Synergistic Modification of Recycled PET Using Halloysite Nanotubes and a Reactive Terpolymer for Enhanced Toughness and Processability

TL;DR

This study improves the toughness and processability of recycled PET by combining halloysite nanotubes and a reactive terpolymer, enabling its use in higher-value applications.

Contribution

The novel synergistic modification using HNTs and E-MA-GMA enhances recycled PET properties beyond individual additives.

Findings

The combined use of HNTs and E-MA-GMA significantly restricts PET molecular chain relaxation.

Recycled PET composites with HNTs and E-MA-GMA show 2.28 times higher impact strength under high-mold-temperature injection molding.

The synergistic modification reduces crystallinity, improving impact resistance after high-temperature treatments.

Abstract

Polyethylene terephthalate (PET) has become the predominant material for single-use packaging owing to its cost and performance advantages. However, massive post-consumer waste leads to environmental concerns, and recycled PET from thermomechanical processing followed by chain extension often suffers from low toughness and poor processability, restricting its use to low-value applications. In this study, halloysite nanotubes (HNTs) and ethylene–methyl acrylate–glycidyl methacrylate random terpolymer (E-MA-GMA) were melt-blended with recycled PET to examine their synergistic modification effects. The DSC results show that HNTs retain a nucleating effect on recycled PET even with the co-addition of E-MA-GMA, albeit with a substantial reduction compared with their effect when used alone. Nevertheless, rheological measurements indicate that the combined introduction of E-MA-GMA and HNTs imposes a significantly stronger restriction on the relaxation behavior of recycled PET molecular chains than the individual addition of either HNTs or E-MA-GMA. This is attributed to the interfacial reactions between E-MA-GMA and the recycled PET matrix, as well as between E-MA-GMA and HNTs, leading to the formation of branching and hybrid structures. This synergistic restraint markedly reduces the crystallization growth rate of PET. As a result, the recycled PET/E-MA-GMA/HNTs composites maintain relatively lower crystallinity compared with the recycled PET/E-MA-GMA composite after high-temperature injection molding or annealing treatment, leading to superior impact resistance. The impact strength of the recycled PET/E-MA-GMA/HNTs composites is 2.28 and 2.14 times that of the recycled PET/E-MA-GMA composite under high-mold-temperature injection molding and annealing conditions, respectively. The approach presented here facilitates the substitution of virgin plastics with recycled PET in demanding applications.

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1. Introduction

Polyethylene terephthalate (PET) has become a dominant packaging material due to its high strength, chemical stability, and low cost. However, the large volume of post-consumer PET leads to resource waste and environmental concerns. Thus, the efficient recycling and high-value utilization of waste PET are crucial to alleviating resource scarcity and promoting the green transformation of industries [1,2,3,4,5,6,7]. Among various recycling methods, a common approach involves thermomechanical recycling [8] followed by chain extension. However, recycled PET often suffers from insufficient toughness and processability, which limit its use to low-value-added applications. Therefore, enhancing the properties of recycled PET and enabling its high-value utilization have become urgent issues to be addressed.

Blending PET with rubber/elastomer [9,10,11,12] is a principal strategy for toughening PET. In this approach, the dispersed elastomer phase acts as a stress-concentration site, inducing shear yielding, crazing, or cavitation in the matrix; the resulting energy dissipation enhances the impact toughness of PET [11]. The effectiveness of the toughening depends on the dispersion morphology of the elastomer and the interfacial adhesion strength of the two phases, with compatibility as the key.

Loyens et al. [11] compared the toughening effects of different elastomer modifiers on PET. Their study demonstrated that systems containing glycidyl methacrylate (GMA) can achieve compatibilization at the phase interface through reactions between its epoxy groups and the terminal carboxyl or hydroxyl groups of PET, thereby significantly enhancing its impact toughness. In contrast, systems containing maleic anhydride (MA) exhibit inferior toughening effectiveness due to the lack of efficient interfacial reactions. The study emphasized that the interparticle distance (IDc) of the dispersed phase is a critical parameter influencing the toughening efficacy, and efficient toughening can be achieved when it falls below a critical value (0.1 μm). Among various systems, the ethylene–glycidyl methacrylate copolymer (E-GMA) provided the most effective compatibilization, resulting in a composite whose impact strength was increased by a factor of 15. Barati et al. [13] investigated the effects of a series of compatibilizers on recycled PET/rubber blends and found that the ethylene-acrylate glycidyl-methacrylate terpolymer (E-MA-GMA) exhibited the most pronounced compatibilizing effect, leading to the largest improvement in impact strength. Raiisi-Nia et al. [10] reported that adding 20 phr of GMA-grafted acrylonitrile butadiene rubber significantly enhanced the interfacial interaction between recycled PET and rubber phases, resulting in an impact strength that is 550% higher than that of neat recycled PET, thereby achieving effective toughening. Walber et al. [12] likewise verified that the material’s impact strength was notably enhanced with the incorporation of 15 or 20 wt% of E-GMA or E-MA-GMA.

