGMA)3 (100/5) blend was also fabricated at the same processing condition as that of the PLLA/PBS-25%/(PLLA-b-PGMA)3

ACCEPTED MANUSCRIPT (60/40/5) blend. The purpose of this experiment was to investigate the effect of single (PLLA-b-PGMA)3 on the properties of PLLA. There was no obviously improvement of elongation at break for PLLA/(PLLA-b-PGMA)3 (100/5) (Table 3).

M AN U

SC

RI PT

3.4 Thermal properties

Fig. 12. DSC curves of PLLA/PBS-25%/(PLLA-b-PGMA)3 of first heating. The right curves gave enlarged DSC curves of the blends.

AC C

EP

TE D

Table 4. Thermal properties and the crystallinity of PLLA/PBS-25%/ (PLLA-b-PGMA)3 blends. Sample Tg (oC) Tc (oC) Tm (oC) (∆Hmc) (J g-1) (∆Hm) (J g-1) Xc (%)a 100/0/0 61.80 126.8 177.80 35.98 39.12 3.33 70/30/0 60.83 177.80 40.91 40.91 70/30/1 60.69 175.21 40.36 37.29 70/30/2 57.36 173.19 39.65 37.24 70/30/5 55.58 172.06 37.33 39.65 a c * -1 Crystallinity: Xc = (∆Hm-∆Hm /∆Hm ) * 100%; ∆Hm*=93 J g ,the melting enthalpy of 100% crystalline PLA, ∆Hmc is cold crystallization enthalpy. It was important to verify how the reactive blending could affect the thermal behavior of PLLA. Fig. 12 presents DSC scans obtained during the first heating of PLLA samples extruded with different contents of (PLLA-b-PGMA)3. As shown in Fig. 13, neat PLLA showed an obvious Tg at about 61.8

o

C and a melting

temperature of 177.8 oC. Also, a broad exothermic peak corresponding to its cold crystallization was presented at 126.8 oC. Tg of PLLA was not obvious after the addition of the PBS-30% and the (PLLA-b-PGMA)3. This was attributed to the

ACCEPTED MANUSCRIPT increase of the crystallinity of PLLA. As shown in Table 4, the crystallinity of neat PLLA was only 3.3%, in contrast, the crystallinity of the PLLA blends was more than 35%. The increase of the PLLA crystalline region resulted in the limitation of the PLLA amorphous region, and the Tg is difficult to be displayed. And the Tg was also

RI PT

decreased, indicating the improvement of the compatibility of the blends.

The Tm showed little change after blending with PBS-25%, however, the cold crystallization peak of PLLA diminished after the addition of PBS-25%, and the

SC

PLLA/PBS-25% (60/40) blend showed melting peak only, indicating that the addition of PBS-25% enhanced the crystallization rate of the PLLA component. Compared the

PLLA/PBS-25%

(60/40)

blend,

the

Tm

M AN U

with

of

PLLA

in

the

PLLA(60)/PBS-25%(40)/(PLLA-b-PGMA)3 was found to shift toward a lower temperature region, which may be due to the less perfection of the PLLA crystals and a reduction in the lamellar thickness resulted from the branching or even

TE D

cross-linking structures in the presence of (PLLA-b-PGMA)3. The restricted chain mobility in the branching structures or cross-linking structures could lead to difficulties in the thickening of the lamella; hence, thinner and imperfect crystals with

EP

lower melting point were formed. The melting enthalpy (∆Hm) reflected the amount of crystallinity developed in each sample. When the (PLLA-b-PGMA)3 contents

AC C

increased from 0% to 2%, the melting enthalpy decreased from 40.91 J g-1 to 39.65 J g-1. The decrease in crystallinity would imply a decrease in the brittleness in the blend, which was in agreement with the tensile experiment. 3.5 Morphology Since mechanical properties of multiphase polymer blends depended largely on the resulting morphologies during melt-blending, SEM was employed to identify the phase structure of the blends. SEM images of fractured surfaces of neat PLLA,

ACCEPTED MANUSCRIPT PLLA/PBS-25%, and PLLA/PBS-25%/(PLLA-b-PGMA)3 blends were shown in Fig. 13. The surface of neat PLLA was extremely flat, indicating the apparent brittle failure of PLLA (Fig. 13A). As compared to the neat PLLA, PLLA/PBS-25% (60/40) blend displayed some coarse surfaces (Fig. 13B). This phenomenon indicated that

RI PT

there was strong interfacial adhesion between PLLA and PBS-25%, which ascribed to the PLLA units in the PBS-25%, and the PLLA units played the role to improve the compatibility of PBS-25% and neat PLLA. After addition of (PLLA-b-PGMA)3 (Fig.

