Localized self-assembly and nucleation: a new strategy for preparing highly toughened polymer blends

Abstract

A facile methodology was developed to induce a core-pompon structure in the PA66/iPP system. By controlling the diffusion and self-assembly behavior of β-nucleating agents (NAs), we successfully localized the self-assembly and heterogeneous nucleation of NAs in the near interface region of PA66 spheres. The PA66-core/β-PP-pompon structure exhibited promising and applicable functionalities as toughening components.

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The wide application of engineering polymers has contributed to heated researches on advanced polymer–polymer blends in recent years.^1–4 Generally, the aim of incorporating a second polymer component into a given polymer matrix is to satisfy the increasing demand on the mechanical, thermal, optical or barrier properties of the material.⁵

Polypropylene exhibits a wide range of favorable properties, such as low-cost, light weight, ease of processing and chemical stability. Blending polyamide 66 (PA66) into isotactic polypropylene (iPP) has been attempted with the goal of preparing composite materials with improved paintability, wear resistance and barrier properties, etc., that would enable new applications.^6–9 However, the mutual repulsive molecular interaction of the two polymers often leads to high interfacial tension in their blend systems, resulting in immiscibility and poor interfacial adhesion throughout the whole range of the composition. Consequently, the mechanical properties of such materials, including tensile strength and impact resistance, have remained unsatisfactory. Compatibilizers such as grafted polymers⁹ or nano-fillers² were adopted to mend the poor interfacial contact between iPP and PA66. Large amounts of these compatibilizers (up to 10 wt%) were added to fully cover the surfaces of discrete PA66 spheres; however, concerns were raised regarding the release of such compatibilizers during use.^10–12

Interfacial crystallization was reported as a promising way to improve the interfacial adhesion in multiphase polymer systems; however, epitaxial crystallization in the PP/PA66 system was doomed to impracticality due to the crystallographic mismatch of the two polymers.^13,14 Inducing the formation of β-crystal form of PP, which usually involves the addition of small amounts of β-nucleating agents, has been shown to improve its impact resistance because β-crystal can absorb a large amount of impact energy via a β–α phase transition.^15–17 However, the β-crystal form often resulted in lessened tensile performance due to its relatively loose packing of molecules as compared to the α-crystals.^18–20 Confined nucleation leads to a hybrid system containing impact-absorbing β-crystals and stress-bearing α-crystals. Localized nucleation in the near-interface region could thus be a very tempting methodology for preparing advanced PA66/iPP materials, not only improving the interfacial adhesion, but also forming a buffer-area around the PA66 spheres.

In this communication, we present a novel and facile technique to improve the interfacial adhesion in PP/PA66 blends. By selective enrichment of the nucleating agents (NAs) in the PA66 phase, followed by elaborate control of its migration and self-assembly, we successfully induced crystallization of the iPP matrix around the near-surface region of the PA66 spheres. The resulting PA66-sphere/β-PP-crystal “core-pompon” structure in the α-crystal iPP matrix featured an impact-resistant area around the hard PA66 spheres while maintaining the 3-dimentional α-PP matrix as a robust and stress-bearing network. This article sheds some light on the structural design of high-performance PA66/iPP composites, as well as other advanced binary or multiphase polymer blend or composite materials.

Granules of PA66 were mixed with β-PP nucleating agents (NAs) in a weight ratio of 100 : 1 via twin-screw extruding. Then, the PA66/NAs master batch was pelletized and pulverized into a fine powder with grain size ∼30 μm. Pure PA66 powder was prepared as the control using the same procedure. We imbedded the neat and NAs-containing PA66 powders in between two pre-prepared iPP films to form sandwich structured composite films denoted as PA66/iPP and NAs-PA66/iPP, respectively, then carefully took them into the sample holder of a polarized optical microscope (POM) with a heating stage (detailed experimental description provided in ESI†).

The iPP crystals were melted at 190 °C; the samples were then heated up to 270 °C, which melted the PA66 phase to enable migration of NAs into the near-interface region. Heating the sample at 270 °C transformed the PA66 phase into its amorphous state, activating the motion of the macromolecular chains. As small molecules, the NAs dissolved in PA66 phase were undergoing intense Brownian movement at this temperature, leading to the migration of NAs from the PA66 phase (high concentration) to the iPP phase (low concentration). As a result, the NAs molecules, trapped in PA66 phase before, were “released” into the iPP phase through the interfacial area. However, due to the viscous nature of both polymer melts and the interfacial tension barrier, this release might be local, i.e., only a portion of the NAs migrated to the near interface region. The diffusion of NAs out of the PA spheres is shown in Fig. 1a. After that, the film was quenched to 140 °C for isothermal crystallization. At this temperature (140 °C), NAs are known to self-assemble (crystallize) into nanostructures,^15,19 thus shutting off their further long-range migration into the i-PP phase. After isothermal treatment for over 180 s, NA nanostructures started to act as heterogeneous nucleating points, inducing epitaxial crystallization of i-PP in the near-interface region (Fig. 1d and e). Given the type of NAs used in the study, this interfacial crystallization layer was composed of β-form crystals. The β-crystals continued to grow to more than 20 μm in thickness after 210 s isothermal treatment (Fig. 1f). Subsequent quenching of the sample to room temperature led to fast self-nucleation and crystallization of the melt. As a comparison, PA66/iPP system was tested using the same temperature ramp. We observed no interfacial crystallization, even after 1800 s of isothermal treatment. However, self-nucleation and crystallization of iPP was observed beginning at 600 s, resulting in spherulites throughout the entire sample as shown in Fig. 1a′–f′. In the case of neat PA66, which exhibited no evident heterogeneous nucleating effects on iPP, the self-crystallization mode of iPP was predominately in the α-form crystals.