In summary, GMA-containing compatibilizers can significantly enhance the impact strength of PET via a reactive compatibilization mechanism, among which the toughening efficiency of GMA-based copolymers (e.g., E-GMA, E-MA-GMA) is particularly prominent [11,13]. In comparison with the binary copolymer E-GMA, the ternary copolymer E-MA-GMA can maintain an amorphous state during processing [12]. However, the epoxy groups in GMA can react with the terminal groups (-OH/-COOH) of PET to form branched or longer linear structures, which retards the ordered arrangement process of PET molecular chains on crystal nuclei [12] and, thereby, prolongs the production cycle.

The most prevalent strategy for modulating the crystallization behavior of semi-crystalline polymers involves the incorporation of nucleating agents [14,15,16,17,18]. These agents accelerate crystallization by reducing the nucleation energy barrier and increasing the nucleation density, with the macroscopic effects of an elevated crystallization temperature (Tc), an enhanced crystallization rate, and a refined crystalline morphology. A wide array of nucleating agents have been developed for PET, with inorganic fibers or particulates constituting a major category. Beyond ameliorating the crystallization kinetics, these fillers can also improve the overall material performance. However, their efficacy is strongly contingent upon the filler dispersion state and the interfacial adhesion with the polymer matrix. To date, numerous PET composites containing nucleating fillers have been developed. Commonly used crystallization-promoting fibers include glass fibers, carbon fibers, and natural fibers, while the effective nucleating nanoparticles cover clays, silica, modified graphene oxide, graphene quantum dots, multi-walled carbon nanotubes (MWCNTs), calcium carbonate, silicon carbide, and metal–organic frameworks [15].

A low loading of MWCNTs can markedly enhance the Tc and crystallinity of PET [19,20]. Furthermore, through synergistic interaction with ethylene–butyl acrylate–glycidyl methacrylate copolymer (E-BA-GMA), MWCNTs can also elevate the impact strength of the resulting composites to 72.3 kJ/m^2^ [19]. Although glass fibers can also raise the Tc of PET, their nucleation efficiency is decreased when combined with E-BA-GMA, which primarily serves as a toughening and compatibilizing agent. In contrast, the addition of a maleic anhydride-grafted ethylene–octene copolymer (POE-g-MAH) can further potentiate the crystallization-promoting effect of glass fibers in PET [21].

Among various fillers, halloysite nanotubes (HNTs) have gained extensive attention due to their high aspect ratio, a unique hollow structure, high surface hydroxyl density, abundant sources, and low cost [22]. Researchers [23,24,25] have confirmed that HNTs are efficient nucleating agents for polymers. For instance, the incorporation of 3–5 wt% HNTs can increase the Tc of PET by more than 15 °C [23]. Similar excellent nucleation effects have been reported in polypropylene [24] and bio-based polyesters [25].

However, HNTs suffer from poor compatibility with recycled PET and are prone to agglomeration within the matrix. These issues adversely affect the crystallization behavior and mechanical properties of PET, even leading to increased material brittleness. In the literature [26,27], coupling agents are frequently employed to modify the surface of HNTs, aiming to improve their compatibility with the polymer matrix. Nevertheless, this method introduces additional processing steps and offers only limited enhancement to the impact performance of the final composite.

Despite the well-established efficacy of E-MA-GMA in toughening recycled PET and the recognized role of HNTs in promoting its crystallization, the potential synergistic effects arising from their combined use remain largely unexplored. In particular, a systematic understanding of how HNTs disperse, interact at interfaces, and ultimately regulate material performance within an elastomer-toughened recycled PET multiphase system is lacking. To address this gap, E-MA-GMA was selected as the toughener in this study, primarily due to the reactivity of its epoxy groups with PET terminal groups, which enables effective interfacial compatibilization and toughness enhancement. HNTs, characterized by abundant surface hydroxyl groups, were employed as a nucleating agent with a dual purpose: to leverage their heterogeneous nucleation capability and to exploit potential interactions—between their surface hydroxyls and the functional groups of E-MA-GMA—that may improve HNTs’ dispersion and interfacial adhesion in the recycled PET matrix. The novelty of this work lies in the first comprehensive investigation of the synergistic contributions of E-MA-GMA and HNTs in recycled PET, with a focus on elucidating their combined influence on the material’s rheological, thermal, and mechanical properties—especially impact toughness. By systematically examining the cooperative modification of recycled PET with an elastomer and a nucleating agent, this research aims to provide new strategies and technical insights for developing high-performance recycled PET composites that integrate superior impact resistance with enhanced processability. Such advances are expected to support the broader application of recycled plastics in value-added sectors, thereby mitigating plastic waste pollution, reducing dependence on virgin petroleum-based materials, and promoting the sustainable utilization of finite resources.

