Synergistic toughening of polypropylene random copolymer at low temperature: β-Modification and annealing

Materials Science and Engineering A 528 (2011) 7052–7059

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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea

Synergistic toughening of polypropylene random copolymer at low temperature: ␤-Modification and annealing Feng Luo a , Jinwen Wang a , Hongwei Bai a , Ke Wang a,∗ , Hua Deng a , Qin Zhang a , Feng Chen a , Qiang Fu a,∗∗ , Bing Na b a b

Department of Polymer Science and Materials, State Key Laboratory of Polymer Materials Engineering, Sichuan University, Chengdu 610065, People’s Republic of China Department of Materials Science and Engineering, East China Institute of Technology, Fuzhou 344000, People’s Republic of China

a r t i c l e

i n f o

Article history: Received 2 April 2011 Accepted 18 May 2011 Available online 26 May 2011 Keywords: Polypropylene random copolymer Toughening Annealing ␤-Nucleating

a b s t r a c t The synergistic effect of ␤-modification and annealing on the impact toughness of a commercially available polypropylene random copolymer (PPR) was investigated. Interestingly, the impact toughness of ␤-nucleated PPR after annealing at low temperature (0 ◦ C) was almost five times as high as that of the virgin PPR without annealing. The crystalline structure, supermolecular structure, phase morphology, the relaxation of chain segments and fracture behavior of matrix were investigated to explore the toughening mechanism related to the ␤-modification and annealing. It was found that annealing improved the mobility of chain segments in amorphous phase as well as the strength of ligament of PPR matrix with profuse ␤ crystals having the intrinsic low plastic deformation resistance, responsible for the superior toughness achieved. This work provides a possible method to toughen semi-crystalline polymers at low temperature by combination of ␤-modification and suitable annealing. © 2011 Elsevier B.V. All rights reserved.

1. Introduction Polypropylene (PP) is one of the most widely used thermoplastic; but its application in some fields is limited especially at low temperature, due to its low impact resistance. The fracture resistance of this semi-crystalline polymer shows dramatic dependence on the ability of shear yielding and crazing [1,2]. The way of promoting extensive shear yielding and/or multiple crazing in the polymer matrix is much beneficial for toughening. As demonstrated in the past studies, various elastomers are thought to be the most efficient toughening agents for PP; but the improvement of fracture toughness is at the cost of deterioration of stiffness and strength [3–5]. Instead of adding elastomers, copolymerization of propylene with ethylene or other olefins is another useful and effective method to produce high-performance polypropylene copolymers, PPR [6–11]. In this copolymer, the homopolymer sequences are semi-crystalline and form a compound crystalline-amorphous biphase; while the propylene-ethylene random segments with high ethylene content tend to coalesce together to form a rubbery phase. It is indeed a multiphase polyolefin system with an excellent rigidity-toughness balance: the crystalline phase of homopoly-

∗ Corresponding author. ∗∗ Corresponding author. Tel.: +86 28 85405402. E-mail addresses: [email protected] (K. Wang), [email protected] (Q. Fu). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.05.030

mer sequences guarantees the moderate strength and modulus, and the well-dispersed rubbery domains offer superior toughness. Therefore, polypropylene random copolymers (PPR) constitute an important class of plastic resins that are widely used as matrix components in pipe, automobile parts, furniture, and other industrial uses in the past two decades, based on the excellent mechanical properties and relatively low production cost [12,13,5]. Even so, it is also found that the toughness of PPR at low temperature (below Tg ) is still too low and thus limits its further applications. On the other hand, polypropylene is a polymorphic material with at least four crystalline forms, namely, the monoclinic ␣-form, trigonal ␤-form, orthorhombic ␥-form, and mesomorphic smectic form [14,15]. The differences in the supermolecular structures and the different crystalline states of PP exhibit different mechanical features. For example, the ␣-PP shows excellent modulus and tensile strength but inferior fracture toughness because the presence of interlocking effect of the radial lamellae by the tangential crystallites makes the plastic deformation very difficult [16,17]. On the contrary, ␤-PP without cross-hatching allows the initiation and propagation of plastic deformation more easily and then enhances the energy dissipation [18–21]. Especially, the enhanced toughness of ␤-PP can be attributed to a stress-induced transformation from less dense (␤-phase) to more dense (␣-phase) crystalline structure at the root of a growing crack [22,23]. An additional factor that should be responsible for the reducing resistance to plastic flow initiation is the crystallographic symmetry of the hexagonal ␤