Fig. 1 POM photos of (a) NAs-PA66/iPP and (a′) PA66/iPP at 270 °C. The samples were isothermally crystallized at 140 °C for different times: (b) 0 s; (c) 120 s; (d) 180 s; (e) 210 s; (f) 270 s; and (b′) 0 s; (c′) 300 s; (d′) 600 s; (e′) 1200 s; (f′) 1800 s. Low-magnification images are provided in ESI-Fig. 1.†

Scanning Electron Microscopy (SEM) was adopted to get more detailed structural information, down to the sub-micron level. As shown in Fig. 2a and a′, radial sheaf-like crystal structures rooted in the interfaces of the PA66 spheres were observed in fully-crystallized NAs-PA66/iPP. The morphology of the crystals was in accordance with previously reported NA-nucleated β-iPP crystals.^15,19 On the contrary, the smooth surface of PA66 particles with iPP spherulites surrounding them was observed in the PA66/iPP system, as shown in Fig. 2b and b′. These observations are in line with the POM observations. Without the assistance of a nucleating agent, iPP crystalized into α-form spherulites under quiescent conditions due to the thermodynamic advantage of α-form over β-form.²¹

Fig. 2 SEM micrographs of the microstructure in near-interface region: (a) NAs-PA66/iPP; (b) PA66/iPP.

Three driving forces are considered as the controlling factors for the migration/diffusion of NAs in polymer composites: interfaces, concentration and temperature. The immiscible morphology of the PA66/iPP blend firstly offered numerous interfaces for inter-phase diffusion of small molecules (NAs). Selectively enriching the NAs in the PA66 phase constructed a high concentration gradient, which directed the “release” of the molecules from NAs-rich PA66 spheres into the iPP matrix. The NAs used in this study, TMB-5, have two temperature-dependent phases itself: the dissolved small-molecule phase and the self-assembled micron-complex phase.^14,18 The former possesses much better mobility. Heating up to 270 °C not only dissolved the NAs in the PA66 phase, but it also melted both PA66 and iPP crystals into the mobilized rubber state. Coupled with the Brownian movement of NAs molecules, heat-induced mobilization of the system triggered the “release” of NAs molecules, as shown in Fig. 3. Subsequent quenching to isothermal temperature (140 °C), which was chosen to activate the self-assembly of NAs, confined the migration of NAs in the near-interface region. Hence, localized self-assembly of NAs in the iPP matrix was achieved around the PA66 spheres, forming micrometer sized complex NAs crystals, which possessed strong heterogeneous nucleating effects on iPP matrix. Areas of iPP distant from the PA66 spheres were not influenced by NAs due to their limited diffusion kinetics (controlled by time, temperature etc.) and therefore crystallized into the expected α-form crystals.

Fig. 3 Schematic illustration of the formation of core-pompon structure in NAs-PA66/iPP system.

Functional composite materials were easily engineered for specific applications using the aforementioned methodology. The composite materials were fabricated by firstly mixing the pre-made NAs-PA66 powder into iPP powder in a weight ratio of 5 : 95. Then, the powder mixture was hot-compressed at 270 °C, followed by isothermal treatment at 140 °C to trigger the migration and localized self-assembly of NAs in near-interface region of the iPP matrix. The iPP β-crystal layer, thus formed through epitaxial crystallization around PA66 spheres, significantly improved the system's toughness. As shown in Fig. 4a, the NA-PA66/iPP composite exhibited impact strength of up to 2.51 kJ m⁻², a significant improvement as compared to neat iPP (1.69 kJ m⁻²) and common blend PA66/iPP (1.84 kJ m⁻²). It is noteworthy that the impact resistance of the ternary blend PA66/iPP/NAs (1.78 kJ m⁻²), which also contained β-crystals, was lower than that of the composite material NA-PA66/iPP. This might be ascribed to the poor interfacial contact between PA66 and iPP, resulting in fast crack formation and propagation. Forming a β-crystal layer around the interfaces of PA66 spheres not only improved the interfacial contact to prevent crack initiation, but also acted as an effective “buffer-cushion” to absorb and dissipate the impact energy and suppress crack propagation.

Fig. 4 Impact strength (a) and elongation at break (b) of neat iPP; binary common blend (PA66/PP); ternary blend PA66/iPP/NAs and “core-pompon” structured NAs-PP/PA66 (TMB-5).

Localized nucleation of β-crystals in the near-interface region gave the composite relatively good flexibility as well. As shown in Fig. 4b, the elongation at the break of NAs-PA66/iPP reached 178%, a dramatic improvement compared to all the other systems. Moreover, the tensile strength and Young's modulus of NAs-PA66/iPP showed no difference to neat iPP, being ∼32 MPa and ∼250 MPa. That's because in iPP matrix areas distant from the interface, the α-crystal still dominated to be a stress-bearing framework.

In summary, a facile strategy to fabricate PA66/iPP composite with novel core-pompon structures was reported. Through localized self-assembly and nucleation, we successfully introduced β-form iPP crystallization on the surfaces of discrete PA66 spheres. The pompon-like β-crystals surrounding the PA66 spheres not only improved interfacial adhesion in the system to give proper flexibility, but also acted as a buffer zone to enhance the impact resistance of the composite. This methodology for structural control could serve as inspiration for fabricating other advanced binary or multiphase polymer blend/composite materials.

Acknowledgements

This work is financed by the National Natural Science Foundation of China (51127003, 51303114, and 51421061) and the Doctoral Fund of Ministry of Education of China (20130181120056), State Key Laboratory of Polymer Materials Engineering (Grant No. sklpme2016-3-05) and the Engineering Research Center of Marine Bioresources Comprehensive Utilization, SOA, (MBRCU201601).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6ra18461h

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