2. Materials and Methods

2.1. Materials

The recycled PET (CB-602R) with an intrinsic viscosity of 0.82 dL/g was supplied by Far Eastern New Century Corporation, Taipei, China. E-MA-GMA (AX8900) was purchased from Arkema Group, Colombes, France. It exhibits a melt flow index of 6.46 g/10 min (190 °C, 2.16 kg). The mass fractions of methyl acrylate and glycidyl methacrylate in this copolymer are 22.86% and 7.37%, respectively. The HNTs employed were the XFI50 grade product provided by Jiangsu Xianfeng Nano-Materials Technology Co., Ltd., Xuzhou, China. X-ray photoelectron spectroscopy analysis confirmed a HNTs purity exceeding 95 at%, with dimensions of 1–5 μm in length and 50–300 nm in diameter.

2.2. Melt Processing

PET and HNTs were dried for more than 8 h at 120 °C in a vacuum oven, while E-MA-GMA underwent the same drying time at 30 °C. Recycled PET, E-MA-GMA, and HNTs were manually preblended in a container prior to being fed into a twin-screw extruder (TDS-35B model, Nanjing Nuoda Extrusion Equipment Co., Ltd., Nanjing, China). Extrusion was performed at 260 °C with a screw speed of 200 rpm, which was determined based on preliminary tests and the stabilization of the motor current reading. The melt-blend was considered homogenized once the current reached a steady state, indicating consistent melt viscosity, coupled with visual uniformity. The extruded pellets were dried and then injection-molded into specimens for mechanical testing via a DH-130SeIII injection molding machine (Donghua Machinery Ltd., Dongguan, China). The processing temperatures were 260 °C at the barrel and 25–30 °C at the mold. To investigate the influence of the cooling rate on the mechanical properties of recycled PET and its blends and composites, specimens were prepared using a BS80-III injection molding machine (Borch Machinery Intelligent Equipment Co., Ltd., Guangzhou, China). The mold employed in this process lacked a cooling system, leading to mold temperatures substantially higher than room temperature. Specimens for rheological tests were fabricated at 270 °C via a QLLHY-50T compression molding machine (Xiamen Qunlong Instrument Co., Ltd., Xiamen, China). The formulations for all samples are listed in Table 1. To facilitate an understanding of the sample formulations, the composition and proportion of each sample are reflected in its name. The letters “P”, “H”, and “E” represent recycled PET, HNTs, and E-MA-GMA, respectively. The number following each letter indicates the proportion of the corresponding material in the formulation. For example, “PE20H2” denotes that in the PET composite, the content of E-MA-GMA is 20 wt% and the content of HNTs is 2 phr (parts per hundred of PET/E-MA-GMA blend). For convenience, recycled PET is uniformly abbreviated to PET in the subsequent sections of this article.

2.3. Scanning Electron Microscope Characterization

The morphology of PET-based compositions was characterized using a TESCAN MIRA3 scanning electron microscope (SEM). Transverse cryogenic fracture (perpendicular to melt flow) was performed on injection-molded specimens. Prior to SEM observation, samples were sputter-coated with a thin gold layer. Additionally, impact-fractured surfaces were examined using a Hitachi SU8600 SEM (Hitachi High-Tech Corporation, Tokyo, Japan).

2.4. Dynamic Rheological Measurements

The rheological behavior of PET and its blends and composites was characterized via an Anton Paar MCR302 rheometer (manufactured by Anton Paar, Graz, Austria) with a 25 mm-diameter parallel plate. The specimen was kept for 5 min at 265 °C in a nitrogen atmosphere to remove the residual stress prior to the test. Frequency sweeps were performed in the range of 0.0628–628 rad/s with a given strain amplitude of 1%. Before each test, the parallel plate compressed the sample to a thickness of 1.00 mm.

2.5. Differential Scanning Calorimetry Analysis

The crystallization and melting behaviors of PET and its blends and composites were investigated via a NETZSCH DSC 214 differential scanning calorimeter (DSC) (manufactured by NETZSCH-Gerätebau GmbH, Selb, Germany) under a nitrogen atmosphere. The DSC analysis involved heating the samples from room temperature to 300 °C at 30 °C/min, followed by a 3 min isothermal hold to erase thermal history. Subsequently, they were cooled to 20 °C at 10 °C/min and then reheated to 300 °C at 10 °C/min. Both heating and cooling curves were recorded to assess the influence of E-MA-GMA and HNTs on the melting and crystallization behavior of the PET matrix. The crystallinity (X_c_) of the PET matrix was calculated using the following equation:

$$[eqn]$$

where $[eqn]$ represents the melting enthalpy of the tested sample; $[eqn]$ denotes the standard melting enthalpy of fully crystalline PET (X_c_ = 100%), which was taken as 140 J/g in this experiment [11]; $[eqn]$ is the mass fraction of the PET matrix in the composites.