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40

30

20

10

-A PR β− P

PR β− P

-A

0

R

Polypropylene random copolymer (PPR) used in this study was a commercial grade R200P supplied by Hyosung (Korea), with Mw = 72.2 × 104 g/mol, Mw /Mn = 4.5 and density of 0.91 g/cm3 . The mass percentage of ethylene component was as low as 3.8 wt.%. A small amount of antioxidant (Irganox 1010) was added into PPR to prevent the thermal decomposition during melt processing. The rare earth ␤-nucleating agent, marked as WBG, was kindly supplied by Winner Functional Materials Co. (Foshan, Guangdong, China).

The notched Izod impact strength of the specimens was measured with a VJ-40 Izod machine according to ASTM D256-04. For low temperature impact test, the specimens were keep at 0 ◦ C for 12 h, and then immediately (less than 5 s) subjected to impact. The average values were obtained over five specimens. Standard tensile tests were performed on a dumbbell-shaped specimen using an SANS Universal tensile testing machine according to the ASTM D638-03 standard. The natural draw ratio n was determined from the separation of ink marks at regular 1 mm intervals preprinted on the specimens. A crosshead speed of 50 mm/min was used. In order to examine the crack-initiation stage, an arrested crack in the specimen was produced with a notch Izod part-impact tester at 0 ◦ C. In detail, the pendulum was raised at an angle of 60◦ from the vertical and then released to hit the specimen. With appropriately chosen impact energy about 0.2 J, the specimen was not broken into two halves, and the propagating crack stopped in the interior of the specimen. For polarized light microscopy (PLM), the slices with a thickness of around 30 ␮m were cut from injection-molded specimens directly using a Leica RM2245 microtome, and then inspected on a Leica DMIP PLM equipped with crossed polarizer. Similarly, the crack-initiation patterns of the slices with 30 ␮m thick cutting from

PP

2.1. Materials and sample preparation

2.2. Characterizations and measurements

PP R

2. Experimental

The PPR and WBG powders were first melt compounded to make a master-batch containing 5 wt.% WBG in a co-rotating twin screw extruder (TSSJ-25 co-rotating twin-screw extruder, China); then, the master-batch and pure PPR were extruded to prepare PPR specimens containing 0.1 wt.% WBG with the setting temperatures of 160–200 ◦ C from hopper to die and the screw speed of 120 rpm. After making droplets, the pellets were dried at 80 ◦ C for 12 h. Then, the pellets were injection-molded into standard specimens for testing, conducted on an injection machine (PS40E5ASE, Japan). The melt temperature was set as 240 ◦ C, and the mould temperature was 30 ◦ C. For comparison, the neat PPR specimens were also prepared by the identical processing parameters. The specimens were annealed in a vacuum oven at 110 ◦ C for 2 h, which was an optimized annealing condition. After being annealed, the specimens were cooled in the ambient air. For short, the samples of pure PPR, annealed PPR, ␤-nucleated PPR and annealed ␤-nucleated PPR were denoted as PPR, PPR-A, ␤-PPR and ␤-PPR-A, respectively.

