Introduction

Isotactic polypropylene (PP) is one of the most important polymers due to its excellent physical properties and low manufacturing cost, which has been widely used in industrial fields [1]. However, the application of PP is still hampered by its low impact toughness, especially at low temperature [2, 3]. In the past three decades, various toughening approaches have been investigated, including blending with rubber or elastomer [4,5,6], adding filler [7, 8] or β-nucleating agent [9] and copolymerizing with olefin monomer [10, 11].

In general, blending is a more economic and effective way among these common methods. It has been reported that PP was blended with various elastomers, such as ethylene–propylene–diene terpolymer (EPDM) [12, 13], ethylene–propylene random copolymer (EPR) [14, 15], styrene–ethylene/butylene–styrene triblock copolymer (SEBS) [16,17,18], ethylene–octene copolymer (POE) [19], etc. Liang and Li [20] summarized the advances in the research of toughening methods and the theoretical explanation of the toughening mechanism for PP/elastomer blends. Although some achievements have been made in the field of elastomer-toughened PP, the modified mechanical properties of the blends cannot fully meet industrial requirements because of strong immiscibility or incompatibility. In order to improve the miscibility between elastomers and PP matrix, many researchers have done a large amount of research. In the 1990s, Dow Chemical commercialized ethylene-α-olefin copolymers with controlled molecular weight, molecular weight distribution and comonomer composition, and proved that the elastomer is more effective for increasing the impact strength of PP due to the lower interfacial tension between the elastomer and the PP matrix [21,22,23,24]. However, the miscibility in binary blends of PP and ethylene-α-olefin copolymer is not as expected, mainly because of the molecular composition and primary structure of ethylene-α-olefin copolymer. Last several years, a new olefin block copolymer (OBC) with alternating arrangement of hard segments and soft segments in the molecular chain was developed by Dow Chemical via a processing called chain-shuttling polymerization [25, 26]. Liu et al. [27, 28] systematically studied the toughening mechanism of OBC for PP and found that a good toughening effect was obtained by proper dispersion and vulcanization of the rubber phase. However, it was difficult to prepare a blend with well-dispersed rubber phase and by controlling the vulcanization.

More recently, the advent of a novel propylene-based random copolymers copolymerized by propylene and ethylene (PEC) has drawn extensive interests of academics and industries, due to the good processability and excellent mechanical properties [29, 30]. PEC elastomer with accurately controllable structure and property was manufactured by ExxonMobil Chemical Co. through special discrete metallocene catalysis and solution polymerization. Such elastomer containing more than 85 mol% propylene shows a chain structure composed of stereoregular PP microcrystalline region and loose amorphous region. Therefore, its properties are different from those of POE and usual EPR due to its low ethylene content.

In practical applications, PEC can be used as finished products independently as well as toughening modifier for olefin polymers. Although PEC has excellent performance and wide application prospects, there are few reports on morphology and properties of the polymer blending system of PEC toughening PP. Therefore, in this work, we systematically studied phase morphology, miscibility, mechanical properties and rheological properties for PP blends toughened by PEC elastomer in order to establish the relationship between morphology and properties. It is hoped that the addition of PEC elastomer can be an effective tool for tailoring the properties of PP and making it possible to prepare new PP-based thermoplastics.

Experimental

Materials

Isotactic polypropylene (PP T30S), with a melt flow index (MFI) of 3.4 g × 10 min−1 (190 °C, 2.16 kg) and weight-average molecular weight (Mw) of 2.5 × 105 g/mol was purchased from Jihua Petrochemical Co., China. PEC elastomer (Vistamaxx™ 6202), a random propylene–ethylene copolymer with an Mw of 1.7 × 105 g/mol was provided by ExxonMobil, USA. It exhibited an MFI of 20 g × 10 min−1 (230 °C, 2.16 kg) and density of 0.863 g/cm3 (23 °C). The content of propylene in the copolymer was 85 mol%. All the raw materials were used without any pretreatment.

Preparation of the blends

Various PP/PEC blends were prepared by a melt mixing in an internal mixer (Haake Rheomix 600, Karlsruhe, Germany). The melt compounding was performed at a screw speed of 60 rpm and 180 °C for a mixing time of 10 min. Then all the samples were compression molded at 190 °C followed by quenching at ambient temperature to form around 1.0-mm-thick sheets for characterization. The neat PP and PEC were also subjected to the same treatment in order to have the same thermal history compared to the blends. The PP/PEC blends contained 10, 20 and 30 wt% PEC and were denoted as PECx, where x represented the weight percentages of PEC.