2.6. Mechanical Properties

Tensile properties of PET and its blends and composites were evaluated according to GB/T 1040.2-2022 [28], using an AI-7000S Universal Testing Machine from Gotech Testing Machines Inc., Dongguan, China. Tensile tests were conducted at crosshead speeds of 50 mm/min and 500 mm/min, respectively. Izod notched impact strength was measured at room temperature using a GT-7045-MDL plastic impact tester (Gotech Testing Machines Inc., Dongguan, China) in accordance with GB 1843-2008 [29]. The pendulum had a speed of 3.46 m/s and a maximum impact energy of 11 J. For each PET-based system, five replicate specimens were randomly selected and evaluated. All test data were averaged and are presented as the mean value ± standard deviation.

2.7. Wide-Angle X-Ray Diffraction Measurements

Wide-angle X-ray diffraction (WAXD) measurements of PET, its blends and composites were conducted on a Rigaku D/max-2200vpc diffractometer (Rigaku Corporation, Tokyo, Japan) with Cu Kα radiation. Measurements were performed at 40 kV and 30 mA, recording diffraction patterns at a scanning rate of 0.02 °/min in the 2θ range of 5–70°. The samples were categorized into two groups: as-molded without post-treatment, and annealed in a vacuum oven at 140 °C for 160 min (where the time started from the onset of heating and represented the total thermal cycle). Raw data were processed and analyzed using Origin software (OriginPro 2025). Specifically, the crystallinity (Xc) [30] was calculated through the Lorentzian peak deconvolution method, following the equation:

$$[eqn]$$

Here, $[eqn]$ represents the crystalline peak area, and $[eqn]$ denotes the amorphous halo area.

3. Results and Discussion

3.1. Distribution and Dispersion of HNTs and E-MA-GMA in PET

The distribution and dispersion of HNTs and E-MA-GMA within the PET matrix are critical to the properties of the resulting composites. Figure 1 shows SEM images of the cryogenically fractured surfaces of the PE20H1, PE20H2, and PE20H4 composites. To distinguish E-MA-GMA from the PET, the fractured surfaces were etched by refluxing in boiling toluene, which selectively removed the E-MA-GMA. Conversely, a phenol/tetrachloroethane mixture in a 1:1 volume ratio was employed to selectively remove the PET. As shown in Figure 1a, PE20H1 exhibits a characteristic “sea–island” morphology, with E-MA-GMA homogeneously dispersed as submicron-scale domains within the PET matrix. Notably, HNTs are selectively localized at the PET/E-MA-GMA interface (the region circled in white in Figure 1a–c), where they remain well-dispersed as individual fibers and establish interfacial connections with the PET matrix. This interface–localization significantly improved the interfacial compatibility between PET and E-MA-GMA, which is consistent with the previously reported findings [19]. When the HNT content increases to 2 phr (Figure 1b), most HNTs remain individually dispersed at the interface, while a small amount is observed in the PET matrix (marked by a white rectangle in Figure 1b). A further increase to 4 phr (Figure 1c) leads to significant aggregation, forming ~30 μm clusters within the PET matrix, though some individual HNTs are still dispersed at the interface. Figure 1d presents the SEM image of PE20H4 after etching with a phenol/tetrachloroethane mixture in a 1:1 volume ratio. The complete removal of the PET phase reveals spherical E-MA-GMA domains, with no HNTs observed within them. This confirms that under the experimental conditions, HNTs preferentially localize at the PET/E-MA-GMA interface. When there are excess HNTs, they exist either as individual fibers or as aggregates within the PET matrix.

To schematically illustrate the distribution of HNTs in the E-MA-GMA-toughened PET system, a phase–structure diagram (Figure 2) was prepared based on the SEM results shown in Figure 1. The E-MA-GMA molecular chain comprises both polar GMA units and non-polar ethylene units. The epoxy groups in E-MA-GMA undergo ring-opening addition reactions with the terminal groups (carboxyl/hydroxyl) of PET or the surface hydroxyl groups of HNTs, leading to the formation of graft copolymers such as PET-g-(E-MA-GMA) or HNTs-g-(E-MA-GMA) at the interface [12]. These grafted structures effectively enhance the dispersion of the dispersed phase (E-MA-GMA or HNTs) within the PET matrix and improve the interfacial strength between the two phases. This structural feature allows E-MA-GMA to form a core–shell structure within the PET matrix: the ethylene units constitute a hydrophobic core, while the GMA units form a hydrophilic shell that interacts with the PET. Due to the strong interaction between HNTs and GMA units, HNTs preferentially localize in the interfacial shell region between PET and E-MA-GMA.