Izod impact strength (kJ/m2)

phase with three equivalent glide planes. This indeed offers a great probability of favourably orientated crystals for slip with regard to the principal shear stress, and allows a more uniform deformation of the ␤ lamellae at reduced yield stress. The amorphous phase has also been supposed to take part in the plastic modification of PP containing ␤ phase crystals through higher intercrystalline tie chain density [24,25]. On the basis of crystallization kinetics and chain-folding regularity, the amorphous phase is believed to provide an uniform stress distribution over the crystalline lamellae in the case of the ␤ phase. A consequence of this is strong strain hardening for the ␤-phase that may notably account for the enlargement of the plastic zone at the crack tip. But, only controlling the amount of ␤-PP in the materials is not enough to obtain the high toughness. What is worse, some works found that the presence of ␤-phase had little effect at low temperature (below Tg ) for isotactic polypropylene [26]. Thus, the combination of rubber particles and ␤-modification has been considered. Largely improved fracture resistance, as well as the shift of brittle-ductile transition to lower rubber particles volume fraction, has been reported in the literatures [27–30]. Recently, the microstructure and mechanical behaviors of ␤PP have also been comparatively researched through annealing process at the elevated temperature between glass transition temperature (Tg ) and melting temperature (Tm ) [31–34]. Compared to ␣-PP with the interlocked structure, it seems more likely that ␤-PP has more potential to be toughened by secondary crystallization occurring in the amorphous phase during annealing. Some works have proved that, at a certain annealing temperature (130 ◦ C), the fracture resistance of ␤-PP can be largely improved [34]. With the improvement of the crystals induced by secondary crystallization, including the degree of crystallinity, molecular arrangement, and lamellae thickness, the fracture toughness in the bulk crystals increases because more fracture energy is required to destroy these improved crystal structure. In addition, the number of chain segments in the amorphous region decreases and some microvoids form, making the lamellae structures looser and more available to slip and/or elongate along the impact direction [34,35]. Although annealing is a useful way to improve the impact strength of ␤-PP at room temperature, the effect of annealing on the low temperature toughness of ␤-PP is few at hand. In current study, we attempted to improve the toughness of PPR at relatively low temperature by combination of ␤-nucleating agent and annealing. The injection-molded bars of PPR with a fixed ␤-nucleating agent content of 0.1 wt.% were prepared at 240 ◦ C, followed by annealing at 110 ◦ C for 2 h. It was found that the synergistic toughening effect at low temperature (0 ◦ C), through ␤-modification and annealing, was achieved. Numerous characterizations, including polarized light microscopy (PLM), X-ray diffraction (XRD), differential scanning calorimetry (DSC), optical microscopy (OM), scanning electron microscopy (SEM) and dynamic mechanical analysis (DMA), were performed to disclose the underlying toughening mechanism.

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Fig. 1. The notch Izod impact strengths for PPR, PPR-A, ␤-PPR and ␤-PPR-A tested at 0 ◦ C.

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Fig. 2. Typical impact-fractured surfaces morphologies of ␤-PPR and ␤-PPR-A. (a1) ␤-PPR and (b1) ␤-PPR-A at low magnifications; (a2), (a3), (a4) enlarged view of subregion marked in (a1); (b2), (b3), (b4) enlarged view of subregion marked in (b1). The arrow shows impact direction.

part-impact specimens were collected by the Leica DMIP optical microscopy. A Philips X’Pert pro MPD apparatus was adopted to acquire WAXD spectra. The analysis method of WAXD spectra was described by Huo et al. [36]. The overall crystallinity, Xc , was calculated according to the following equation:

Xc =





Acryst

Acryst +



Aamorp

(1)

where Acryst and Aamorp are the fitted areas of crystal and amorphous region, respectively. While the relative amount of the ␤-form crys-

tal, Kˇ , was evaluated according to the method of Turner-Jones et al. [37]

Kˇ =

Aˇ(3 0 0) A˛(1 1 0) + A˛(0 4 0) + A˛(1 3 0) + Aˇ(3 0 0)

(2)

where Aˇ(3 0 0) is the area of the (3 0 0) reflection peak of ␤-form at 2 = 16.1 ◦ ; A˛(1 1 0) , A˛(0 4 0) and A˛(1 3 0) are the areas of the (1 1 0), (0 4 0), and (1 3 0) reflection peaks of ␣-form, and correspond to 2 = 14.1◦ , 16.9◦ and 18.6◦ , respectively.