Characterizations

Dynamic mechanical analysis (DMA) was performed using a dynamic mechanical analyzer SDTA861e (Mettler Toledo) in a mode of tensile–compression. The gauge dimensions of the test samples were 20 × 10 × 1 mm3. All samples were heated from −60 to 100 °C. The heating rate and frequency were 3 °C/min and 10 Hz, respectively. The dynamic loss factor (tan δ) and the storage modulus (\( E^{\prime } \)) were determined as a function of temperature.

The morphology of the cryo-fractured and impact-fractured surface was investigated using a field emission scanning electron microscopy (FEI Co., Eindhoven, Netherlands). Before SEM characterization, the surfaces of the samples were sputtered with a thin layer of gold.

Thermal analysis was investigated by a Q20 differential scanning calorimeter (DSC) (TA Instruments, USA). The indium standard was used to calibrate the temperature and enthalpy. The samples had a nominal weight of 5–8 mg, and the measurements were taken in N2 atmosphere. All the samples were first heated form 25 to 190 °C at a heating rate of 40 °C/min, held for 5 min to erase the thermal history. Then they were cooled to 40 °C at a cooling rate of 20 °C/min (first cooling); the second heating scan was performed from 25 to 190 °C at a heating rate of 10 °C/min. The second melting and first cooling curves of the samples were recorded. The melting temperature (Tm) and melting enthalpy (ΔHm) were determined from the melting curves. The crystallization temperatures (Tc) and crystallization enthalpy (ΔHc) were determined from the cooling curves.

The degree of crystallinity (Xc) of the blends was calculated using Eq. (1).

$$ X_{\text{c}} = \frac{{\Delta H_{\text{m}} }}{{\left( {\Delta H_{\text{m}}^{0} \varphi } \right)}} \times 100\% $$
(1)

where φ is the mass fraction of PP in the blends, and \( \Delta H_{\text{m}}^{0} \) is 207 J/g for 100% crystalline PP [31].

Wide angle X-ray diffraction (WAXD) was carried out with an X-ray diffractometer (Rigaku model D/max 2500 VB2t/PC). The Cu Kα radiation (λ = 0.154 nm) source was operated at 200 mA and 40 kV. The measurement was operated in the range of 5°–40° at a scanning rate of 3°/min.

Uniaxial tensile tests of neat PP and various blends were determined by an Instron 1211 testing machine (Canton, MA). The measurements were conducted at a crosshead speed of 200 mm/min with dumbbell-shaped specimens (20 × 4×1 mm3). The notched Izod impact strength was measured by using a CEAST impact machine according to GB1843-93 (China). All tests were conducted at room temperature and relative humidity of 50%. The obtained values were the average of at least 5 measurements. The fracture surfaces of samples after impact tests were observed using SEM.

Rheological measurements were carried out by a rotational rheometer (TA Series AR2000ex, TA Instrument, USA) equipped with 25 mm plate–plate geometry. The samples were pressed into 1.0 mm thick at 190 °C. Frequency sweep for the specimen was performed at N2 atmosphere in the fixture. The linear viscoelastic region of the samples was determined by a dynamic strain sweep test. The strain and angular frequency ranges in this study were 5% and 0.05–100 rad/s, respectively.

Results and discussion

Miscibility

It is well known that in DMA measurement, the miscible polymer blends exhibit a single glass transition temperature (Tg) due to the homogeneous phase in amorphous region, whereas the immiscible blends exhibit two Tg’s corresponding to each individual component.

Figure 1 shows the DMA traces of PP/PEC blends with different PEC contents. Glass transition, that is, relaxation of motion units from a ‘frozen’ state to a ‘free’ state, is characterized by the abrupt transformation of tan δ in Fig. 1a and listed in Table 1. Neat PEC exhibits a single sharp relaxation peak at around − 24 °C, corresponding to the glass transition. Neat PP shows a glass transition of around 2 °C. As for PP/PEC blends, two peaks corresponding to the glass transition of each component are observed, suggesting that the blends are not thermodynamically miscible. On the other hand, it can be found that the Tg of PEC shifts to higher temperatures, while the Tg of PP shifts slightly to lower temperature with the increase in PEC content. Such variation of Tg’s with PEC content indicates that there is a inter-compositional interaction between PP and the PEC in amorphous region [32]. Parts of PP sequences in the PEC chains are incorporated into the crystal of PP and the remainder is extruded into the amorphous region in the blends. This structure enhances the efficient diffusion and mutual incorporation between the PP amorphous region and the PEC molecules. The flexible PEC molecules enhance the motion of the amorphous PP segments. On the contrary, the rigid PP microcrystalline weakens the motion of the PEC components. Therefore, the decrease in the Tg for PP and the increase in Tg for PEC are caused by such inter-compositional interaction, and the change in Tg is more pronounced with the increase in the PEC content [33]. So we suggest that the blends composed of PP and PEC components are partial miscibility.