3.2. Effect of HNTs and E-MA-GMA on PET Chain Relaxation

Rheological testing serves as an effective method for characterizing the topological structure of polymers, the dispersion state of fillers, and the relaxation behavior of macromolecular chains. In this study, the oscillatory shear mode was employed to investigate the effects of HNTs and/or E-MA-GMA on the viscoelastic behavior of PET and its blends/composites. Figure 3a–c presents the frequency-dependent curves of the storage modulus (G′), loss modulus (G″), and complex viscosity (η*) for these materials. The frequency sweep curves of G′ and G″ reveal a solid-like plateau at low frequencies, where G′ or G″ remains nearly constant with the increasing frequency. This behavior is a characteristic of the formation of substantial long-chain branching structures in polymer systems [31] or the development of physical networks of nanofillers in composites [20,32]. Typically, the long-chain branching structures or the nanofiller networks can substantially enhance the G′, G″, and η* of the sample melt [20,31,32]. As shown in Figure 3a,b, the G′ and G″ curves of PE20 exhibit a distinct plateau in the low-frequency region. This phenomenon is attributed to the reaction between the epoxy groups in E-MA-GMA and the terminal groups of the PET, which form long-chain branching structures. These structures increase the entanglement density and enhance intermolecular interactions, leading to a substantial increase in G′, G″, and η* compared with PET. For PE20H2 and PE20H4, the epoxy groups in E-MA-GMA react not only with PET terminal groups but also with the surface hydroxyl groups of HNTs. This dual interaction forms a composite system featuring both long-chain branching structures and a hybrid network, resulting in an even higher G′, G″, and η* than in PE20. In contrast, PH2 (containing only 2 phr HNTs without E-MA-GMA) exhibits a lower G″ and η* than PET. This phenomenon is attributed to the absence of chemical interactions between HNTs and PET without E-MA-GMA, preventing the formation of a stable hybrid network. Furthermore, the solid HNTs act as physical barriers, reducing PET intermolecular entanglement and promoting PET chain slippage, thereby lowering the G″ and η* below those of the PET. A similar mechanism explains why PE20H4 exhibits lower G′, G″, and η* than PE20H2, particularly at low frequency. Excessive HNTs tend to agglomerate (Figure 1c), which compromises the effectiveness of the hybrid network compared with the well-dispersed state in PE20H2.

The van Gurp–Palmen (vGP) and Cole–Cole plots allow a more detailed correlation between the dynamic rheological properties of blend/composite materials and the evolution of their microstructure and component compatibility, providing crucial insights into the polymer microstructure and molecular chain relaxation mechanisms [33]. The vGP plot depicts the phase angle (δ) versus the complex modulus (G*). As noted in reference [34], a smaller phase angle (δ) in the low G* region indicates a higher degree of long-chain branching in polymers. For polymer-based composites, this characteristic suggests the formation of a network-like structure [33,35]. Both long-chain branching and network structures can impede the relaxation of molecular chains. As shown in Figure 3d, the phase angle of PET remains at a 90° plateau across the entire range of complex moduli, exhibiting the typical characteristics of linear polymers. The vGP curve of PH2 closely follows and nearly overlaps with that of the PET. This indicates that the addition of HNTs alone does not significantly affect the relaxation of PET molecular chains. In contrast, the vGP curve of PE20 deviates from that of the PET and shifts toward lower phase angle regions, suggesting that E-MA-GMA forms long-chain branched structures with PET, thereby slowing down the chain relaxation. The vGP curves of PE20H2 and PE20H4 are located to the lower right of PE20, with smaller phase angles at the same complex modulus. This indicates that the synergistic effect of E-MA-GMA and HNTs further slowed down the relaxation of PET chains and promoted quasi-solid behavior in the system.

The synergistic effect of E-MA-GMA and HNTs in retarding PET chain relaxation is further supported by the Cole–Cole plot in Figure 3e. This plot displays the imaginary part (η″) against the real part (η′) of the complex viscosity. A single semicircle is typically observed for homogeneous or fully compatible systems. Deviation from a semicircular shape indicates system heterogeneity or poor compatibility, and an increase in the semicircle diameter corresponds to a longer average relaxation time [33]. The Cole–Cole curves of PET and PH2 exhibit irregular shapes (Figure 3e), which can be attributed to their relatively lower viscosity. The curves for PE20, PE20H2, and PE20H4 deviate from the ideal semicircle and display a noticeable upward trend in the low-frequency region. This phenomenon arises from the synergistic effect of long-chain branching and the hybrid-network structure, which introduces longer relaxation times at low frequencies, and thus modifies the rheological response [31,33,34]. The semicircle diameter is larger for PE20H2 and PE20H4 than for PE20, indicating that the combined E-MA-GMA and HNT system hinders PET chain mobility more strongly than E-MA-GMA alone, leading to increased relaxation times. A closer inspection shows that the arc radius of PE20H4 falls between those of PE20 and PE20H2, suggesting that its relaxation time is longer than that of PE20 but shorter than that of PE20H2. When the HNT content was increased to 4 phr, it tended to agglomerate into larger-sized particles (Figure 1c). These large particles are less effective in contributing to the hybrid-network structure compared with the well-dispersed 2 phr HNTs system, which results in a shorter relaxation time for PE20H4 relative to PE20H2. These results confirm that an efficient dispersion of HNTs is essential for effectively retarding the relaxation of PET chains.