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Therefore, the crystallinity of ␤-form (Xˇ ) and ␣-form (X˛ ), could be calculated according to the following equation: (3)

X˛ = Xc − Xˇ

(4)

The calorimetric analysis was conducted on a Perkin-Elmer pyris-1DSC calorimeter calibrated by indium. The mass of testing specimen was around 5 mg. The specimens cut from the injectionmolded specimens were directly heated from 30 to 200 ◦ C at a rate of 10 ◦ C/min. The scanning electron microscopy (SEM) experiments were performed using a Hitachi S3400+EDX SEM instrument with an acceleration voltage of 20 kV. The impact fracture surfaces after low temperature test are examined. Dynamic mechanical analysis (DMA) testing was carried out using a DMA Q800 analyzer (TA instruments, USA). The three-pointbend mode was used, and the measurement was carried out on a rectangular shaped part in the size of 30 mm × 10.2 mm × 4.2 mm (length × wide × thickness) from −40 to 100 ◦ C at a heating rate of 3 ◦ C/min and an oscillatory frequency of 1 Hz. 3. Results and discussion

β-PPR-A

Relative Intensity

Xˇ = Xc × Kˇ

β-PPR

PPR-A PPR

10

15

It is well known that the overall crystallinity and polymorphic composition are two important factors determining the macroscopic toughness of semi-crystalline polymers. To quantitatively estimate the crystallinity (Xc ) and polymorphic composition in the samples, wide-angle X-ray diffraction (WAXD) was used. Fig. 3 illustrates the WAXD spectra of PPR, PPR-A, ␤-PPR and ␤-PPR-A. In these profiles, the (1 1 0) plane at 2 = 14.1◦ , (0 4 0) plane at 16.9◦ , and (1 3 0) plane at 18.6◦ are the principal reflections for ␣-crystals in PPR, while the (3 0 0) plane at about 16.1◦ is the principle reflection of the ␤-crystals. Quantitative estimations of total crystallinity

30

2θ( ) Fig. 3. The WAXD spectra of PPR, PPR-A, ␤-PPR and ␤-PPR-A. Table 1 Values of total crystallinity (Xc ), relative fraction of ␤-form (Kˇ ), and crystallinity of ␣-form (X˛ ) and ␤-form (Xˇ ) obtained by WXRD.

PPR PPR-A ␤-PPR ␤-PPR-A

Xc (%)

Kˇ

X˛ (%)

Xˇ (%)

40 44 45 51

0 0 0.35 0.32

40 44 29 35

0 0 16 16

(Xc ), relative fraction of ␤-form (Kˇ ) and crystallinity of ␣-form (X˛ ) and ␤-form (Xˇ ) have been done according to the methods mentioned in Section 2, and the values are reported in Table 1. It is shown that the total crystallinity increases after annealing for both virgin PPR and ␤-PPR. Moreover, the increase in the total crystallinity after annealing is more likely to result from the increase in the ␣-form, because the crystallinity of ␤-form is almost equal for both ␤-PPR and ␤-PPR-A (16% of ␤-form crystals). It is similar to some other reports that annealing of PP induces transformation of smectic mesophase into a polymorph without noticeable development of ␤ structures [38,39]. On the other hand, some similar results have been achieved through calorimetric analysis (DSC), as shown in Fig. 4. It is particularly worth noting that an obvious extra annealing melting peak