SC

13C), the in-situ reactive blending of PBS-25% and (PLLA-b-PGMA)3 in the PLLA matrix dramatically changed its surface characteristic. The fractured surface of the

M AN U

PLLA/PBS-25%/(PLLA-b-PGMA)3 (60/40/1) exhibited considerable ductile tearing, and surface roughness. The good compatibility and the entanglement between the PLLA and PBS-25% phase resulted in significant toughness enhancement of PLLA. The fracture surface became more rough and some stubbly fibers were observed with

TE D

the contents of (PLLA-b-PGMA)3 increased to 2% (Fig. 13D). This indicates that an ideal interfacial adhesion between PLLA and PBS-25% was obtained via in situ reactive blending, thus the shear yielding of the PLLA matrix occurred during the

EP

interfacial debonding. The toughness of the blend was also significantly improved.

AC C

With the further increase of (PLLA-b-PGMA)3 (5%), the fracture surface produced significantly plastic deformation (Fig. 13E), which was induced by the finer interfacial adhesion and stronger entanglement between the two polymer domains. The plastic deformation induced energy dissipation and therefore led to the improvement in tensile toughness of the PLLA/PBS-25%/(PLLA-b-PGMA)3 blends. The SEM observation was well consistent with the results of mechanical properties.

(B)

(C)

(D)

AC C

EP

TE D

(E)

M AN U

SC

(A)

RI PT

ACCEPTED MANUSCRIPT

Fig. 13. SEM micrographs of the cryogenically fractured surface of PLLA/PBS-25%/ (PLLA-b-PGMA)3 blends: (A) 100/0/0, (B)60/40/0, (C) 60/40/1; (D) 60/40/2; (E) 60/40/5.

During tensile testing, there was stress-whitening in the subfracture region in

PBS-25%, PLLA/PBS-25% and PLLA/PBS-25%/(PLLA-b-PGMA)3 blends, as shown in Fig. 14. There was the possibility of the creation of numerous tiny crazes in the PLLA matrix. The stress-whitening occurred because of the scattering of light by these tiny crazes created by block copolymer or branched polymer. Such intense

ACCEPTED MANUSCRIPT stress-whitening was absent in the neat PLLA sample (Fig 14a). The phenomenon of multiple crazing was likely the possible energy dissipation mechanism behind the toughening of PLLA. The good compatibility of PLLA with PBS-25% results in the formation of numerous crazes in the PLLA matrix. Also, the in-situ reactive blending

RI PT

of PBS-25% and (PLLA-b-PGMA)3 in the PLLA matrix resulted in the formation of a branched copolymer at the PLLA/PBS-25% interface, which further improved compatibility between the two phases. The creation of more stable crazes in the PLLA

SC

matrix by branched structure was helpful for substantial drawing and orientation of

TE D

M AN U

PLLA chains under tensile loading.

Fig. 14. Representative deformed samples PLLA/PBS-25%/(PLLA-b-PGMA)3

blends are

60/40/5.

EP

shown in the photograph: (a) 100/0/0; (b) 0/100/0; (c) 60/40/0; (d) 60/40/1; (e) 60/40/2; (f)

AC C

To further explain the ductile tearing and plastic deformation of PLLA/PBS-25%/ (PLLA-b-PGMA)3 blends, the SEM of the tensile fracture surfaces was also investigated, as shown in Fig. 15. The neat PLLA showed the brittle tensile surface. The addition of PBS-25% and (PLLA-b-PGMA)3 significantly changed its fracture surface

after

tensile.