The analysis of vGP and Cole–Cole plots reveals that the addition of HNTs alone does not significantly alter the PET chain relaxation. In contrast, incorporating E-MA-GMA, either alone or in combination with HNTs, substantially slows the relaxation process. Notably, the synergistic effect of E-MA-GMA and HNTs leads to a more pronounced hindrance of PET chain mobility, as evidenced by the extended relaxation times.

3.3. Effect of HNTs and E-MA-GMA on PET Crystallization Behavior

Figure 4 presents the DSC curves of PET, PH2, PE20, and PE20H2, with the corresponding crystallization and melting parameters summarized in Table 2. As shown in Figure 4a and Table 2, compared with PET, PH2 containing 2 phr HNTs exhibits an 11.2 °C increase in crystallization temperature from melt (Tc). However, PE20 containing 20 wt% E-MA-GMA shows an increase of only 2.1 °C. This indicates that HNTs serve as effective heterogeneous nucleating agents for PET, promoting the initiation of PET chain crystallization at higher temperatures, whereas E-MA-GMA exhibits a much weaker nucleating effect. PE20H2 shows a Tc of 193.4 °C, merely 3.4 °C higher than that of PET. This suggests that the nucleating effect of HNTs on PET is markedly diminished when HNTs are added together with E-MA-GMA. As reported in the literature [36], the epitaxial growth of polymer chains on nucleating agents is a critical step in heterogeneous nucleation crystallization. This process refers to the growth of one phase on the surface of another with a specific crystallographic orientation, which depends on the lattice matching between the two phases. In the PE20H2 system, the hydroxyl groups on the HNTs surface readily undergo chemical reactions with the epoxy groups in the E-MA-GMA molecular chains, causing the HNTs to preferentially disperse within the GMA segments of E-MA-GMA (Figure 1). The encapsulation of HNTs by E-MA-GMA is the primary reason for the reduced nucleating efficiency.

In polymer crystallization studies, Tc and the full width at half maximum height of the crystallization peak (FWHMH, Figure 4c) are key indicators of crystallization kinetics. A higher Tc together with a narrower FWHMH corresponds to a faster crystallization rate [37,38]. Table 2 shows that the FWHMH of PH2 is 55.8 s, which is 11% narrower than that of PET. This suggests that HNTs effectively enhance the crystallization rate of PET, further confirming their role as an efficient nucleating agent. However, PE20 exhibits a significantly increased FWHMH of 77.3 s, which is 23% wider than that of PET. This is caused by the reaction between the epoxy groups in E-MA-GMA and the terminal groups of PET. The resulting long-chain branched structures restrict the mobility of PET chains and, thereby, reduce the crystallization rate. Similar phenomena were observed by Loyens et al. [11] in rubber-toughened PET systems. Martuscelli et al. [39] also confirmed that reactive compatibilization increases the system viscosity, consequently decreasing the crystallization rate. In PE20H2, HNTs and E-MA-GMA exert opposing effects on the crystallization rate of PET. Consequently, the FWHMH of PE20H2 is 59.8 s, only 5% narrower than that of the PET. This result demonstrates that even with the incorporation of E-MA-GMA, the HNTs still retain a certain degree of nucleating capability for PET, although their overall crystallization-promoting effect is significantly reduced.

Figure 4b presents the DSC remelting curves of PET, PH2, PE20, and PE20H2, where the endothermic peak in the range of 225~260 °C corresponds to the melting peak of PET. Compared with PET, the melting points (Tm) of PH2, PE20, and PE20H2 show minimal variation (ranging from 249.0 to 251.9 °C, as shown in Table 2), and the crystallinity (X_c_) calculated from the melting enthalpy (ΔH_m_) also changes insignificantly (28.3–29.1%, Table 2). These results suggest that HNTs and/or E-MA-GMA exert negligible effects on the ultimate crystal perfection and the maximum crystallizable fraction of PET.

The effect of HNT content on the crystallization and melting behavior of the PET/E-MA-GMA system is summarized in Figure 5 and Table 3. As the HNT content increases from 1 to 4 phr, its influence on the PET crystallization temperature and crystallization rate is minimal and far less pronounced than the effect observed in PH2. This is because the E-MA-GMA coats the surface of HNTs, thereby reducing its nucleating effect on PET.

3.4. Effect of HNTs and E-MA-GMA on the Tensile Properties of PET

Figure 6a shows the elongation at break of PET, PH2, PE20, and PE20H2 obtained at crosshead speeds of 50 and 500 mm/min. At the slower speed of 50 mm/min, the elongation at break of PET exceeds 700%, whereas it drops below 200% at the higher speed of 500 mm/min. The samples were injection-molded at a relatively low mold temperature (25–30 °C), resulting in an amorphous state of PET. Under slow deformation, the polymer chains have adequate time to disentangle from their initial disordered entanglements and orient along the stretching direction. Moreover, PET—being a semi-crystalline polymer—can undergo stress-induced crystallization during stretching, which further dissipates stress and extends the deformation process prior to fracture. In contrast, at the crosshead speed of 500 mm/min, the duration of applied stress is shorter than the relaxation time of the PET chains, leading to stress concentration and a pronounced decrease in elongation at break.