α β αΑ

β−PPR-A

Endo

3.2. Crystallinity and polymorphic composition

25

o

3.1. Toughness and fracture morphology On the basis of the research goal of this study, the notch Izod impact strengths for PPR, PPR-A, ␤-PPR and ␤-PPR-A were studied at 0 ◦ C, and the values are plotted in Fig. 1. It is found that the sample of PPR shows relative low impact strength at this testing temperature. It also shows that the impact strength of PPR has been improved a little by either annealing or ␤-modification. However, the sample of ␤-PPR-A shows obviously high impact strength (30.2 kJ/m2 ), which is almost five times as high as that of the virgin PPR (6.5 kJ/m2 ) and about three times as high as ␤-PPR (10.4 kJ/m2 ). So, it means synergistic effect of annealing and ␤-modification for toughening PPR. To clearly understand the variation of fracture resistance at low temperature related to the annealing process, the typical impactfractured surfaces of ␤-PPR and ␤-PPR-A were characterized and the results are shown in Fig. 2. Before annealing, the sample exhibits the typical surface feature of brittle fracture mode. The plastic deformation zone is relative small and smooth, and the weak plastic deformation is concentrated in the region not far away from the crack root. However, after being annealed, the whole fracture surface shows obvious plastic deformation and the extent of plastic deformation increases significantly. At the edge of the sample, a serrated surface can be distinguished, while in the center of the sample, shear deformation is observed. At higher magnification (see Fig. 2(a2–4, b2–4)), one can observe shear deformation of the matrix, indicating more energy dissipation during impact fracture. Therefore, annealing induces the brittle-ductile transition of ␤-modified PPR.

20

β−PPR

PPR-A PPR

100

110

120

130

140

150

Temperature (ºc) Fig. 4. The DSC melting curves of PPR, PPR-A, ␤-PPR and ␤-PPR-A.

160

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Fig. 5. PLM photographs of crystalline morphology of PPR, PPR-A, ␤-PPR and ␤-PPR-A. (a) PPR, (b) PPR-A, (c) ␤-PPR and (d) ␤-PPR-A.

appears in the annealed samples at relatively low temperature. Meanwhile, the original ␣ melting peak moves to higher temperature slightly. According to other’s works, the low-temperature melting peak was attributed to the melting of thin lamellae formed during the secondary long-time annealing process [34,40], while the high-temperature melting peak was attributed to the melting of thick lamellae generated from primary crystallization of the injection molding process and then thickening by annealing [41]. Therefore, it can be deduced that annealing at this temperature hardly changes the polymorphic composition, irrespective of adding ␤-nucleating agent. According to the DSC result, annealing leads up to somewhat increase in the amount of ␣-form, especially thin lamellae of ␣ crystal.

3.3. Crystalline morphology Since adding ␤-nucleating agent can affect the polymorphic composition and crystal morphology of PPR, the increase in the toughness can also be partly ascribed to the change of crystal morphology induced by ␤-nucleating agent. Also, whether the annealing might affect the crystal morphology needs to be further clarified. Thus, the effect of ␤-modification and annealing on the crystalline morphology of PPR has been investigated in detail by PLM, as shown in Fig. 5. For virgin PPR, typical ␣-spherulitic morphology is observed for both annealed and unannealed samples. For ␤-PPR, the size of crystal sharply decreases while compared with that of virgin PPR. More importantly, a lot of ␤-crystals with high birefringence character form into net-framework in the samples. For semi-crystalline polymers, the size and the integrity of spherulites, as well as the boundary strength between spherulites, also influence the fracture resistance [42]. That is to say, the fine-

ness of the crystal structure could be beneficial for increasing of fracture resistance. On the other hand, formation of ␤-phase can improve the molecular entanglement between the spherulites, thus giving rise to the increase of the boundary strength between the spherulites. This also contributes to the improvement of fracture resistance [43]. 3.4. Change of amorphous phase Deformation of a semi-crystalline polymer can be viewed as stretching of amorphous network with crystallites as physical cross-links. As Maiti et al. pointed out that annealing of PP at an elevated temperature results in (i) secondary crystallization of a part of the amorphous phase, (ii) thickening of radial lamellae, (iii) development of subsidiary lamellae, and (iv) growth of the crystal perfection [44]. Therefore, the amorphous network must be changed during annealing, reflected by variation of intrinsic deformation parameters, and thus could be responsible for the toughness enhancement. To explore the relationship between the fracture toughness and intrinsic deformation parameters, such as true yield stress and strain hardening modulus, would give a deep insight into the fracture and toughness of semicrystalline polymers. The deformation and failure of polymers can be described with a simplified model based on a stability analysis [45]: ty G