Ductile

PLLA/PBS-25%/(PLLA-b-PGMA)3

tearing

and

rough

surface

of

blends indicated that the reactive blending and

block copolymerization significantly improved the interfacial interaction between the PLLA and PBS-25%, and remarkably improved the toughness of PLLA.

(B)

(C)

(D)

M AN U

SC

(A)

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

(E)

Fig. 15. SEM micrographs of the tensile fracture surface of PLLA/PBS-25%/(PLLA-b-PGMA)3 blends: (A) 100/0/0, (B)60/40/0, (C) 60/40/1; (D) 60/40/2; (E) 60/40/5.

4. Conclusions High melt strength and high toughness PLLA/PLLA-b-PBS-b-PLLA/(PLLA-bPGMA)3

blends

were

prepared

by

in

situ

reactive

blending.

The

PLLA-b-PBS-b-PLLA block copolymer contributed to the enhancement of the

ACCEPTED MANUSCRIPT toughness of the PLLA, and had good compatitility with PLLA. The introduction of PLLA-based epoxy-functional (PLLA-b-PGMA)3 block copolymer improved the melt strength of PLLA greatly due to the alternation of the topology of the blend components from linear structures to long chain branched structures, and

RI PT

PLLA/PLLA-b-PBS-b-PLLA/(PLLA-b-PGMA)3 blends showed obvious strain hardening behavior. The generation of new blends improved the toughness and elongation at break greatly and remarkably maintained high strength and modulus

demonstrated

by

the

lowered

melt

point

SC

values. DSC studies also provided evidence of low crystallization capacity as and

melting

enthalpy

of

M AN U

PLLA/PLLA-b-PBS-b-PLLA/(PLLA-b-PGMA)3 blends. The interfacial adhesion between PLLA and PLLA-b-PBS-b-PLLA was improved greatly after addition of the (PLLA-b-PGMA)3, and further confirmed the results of the mechanical properties. The incorporation of PLLA-b-PBS-b-PLLA and (PLLA-b-PGMA)3 into the PLLA

TE D

imparted enhanced toughness, strain harderning and melt strength to the system. To the best of our knowledge, the combination of copolymerization and in situ reactive blending to modify PLLA had not been reported elsewhere. Thus, such an approach

EP

developed a system with the enhanced performance and processability of PLLA

AC C

blends, and had great meaning for the expanding application of PLLA.

Acknowledgements

This work was supported by grants from National Natural Science Foundation of

China (Key Project 51303176), National High-tech R&D Program of China (863 Program Project 2015AA034004).

ACCEPTED MANUSCRIPT

References [1] K. Hamad, M. Kaseem, Y.G. Ko, F. Deri, Biodegradable Polymer Blends and Composites: An overview, Polym. Sci. Ser. A 56(6) (2014) 812-829. [2] N. Peelman, P. Ragaert, B.D. Meulenaer, D. Adons, R. Peeters, L. Cardon, F.V.

Food. Sci. Tech. 32 (2013) 128-141.

RI PT

Impe, F. Devlieghere, Application of bioplastics for food packaging, Trends.

[3] H.Z. Liu, J.W. Zhang, Research progress in toughening modification of

SC

poly(lactic acid), J. Polym. Sci. Part. B. Polym. Phys. 49 (2011) 1051-1083.

[4] X. Pang, X.L. Zhuang, Z.H. Tang, X.S. Chen, Polylactic acid (PLA): Research,

M AN U

development and industrialization, Biotechnol. J. 5 (2010) 1125-1136. [5] D. Garlotta, A literature review of poly(Lactic Acid), J. Polym. Environ. 9(2) (2001) 63-84.

[6] R. Auras, B. Harte, S. Selke, An overview of polylactides as packaging materials,

TE D

Macromol. Biosci. 4 (2004) 835-864.

[7] K.S. Anderson, K.M. Schreck, M.A. Hillmyer, Toughening polylactide, Polym. Rev. 48 (2008) 85-108.

EP

[8] K.M. Nampoothiri, N.R. Nair, R.P. John, An overview of the recent

AC C

developments in polylactide (PLA) research, Bioresour. Technol. 101 (2010) 8493-8501.