Fig. 1
figure 1

DMA traces of PP/PEC blends with different PEC contents: a tan δ versus temperature; b storage modulus versus temperature

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Table 1 Thermodynamic parameters of PP and PP/PEC blends
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The storage modulus (\( E^{\prime } \)) curves of the blends with different PEC contents are shown in Fig. 1b. For all blends, significant depression of \( E^{\prime } \) can be seen after the temperature higher than the Tg, suggesting the addition of PEC elastomer has a great influence on the elastic property of neat PP attributed to the low stiffness of the PEC.

Phase morphology

It is well known that, for polymer blend, the interfacial structure and the dispersed phase structure strongly affect its mechanical properties. Therefore, it is necessary to clarify the phase morphology of PP/PEC blend and its effect on the mechanical properties. Figure 2 shows the SEM micrographs of cryo-fractured surfaces for neat PP and PP/PEC blends. As shown in Fig. 2a, the fractured surface of neat PP appears smooth, typical continuous phase structure. Figure 2b–d shows a much fine and uniform dispersion of the PEC dispersed phase, attributed to the inter-compositional interaction between PP and the PEC. It is obvious that the size of elastomer particles is not affected by the increasing of PEC content, whereas the number of the elastomer particles increases and thus result in decrease in distance between particles. In general, for the traditional PP/elastomer blends, such as PP/EPDM blends, the size of the dispersed phase particles is significantly affected by the EPDM content [13]. As well known, the increase in the content of elastomer is favorable for enlarging the size of dispersed particles and broadening size distribution [34, 35]. On the other hand, the PP sequences in the PEC chains have the compatibilizing effect, which not only reduce the interfacial tension between PEC dispersed phase and PP matrix but also decrease the viscosity ratio of the two components, thereby restricting increase in the size of dispersed particles and widening of size distribution. The evolution of interfacial tension and viscosity ratio after addition of the PEC elastomer are favorable for uniform dispersion of elastomer particles. Therefore, it is the synergy of blend composition, interfacial tension and viscosity ratio that determines the phase morphology of the blend.

Fig. 2
figure 2

SEM micrographs of cryo-fractured surface of neat PP and PP/PEC blends: a Neat PP, b 10% PEC, c 20% PEC, d 30% PEC

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Thermal properties

Figure 3 shows the DSC traces of PP/PEC blends. The crystallization and melting temperatures obtained from the first cooling and second heating run, respectively, are listed in Table 1. These parameters are helpful to further understand the interaction/miscibility between neat PP and PEC components. As shown in Fig. 3, neat PP exhibits a melting temperature of 163 °C and a crystallization temperature of 113 °C, respectively. For all the blends, Tm’s decrease slightly with the addition of PEC. Such slight decrease in Tm may be due to the partial miscibility between PP matrix and PEC elastomer. As discussed above, the PP sequences in PEC molecular chains are involved in the crystals of neat PP, which in turn interferes with the perfection of the PP crystals to some extent. Therefore, the melting temperature of the blends is decreased compared with that of pure PP. On the other hand, Tc’s decrease with the increase in PEC content. We believe that the variation of Tc is caused by two reasons. First is the imperfection of PP crystals caused by the insertion of the PP sequences from the PEC chains, and second is the dilution effect from the remainder of the PEC molecular chains. Both reasons slow down the diffusion of the PP chains, inhibit their orderly arrangement, and ultimately lead to a slower crystal growth rate. This is also the reason why the degree of crystallinity for the blends is less than that of pure PP.

Fig. 3
figure 3

DSC thermograms for neat PP and PP/PEC blends: a first cooling at 20 °C/min, b second heating at 10 °C/min

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In order to reveal the effect of PEC elastomer on the crystalline structure of PP, WAXD experiments were performed and the results are shown in Fig. 4. Neat PP presents five strong diffraction peaks at around 14.0°, 16.8°, 18.4°, 21° and 21.8°, corresponding to (110), (040), (130), (111) and (041) lattice planes of α-form crystals of PP [36, 37]. It can be noted that the position of characteristic diffraction peaks of PP crystal in the blends hardly change regardless of the composition variations, suggesting that the PP unit cell never change. Therefore, the PP sequences in the PEC chains are incorporated into the neat PP crystals according to the α-form, without changing the crystalline structure of PP, but the crystal perfection is reduced.