=

2 n −(1/n ) n −1

(5)

where  t and G are respectively the intrinsic yield stress and strain hardening modulus and n is the estimated “natural draw ratio” of the neck. Since the stress–strain curves of most injection-molded samples tested show similar behavior with respect to neck drawing,

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3

Storage modulus G' (MPa)

3.0x10

a

β−PPR β−PPR-A

3

2.5x10

3

2.0x10

3

1.5x10

3

1.0x10

2

5.0x10

-40

-20

0

20

40

60

80

100

o

Temperature ( c)

-1

1.5x10

b

β−PPR β−PPR−A

-1

1.0x10

representative nominal stress–strain curves of ␤-PPR and ␤-PPR-A are plotted only in Fig. 6, respectively. Under stretching, the samples behave as homogenous elastic deformation, followed by the formation of a stable neck. Stable necking provides an effective way to obtain the ratio of true yield stress  t to strain hardening modulus G as long as the natural draw ratio in the neck n is known. When equilibrium is reached between the load transferred in the neck and that in the undeformed zone, a stable neck is formed. At this point, the load in the neck will be large enough to induce yield in the adjacent undeformed material. So the true yield stress  t can be obtained from the maximum curvature of true tress strain curve around nominal yield point. The results of n ,  t , G and  t /G are summarized in Table 2. It is indicated that lowest values of n and  t /G exist in the samples with ␤ modification and being annealed. The main reasons for these changes might be clarified as follows: firstly, adding ␤-nucleating agent in PPR increases the crystallinity, decreases crystalline size and promotes the formation of ␤-crystals. Secondly the lamellar structure of ␤-crystal is looser than that of ␣-crystal. Therefore, the shear deformation ability in crystalline component has been improved. Thirdly, during annealing, secondary crystallization of a part of the amorphous phase increases the junctions of network, which makes large amorphous region can deform as an integrity. Therefore, it can be concluded that the improved toughness in the ␤-PPR-A samples is related to the enhanced amorphous phase that could effectively transmit stress from amorphous strands to the adjacent crystallites. On the other hand, the molecular mobility in the amorphous phase is another factor affecting the deformation of amorphous

Tan δ

Fig. 6. Typical nominal stress–strain curves of ␤-PPR and ␤-PPR-A.

-2

5.0x10

0.0 -40

-20

0

20

40

60

80

100

o

Temperature ( c) Fig. 7. (a) Storage modulus (G ) and (b) mechanical loss factor (Tan ı) as a function of temperature for ␤-PPR and ␤-PPR-A.

strands and thus toughness. It can be deduced from the mechanical relaxation, as shown in Fig. 7. After being annealed, the storage modulus (G ) value of the specimen increases. This can be ascribed to the improvement of crystallinity and crystalline structure during annealing. More importantly, the mechanical loss factor (Tan ı) curve is thought to more useful to exhibit the microstructure change [46]. The maximum at low temperature 15 ◦ C is related to ␤ relaxation, accounting for the glass transition of the unrestricted amorphous PP (Tg ), while the peak at the higher temperature 70 ◦ C is related to ␣c -relaxation, accounting for the relaxation of restricted PP amorphous chains in the crystalline phase (defects), also known as rigid amorphous molecules [47,48]. The Tg value of annealed sample moved to lower temperature compared with that of unannealed sample, indicating that annealing promotes the mobility of amorphous chain segments. This can be ascribed to a decrease in the concentration of the chain segments in the amorphous phase due to the second crystallization process

Table 2 The intrinsic deformation parameters deduced from tensile tests. Natural draw ration n PPR PPR-A ␤-PPR ␤-PPR-A a

a

5.7 (0.4) 4.8 (0.2) 4.4 (0.3) 2.9 (0.2)

Values in parentheses indicate standard deviation.