[9] S. Nouri, P.G. Lafleur, C. Dubois, Enhanced film blowing of polylactide by incorporating branched chains and stereocomplex crystals, Int. Polym. Proc. 30(4) (2015) 500-510. [10]M.T.

Martello,

M.A.

Hillmyer,

Polylactide-poly(6-methyl-epsilon-

caprolactone)-polylactide thermoplastic elastomers, Macromolecules 44 (2011) 8537-8545.

ACCEPTED MANUSCRIPT [11] H. Xiu,

H.W. Bai, J. Jiang, F. Chen, H. Deng, K. Wang, Q. Zhang, Q. Fu,

Improving impact toughness of polylactide/poly(ether)urethane blends via designing the phase morphology assisted by hydrophilic silica nanoparticles, Polymer 55(6) (2014) 1593-1600.

RI PT

[12] Q. Peng, K. Mahmood, Y. Wu, L. Wang, Y. Liang, J. Shen, Z. Liu, A facile route to realize the copolymerization of L-lactic acid and epsilon-caprolactone: sulfonic acid-functionalized Bronsted acidic ionic liquids as both solvents and catalysts,

SC

Green. Chem. 16(4) (2014) 2234-2241.

[13] M. Kowalczyk, E. Piorkowska, S. Dutkiewicz, P. Sowinski, Toughening of

Polym. J. 59 (2014) 59-68.

M AN U

polylactide by blending with a novel random aliphatic-aromatic copolyester, Eur.

[14] M.K. Kang, Y. Jung, S.H. Kim, Biodegradable stereocomplex polylactide having flexible ε-caprolactone unit, Macromol. Res. 21(9) (2013) 1036-1041.

TE D

[15] X.Q. Zeng, B.S. Wu, L.B. Wu, J.J. Hu, Z.Y. Bu, B.G. Li, Poly(L-lactic acid)block-poly(butylene succinate-co-butylene adipate) multiblock copolymers: from synthesis to thermo-mechanical properties, Ind. Eng. Chem. Res. 53 (2014)

EP

3550-3558.

AC C

[16] D. Wu, Y. Zhang, M. Zhang, Phase Behavior and its viscoelastic response of polylactide/poly(ε-caprolactone) blend, Eur. Polym. J. 44 (2008) 2171-2183.

[17] T. Qiang, D.M. Yu, A.J. Zhang, H.H. Gao, Z. Li, Z.C. Liu, W.X. Chen, Z. Han, Life cycle assessment on polylactide-based wood plastic composites toughened with polyhydroxyalkanoates, J. Clean. Prod. 66 (2014) 139-145. [18] M. Shibata, Y. Inoue, M. Miyoshi, Mechanical properties, morphology, and crystallization behavior of blends of poly(l-lactide) with poly(butylenes succinate-co-l-lactate) and poly(butylene succinate), Polymer 47 (2006) 3557-

ACCEPTED MANUSCRIPT 3564. [19]L. Jiang, B. Liu, J. Zhang, Properties of poly(lactic acid)/poly(butylene adipate-co-terephthalate)/nanoparticle ternary composites, Ind. Eng. Chem. Res. 48 (2009) 7594-7602.

properties

of

immiscible

blends

of

RI PT

[20]V. Ojijo, S.S. Ray, R. Sadiku, Role of specific interfacial area in controlling biodegradable

polylactide and

6689-6700.

SC

poly[(butylene succinate)-co-adipate], ACS Appl. Mater. Inter. 4(12) (2012)

[21]M. Shibata, N. Teramoto, Y. Inoue, Mechanical properties, morphologies, and behavior

of

plasticized

poly(L-lactide)/poly(butylene

M AN U

crystallization

succinate-co-L-lactate) blends, Polymer 48(9) (2007) 2768-2777. [22]M. Nerkar, J.A. Ramsay, B.A. Ramsay, A.A. Vasileiou, M. Kontopoulou, Improvements in the melt and solid-state properties of poly(lactic acid),

TE D

poly-3-hydroxyoctanoate and their blends through reactive modification, Polymer 64 (2015) 51-61.