Fig. 4
figure 4

WAXD patterns of neat PP and PP/PEC blends

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Mechanical properties

The effects of PEC elastomer on the mechanical properties of PP were investigated by tensile and impact test. Figure 5 shows the stress–strain curves of PP/PEC blends, and the corresponding mechanical data are summarized in Table 2. It can be seen that the neat PP exhibits typical characteristics of brittle and stiff material with an elongation at break of only 16%, and there is no distinct yield point before tensile fracture. In contrast, the blends show a distinct yield process and a stable necking growth. When the PEC content was 10%, the elongation at break was 350%, which was increased by more than 20 times in the blends. At the same time, the yield tensile strength was 38.4 MPa, only a decrease in < 6%. Blending a small amount of PEC can significantly improve flexibility of PP without reducing the yield tensile strength apparently, which indicates that PEC is an effective toughening agent for improving the mechanical properties of PP.

Fig. 5
figure 5

Tensile stress–strain curves of neat PP and PP/PEC blends

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Table 2 Mechanical properties of the PP and PP/PEC blends
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The impact strength analysis is an important tool for studying the toughening mechanism of elastomer modified polymer blends. Figure 6 shows the effect of PEC content on the impact strength of PP/PEC blends and Fig. 7 shows the morphology of the impacted PP blends. The impact strength of pure PP is only about 5 kJ/m2, and there is no stress-whitening zone (Fig. 7a). Obviously the samples fractured in a brittle manner. An increase in PEC content in blends leads to a gradual improvement in toughness, accompanied by a certain degree of stress-whitening (Fig. 7b–c), and the brittle–ductile transition occurs when the PEC content is about 20 wt% because the impact strength sharply rises to 40 kJ/m2 at this point. Blend with 30 wt% PEC could not be fractured completely under the same test conditions because of its good impact toughness, as shown in Fig. 7d. (The cross-sectional area used to calculate the impact strength for the blends with 30 wt% PEC was obtained from the fracture part.) Figure 8 shows SEM micrographs of notched impact-fractured surfaces of neat PP and the blends. Neat PP shows good rigidity due to its perfect crystal structure. However, when it is subjected to impact, brittle fracture occurs immediately because of stress concentration, and the fracture surface is smooth and flat, as shown in Fig. 8a. For all the PP/PEC blends, the rough surface and no obviously separate particles can be observed (Fig. 8b–d), indicating good adhesion between the PP matrix and elastomer particles. On adding PEC elastomer, the PP sequences in the PEC chains act as a tie molecules, connecting adjacent lamellae of the PP crystals, and the rest of PEC chains play a role in softening, improving the mobility of PP chains, all of which help to improve the impact strength. In the case of PP/PEC blends with PEC content is 10 wt%, despite the presence of elastomer particles, there is still some stress concentration. This is due to the fact that the particles are greatly separated in the PP matrix, so the stress fields around the elastomer particles have little effect on each other. Such phase morphology eventually leads to a small increase in impact strength (Fig. 8b). Otherwise, along the crack-propagation direction, many fibrils can be observed on the fractured surfaces for the blend with 20 wt% PEC (Fig. 8c). As the interparticle distance shortens, the stress fields around neighboring elastomer particles would interact and overlap, which contributes to generation of shear yielding and plastic deformation of PP [33]. In addition, the rearrangement of PP segments at the interphases during impact test, including the merged the PP sequences in the PEC chains, and the dilution of the non-crystallizable portion of the PEC chains make it difficult for the blends to form stress concentration. Therefore, ductile fracture occurs for this blends. As far as the blend with 30 wt% PEC is concerned, the plastic deformation of PP matrix due to the increase in the number of rubber particles is more pronounced, manifesting with the appearance of discernible cavities on the fractured surfaces (Fig. 8d), which enables the blends to possess higher impact toughness.