True yield stress  ty (MPa)

Strain hardening modulus G (MPa)

 ty /G

23.5 (0.4) 26.5 (0.7) 24.0 (0.5) 25.1 (0.5)

3.4 (0.2) 4.4 (0.4) 4.3 (0.3) 5.9 (0.3)

6.9 (0.2) 6.0 (0.3) 5.6 (0.3) 4.2 (0.3)

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Fig. 8. Crack-initiation pattern of PPR, PPR-A, ␤-PPR and ␤-PPR-A after the notch Izod part-impact test. (a) PPR, (b) PPR-A, (c) ␤-PPR and (d) ␤-PPR-A.

relating to the change of some chain segments from amorphous state to crystalline state during the annealing process [34,49,50]. Furthermore, the intensity of ␤-relaxation peak of annealed sample exhibits higher value than that of unannealed sample. As is well known, the intensity of ␤-relaxation is representative of the maximum energy dissipated because of the viscoelastic relaxation of PP component. This means that, after being annealed, the fracture resistance and damping capacity are improved simultaneously.

3.5. Crack-initiation term and toughening mechanism To determine the synergistic toughening mechanism of ␤modification and annealing for PPR, the crack-initiation term for PPR, PPR-A, ␤-PPR and ␤-PPR-A has been studied further. Fig. 8 shows the crack-initiation pattern of these samples after the notch Izod part-impact test. For both annealed and unannealed PPR, brittle behavior was observed, and ahead of the crack tip a single craze was visible. It may be caused by the ␣-form crystal with weak interaction between the spherulites with a little tie-molecules. On the other hand, for ␤-PPR, crack shows a broad craze and a certain amount of stress whitening at the end of the crack, indicating the presence of ␤ phase may benefit to the stress whitening. Interestingly, totally different from the front propagation behaviors, multiple crazing and a certain amount of stress whitening remain visible along the tips of the arrested crack for ␤-PPR-A. The trajectory of these features at the tip of crack shows obvious shear deform of the matrix. As a result big deform zone and large energy dissipation can be obtained. The toughening mechanism according to the behavior of crack initiation and propagation for ␤-PPR-A can be contributed by

several factors. Firstly, adding ␤-nucleating agent decreases the size of spherulites resulting in more tie molecules between the inter-lamellae. Secondly, as ␤ form crystal, the primary lamellar structure is much looser and the entangled chains of inter-lamellae are more than that of ␣ form. The effect of the ␤-modification is very likely associated with a specifics structure of the amorphous phase induced ␤-crystallites [51]. In addition, the ␤ phase favors the development of microcavites in deformed PP, which has been widely reported in literature [26,52,53]. Thirdly, after being annealed, the strength of amorphous phase increases according to the increase of junctions of the network by secondary crystallization. Also, the mobility of the chains in amorphous region has been largely improved. Therefore, the craze patterns of the annealed ␤-PPR sample show well-propagated branches. Combining all of these factors, the good low temperature impact toughness has been achieved.

4. Conclusion In this work, excellent low temperature impact toughness of polypropylene random copolymer (PPR) was obtained by combining ␤-modification and annealing. On the one hand, the decrease in the crystalline size and formation of ␤ form reduce the plastic resistance of crystals. On the other hand, because of the second crystallization, the strength of amorphous phase enhances and can transmit stress to the adjacent crystallites. What is more, the enhanced molecular mobility in the amorphous phase is another factor contributing to the high energy dissipation during fracture. This work proves a considerable way to increase the low temperature resistance of semicrystalline polymer by increasing the ability

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