[23] X. Li, H.L. Kang, J.X. Shen, L.Q. Zhang, T. Nishi, K. Ito, C.M. Zhao, P. Coates,

EP

Highly toughened polylactide with novel sliding graft copolymer by in situ

AC C

reactive compatibilization, crosslinking and chain extension, Polymer 55(16) (2014) 4313-4323.

[24] P. Ma, X. Cai, Y. Zhang, S. Wang, W. Dong, M. Chen, P.J. Lemstra, In-situ compatibilization of poly(lactic acid) and poly(butylene adipate-co-terephthalate) blends by using dicumyl peroxide as a free-radical initiator, Polym. Degrad. Stab. 102 (2014) 145-151.

ACCEPTED MANUSCRIPT [25] T. Semba, K. Kitagawa, U.S. Ishiaku, H. Hamada, The effect of crosslinking on the mechanical properties of polylactic acid/polycaprolactone blends, J. Appl. Polym. Sci. 101(3) (2006) 1816-1825. [26] D.Y. Ji, Z.Y. Liu, X.R. Lan, F. Wu, B.H. Xie, M.B. Yang, Morphology, rheology, behavior,

and

mechanical

properties

of

poly(lactic

RI PT

crystallization

acid)/poly(butylene succinate)/dicumyl peroxide reactive blends, J. Appl. Polym. Sci. 131(3) (2014) 39580,

SC

[27] G.C. Liu, Y.S. He, J.B. Zeng, Y.Z. Wang, In situ formed crosslinked polyurethane

M AN U

toughened polylactide, Polym. Chem. 5(7) (2014) 2530-2539.

[28] C. Kaynak, Y. Meyva, Use of maleic anhydride compatibilization to improve toughness and other properties of polylactide blended with thermoplastic elastomers, Polym. Adv. Technol. 25(12) (2014) 1622-1632. [29] Y.S.

Wang,

Z.Y.

Wei,

Leng,

epoxidized

TE D

toughened polylactide with

X.F.

K.H.

Shen,

polybutadiene

Y.

by

Li,

Highly

in-situ

reactive

compatibilization, Polymer 92 (2016) 74-83. [30] V. Ojijo, S.S. Ray, Super toughened biodegradable polylactide blends with

EP

non-linear copolymer interfacial architecture obtained via facile in-situ

AC C

reactive compatibilization, Polymer 80 (2015) 1-17. [31] W.F. Dong, B.S. Zou, Y.Y. Yan, P.M. Ma, M.Q. Chen, Effect of chain-extenders on

the

properties

and

hydrolytic

degradation

behavior

of

the

poly(lactide)/poly(butylene adipate-co-terephthalate) blends, Int. J. Mol. Sci. 14 (2013) 20189-20203.

[32] R. Al-Itry, K. Lamnawar, A. Maazouz, Improvement of thermal stability, rheological and mechanical properties of PLA, PBAT and their blends by reactive extrusion with functionalized epoxy, Polym. Degrad. Stab. 97 (2012) 1898-1914.

ACCEPTED MANUSCRIPT [33] X.C. Bian, B. Zhang, Z. Q. Sun, S. Xiang, G. Li, X.S. Chen, Synthesis of multi-arm poly(L-lactide) and its modification on linear polylactide, Polym. Bull. (2016) doi: 10.1007/s00289-016-1713-4. [34] X.C. Bian, B. Zhang, B. Sun, Z. Q. Sun, S. Xiang, G. Li, X.S. Chen, Preparation

RI PT

of high toughness and high transparency polylactide blends resin based on multiarmed polycaprolactone-block-poly(L-lactide), Polym. Eng. Sci. 56(10) (2016) 1125-1137.

SC

[35] G.X. Liu, H.W. Ma, H. Lee, H.D. Xu, S.W. Cheng, H. Sun, T.Y. Chang, R.P. Quirk, S.Q. Wang, Long-chain branched polymers to prolong homogeneous

M AN U

stretching and to resist melt breakup, Polymer 54 (2013) 6608-6616. [36] L.Y. Wang, X.B. Jing, H.B. Cheng, X.L. Hu, L.X. Yang, Y.B. Huang, Rheology and crystallization of long-chain branched poly(L-lactide)s with controlled branch length, Ind. Eng. Chem. Res. 51 (2012) 10731-10741.