For PH2, the elongation at break is lower than that of PET under both crosshead speeds, which is ascribed to the poor compatibility between HNTs and PET. This indicates that the addition of HNTs alone does not toughen PET. The relatively large error bars for PET and PH2 stem primarily from the inherent structural heterogeneity of recycled PET and the poor PET/HNTs compatibility. Unlike PET and PH2, the crosshead speed exerts a relatively minor influence on the elongation at break of PE20 and the PET composites containing E-MA-GMA. This behavior originates from the short relaxation time of the flexible chain segments in E-MA-GMA, enabling a prompt response to external stress under the tested conditions. Through their own deformation, these segments effectively dissipate stress, thereby enhancing the elongation at break. The results demonstrate that E-MA-GMA can effectively toughen PET.

Notably, when the HNT content increases from 0 to 2 phr, the elongation at break measured at 50 mm/min rises from 468% (PE20) to 580% (PE20H2), corresponding to an increase of 23%. Even at 500 mm/min, PE20H2 still exhibits about 20% higher elongation at break compared with PE20. These findings suggest a synergistic effect between HNTs and E-MA-GMA in enhancing the toughness of PET. However, further increasing the HNT content to 4 phr leads to a sharp decline in the elongation at break of PE20H4. This is mainly attributed to the poor dispersion of the excessive HNTs within the PET matrix, which form aggregates (as illustrated in Figure 1c) and consequently deteriorate the material performance.

Figure 6b shows the yield stress and fracture stress of the PET and PET-based blends/composites at different crosshead speeds. A consistent trend is observed for all samples: the yield stress values obtained at 500 mm/min are higher than those at 50 mm/min. Notably, PH2 exhibits a higher yield stress than PET under both crosshead speeds, which is primarily attributed to the high intrinsic strength and stiffness of HNTs [40], confirming their reinforcing function in the matrix. However, compared with earlier reports on carbon nanotubes (CNTs) [41], the reinforcement efficiency of HNTs appears less pronounced. For instance, at 500 mm/min, the yield stress of PH2 is only about 6% higher than that of PET. This lower efficiency can be explained by the dimensional difference between HNTs (diameter 50~300 nm) and CNTs (diameter 17.36 nm); the larger aspect ratio and smaller size of CNTs generally favor more effective load transfer. The addition of 20 wt% E-MA-GMA—a conventional toughening agent with inherently low yield stress—leads to a considerable decrease in yield stress for PE20, PE20H1, PE20H2, and PE20H4 relative to PET. Moreover, within the experimental range studied, variations in HNT content show only a minor influence on the yield stress of the PET/E-MA-GMA system.

Interestingly, the fracture stress values of all samples measured at 500 mm/min are lower than those obtained at 50 mm/min. This trend can be explained by the different molecular-chain dynamics under varied strain rates. At the lower crosshead speed (50 mm/min), there is sufficient time for PET chains to align along the stretching direction and to undergo thorough reorganization and recrystallization. In contrast, during high-speed deformation (500 mm/min), the relaxation of the polymer chains lags far behind the applied strain rate; hence, the chains cannot reach full orientation before fracture. Because crystalline regions generally possess higher mechanical strength than amorphous domains, PET and PH2 exhibit markedly higher fracture stress at 50 mm/min, owing to the development of a highly oriented structure. The incorporation of E-MA-GMA into the PET system introduces flexible chain segments with short relaxation times, which respond promptly to external loading and alleviate stress concentration through their own deformation, thereby markedly increasing the elongation at break of the specimens. In this process, the PET molecular chains gain sufficient time to undergo orientation and crystallization, leading to an enhancement in the fracture stress of PET-based blends/composites. When the HNT content increases from 0 to 2 phr, the fracture stress of PET-based composites changes only slightly. However, a further increase in HNT content to 4 phr results in a pronounced decrease in fracture stress, which is attributed to the agglomeration of excess HNTs within the matrix (as shown in Figure 1c).