Fig. 6
figure 6

Variation of impact strength for PP/PEC blends with PEC content

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Fig. 7
figure 7

The sample morphology of the impacted PP blends sheets

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Fig. 8
figure 8

SEM images of impact-fracture surface of PP and PP/PEC blends: a neat PP, b 10% PEC, c 20% PEC, d 30% PEC. The arrow indicates the impacting direction

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In rubber-toughened polymer blends, the dispersed rubber particles are likely to elongate along the stress direction under impact. Adjacent stress fields around rubber particles interact and overlap, which induces shear yield of the matrix. Then, matrix deforms with the rubber particles as a whole. The cooperative motion of matrix and rubbery phase can effectively dissipated large amounts of fracture energy, resulting in a highly increase in impact strength [38, 39].

Viscoelastic behaviors

In general, the rheological properties play an important role in optimizing the appropriate processing conditions. In the melt mixing process, the rheological characteristics of the components will control the morphology development of the blends. Figure 9 shows the complex viscosity (\( \eta^{*} \)), dynamic storage modulus (\( G^{\prime } \)) and dynamic loss modulus (\( G^{\prime \prime } \)) of neat PP, PEC and PP/PEC blends as a function of the frequency at 190 °C. It is obvious that \( \eta^{*} \) of neat PP exhibits non-Newtonian shear-thinning behavior, while the \( \eta^{*} \)of neat PEC shows slight shear-thinning behavior at high frequency and near-Newtonian behavior at low frequency as shown in Fig. 9a. In addition, \( \eta^{*} \) of neat PP is greater than that of PEC at all frequency range. It is also observed that \( \eta^{*} \) of the blends decrease with the increase in PEC content and show shear-thinning behavior. The viscosity decrease in the blends compared with neat PP is due to the dilution effect of the PEC elastomer on the PP melt, which is caused by the higher melt mass flow rate of the PEC elastomer. It is well known, the storage modulus (\( G^{\prime } \)) represents the stored energy, while loss modulus (\( G^{\prime \prime } \)) can be considered as the dissipated energy. The effect of frequency on the \( G^{\prime } \) and \( G^{\prime \prime } \) may infer the relative motion of molecules and show significant information about flow behavior of the melts. Figure 9b and c show the dependence of \( G^{\prime } \) and \( G^{\prime \prime } \) for neat PP, PEC and various PP/PEC blends at 190 °C. From Fig. 9b, it can be seen that neat PEC chains relax fully and present typical terminal behavior, that is, at low frequencies G is approximately proportional to ɷ2, which is consistent with linear viscoelastic theory. Moreover, the \( G^{\prime } \) of neat PP exhibits more obvious non-terminal behavior, indicating a slower relaxation behavior due to the longer molecular chain and more regular chain structure of PP. For the blends, with increasing PEC content, the \( G^{\prime } \) decreases over the whole frequency range compared with neat PP, which could be attributed to the dilution effect of the PEC elastomer that enhance the long-range motions of PP chains [40]. The dependence of \( G^{\prime \prime } \) on the PEC content is also better: \( G^{\prime \prime } \) decreases slightly with the increase in PEC content. The results are significant because they show that the optimum processing conditions of the blends may be very different compared to neat PP. As we all know, the processing methods of PP are diverse, and therefore, the decrease in \( \eta^{*} \) and \( G^{\prime } \) of the blends is beneficial to some processing properties, such as injection molding, cast film, etc.

Fig. 9
figure 9

Rheological properties of neat PP and PP/PEC blends at 190 °C: a the complex viscosity \( \eta^{*} \), b the storage modulus \( G^{\prime } \), c the loss modulus \( G^{\prime \prime } \)

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Conclusions

In this work, we systematically investigated the morphology, miscibility, mechanical properties and rheological properties of the polymer blending system of PEC elastomer-toughened PP. It was found that the miscibility of the two immiscible components was improved to some extent due to the intermolecular interactions between PEC components and neat PP. The crystallizable PP sequences in PEC chains could participate in the crystallization process of neat PP, and they together acted as the matrix of PP/PEC blends. At the same time, the remainder of the PEC molecular chains served as the dispersed phase. In addition, the PEC content significantly influenced the phase morphology of the PP/PEC blends. With the addition of PEC, the blends exhibited a homogeneous phase morphology, namely the size and size distribution of the elastomer particles remained unchanged, and only the distance between the particles became smaller. Such evolution of phase morphology with the content of PEC was beneficial for the improvement of impact resistance. The complex viscosity \( \eta^{*} \), storage modulus \( G^{\prime } \) and loss modulus \( G^{\prime \prime } \) of the PP/PEC blends were decreased with the increase in PEC elastomer due to the dilution effect of the PEC elastome, which was beneficial to the processing properties of PP.