TE D

[37] R.M. Rasala, A.V. Janorkarc, D.E. Hirta, Poly(lactic acid) modifications, Prog.

AC C

EP

Polym. Sci. 35 (2010) 338-356.

ACCEPTED MANUSCRIPT

Figure captions Scheme 1. Synthesis of PLLA-b-PBS-b-PLLA. Scheme 2. Synthesis of 3-arm block copolymer (PLLA-b-PGMA)3. Fig. 1. 1H NMR spectra of PBS (A), and PLLA-b-PBS-b-PLLA (PBS-25%) (B) were

RI PT

recorded at room temperature in CDCl3. Fig. 2. IR spectrum of PBS.

PLLA-b-PBS-b-PLLA (Mn=62000，PDI=1.5).

SC

Fig. 3. GPC traces of PBS and PLLA-b-PBS-b-PLLA. PBS (Mn=17000, PDI=1.7),

M AN U

Fig. 4. 1H NMR spectra of 3-arm PLLA (A) macroinitiator (B) and block copolymer (PLLA-b-PGMA)3 (C) was recorded at room temperature in CDCl3. Fig. 5.

13

C NMR spectrum of 3-arm PLLA was recorded at room temperature in

CDCl3.

Fig. 6. GPC traces of 3-arm PLLA (a), macroinitiator (b), and block copolymer

TE D

(PLLA-b-PGMA)3 (c): (a) Mn=9200 ， PDI=1.51; (b) Mn=9500 ， PDI=1.42; (c) Mn=15000，PDI=1.38.

EP

Fig. 7. Stress-strain curves obtained at a cross-head speed of 20 mm/min for the PLLA-b-PBS-b-PLLA with different PBS contents: (a) neat PLLA; (b) 5%; (c) 15%;

AC C

(d) 25%; (e) 35%. The right curves gave details of stress-strain of the blends in the neighborhood of yield points. Fig. 8. Stress-strain curves obtained at a cross-head speed of 20 mm/min for the PLLA/PBS-25% blends with different PBS-25% contents. Fig. 9. Frequency dependence of A complex viscosity (η*), B storage modulus (Gʹ ) and C loss modulus (Gʺ) of PLLA/PBS-25%/(PLLA-b-PGMA)3 blends: (a)100/0/0; (b)60/40/0; (c)60/40/1; (d) 60/40/2;(e) 60/40/5. Fig. 10. Elongational viscosity of the PLLA/PBS-25%/(PLLA-b-PGMA)3 blends at

ACCEPTED MANUSCRIPT strain rate of 3 s-1: (a) neat PLLA; (b) 60/40/0; (c) 60/40/1; (d) 60/40/2; (e) 60/40/5. Fig. 11. Stress-strain curves of PLLA/PBS-25%/(PLLA-b-PGMA)3 obtained at a cross-head speed of 20 mm/min. Fig. 12. DSC curves of PLLA/PBS-25%/(PLLA-b-PGMA)3 of first heating. The right

RI PT

curves gave enlarged DSC curves of the blends.

Fig. 13. SEM micrographs of the cryogenically fractured surface of PLLA/PBS-25%/ (PLLA-b-PGMA)3 blends (A) 100/0/0, (B)60/40/0, (C) 60/40/1; (D) 60/40/2; (E)

SC

60/40/5.

Fig. 14. Representative deformed samples PLLA/PBS-25%/(PLLA-b-PGMA)3

(e) 60/40/2; (f) 60/40/5. Fig.

15.

SEM

micrographs

M AN U

blends are shown in the photograph: (a) 100/0/0; (b) 0/100/0; (c) 60/40/0; (d) 60/40/1;

of

the

PLLA/PBS-25%/(PLLA-b-PGMA)3 blends:

AC C

EP

TE D

(D) 60/40/2; (E) 60/40/5.

tensile

fracture

surface

of

(A) 100/0/0, (B)60/40/0, (C) 60/40/1;