3.5. Effect of HNTs and E-MA-GMA on the Impact Strength of PET

Figure 7 shows the impact strength of the PET and PET-based blends and composites before and after annealing. Here, “unannealed” samples refer to samples kept at room temperature for at least 48 h after injection molding; “annealed” samples were annealed at 140 °C for 160 min in a vacuum oven, with the time starting from the onset of heating and representing the total thermal cycle. The impact strength of unannealed PH2 is lower than that of PET, indicating that HNTs alone do not effectively toughen PET. In contrast, the unannealed PE20 sample exhibits an impact strength of 77.2 kJ/m^2^, which is about 13.9 times that of PET, confirming E-MA-GMA as an efficient toughening agent for PET. Figure 1 shows that E-MA-GMA is dispersed uniformly in the PET matrix in nearly spherical droplets. This “discrete dispersed phase–uniform distribution” microstructure provides the structural basis for the improved toughness of the polymer matrix [9]. The impact fracture surface of PE20 (Figure 8a) displays a rough undulating topography with highly stretched fibrous bundles of the PET matrix, together with voids formed by the deformation of E-MA-GMA droplets under stress. The energy dissipation associated with matrix deformation and void formation constitutes the key mechanism through which E-MA-GMA toughens PET [9]. When HNTs are incorporated into the PET/E-MA-GMA system, the impact strength remains above 76.0 kJ/m^2^, as long as the HNT content does not exceed 2 phr. This impact strength value is comparable with that reported by Huang et al. [19] for MWNTs-filled PET composites. However, when the loading is increased to 4 phr, the impact strength of PE20H4 drops markedly to 68.7 kJ/m^2^. This decline is directly attributed to the excessive aggregation of HNTs within the matrix (Figure 1c). It should be noted that the actual impact strengths of PE20, PE20H1, and PE20H2 are significantly higher than the values recorded in Figure 7, as the specimens did not undergo complete fracture during testing. This phenomenon has also been reported in the work of Huang et al. [19] and Sun et al. [20].

Upon annealing at 140 °C for 160 min, the impact strength of PET and its blends and composites exhibited a significant decrease compared with that of the unannealed samples. In particular, the four samples containing E-MA-GMA showed a pronounced reduction in impact resistance, with the decline ranging from 45% to 74%. To clarify the structural origin of this mechanical deterioration, WAXD analysis was performed on both the annealed and unannealed specimens (Figure 9). The distinct sharp peaks, observed at 2θ = 12.0° and 24.5°, correspond to the diffraction from the (001) and (002) crystal planes of HNTs [40], respectively. In the XRD patterns of pristine HNT powder and PET/HNTs composites with different HNT loadings (PH2, PE20H1, PE20H2, and PE20H4), the characteristic diffraction peaks corresponding to the (001) and (002) crystal planes of HNTs exhibit no discernible shift in 2θ position and no change in peak width. These results indicate that the HNTs are dispersed as intact tubular particles in the PET matrix, without exfoliation into individual nanosheets or intercalation of PET chains into the interlayer spacing of HNTs. This is consistent with the structural features observed in the SEM images of Figure 1a–c. In the unannealed samples (Figure 9a), two broad diffraction peaks appear within the 2θ range of 10–60°, which are attributed to the primary and secondary amorphous halos of PET [30]. This indicates that the PET matrix remains predominantly amorphous, a consequence of rapid quenching during injection molding at low mold temperatures (25–30 °C), which effectively freezes the polymer chains before substantial crystallization can take place.

In contrast, annealing induced notable alterations in the crystalline morphology. Figure 9b shows that sharp diffraction peaks superimposed on the primary amorphous halo were detected at approximately 2θ = 16–17°, 21–22°, and 26°, corresponding respectively to the (100), (−110), and (010) crystal planes of PET [42]. These findings demonstrate that the annealing process promoted recrystallization of the PET matrix. At 140 °C, the enhanced segmental mobility of PET chains facilitates conformational rearrangement, enabling previously frozen non-crystalline regions to undergo full crystallization. Moreover, imperfect crystalline domains were reorganized into more regular and well-defined crystalline structures. Consequently, the annealing treatment led to a significant increase in the crystallinity of the PET-based materials and improved the structural integrity of the crystalline phases [43]. The WAXD curves presented in Figure 9b were fitted to a Lorentzian function using the peak analysis module in Origin software. The deconvolution results for the annealed PET samples are shown in Figure 10; analogous fitting procedures were applied to the remaining samples, although their detailed spectra are omitted here for brevity. The crystallinity was calculated based on the integrated intensities of the six most intense crystalline peaks within the 2θ range of 10–35°, in accordance with the method reported by Bai et al. [30], using the formula given in Equation (2). Figure 11 shows that the crystallinity of the annealed PE20, PE20H1, PE20H2, and PE20H4 increased markedly from an initially near zero level (predominantly amorphous in Figure 9a) to approximately 36–43% after annealing. To further examine the impact of annealing on the crystallization behavior of the PET-based blends and composites, DSC was conducted on both unannealed (as-injection-molded) and annealed specimens. The DSC results (Figure S1c) reveal a pronounced increase in crystallinity following annealing, which is consistent with the trend reflected by WAXD shown in Figure 9 and Figure 11.

It is well established that the degradation of PET leads to increased crystallinity and reduced impact resistance. In this study, however, the rise in crystallinity observed in PET-based blends and composites is attributed solely to the annealing treatment rather than to PET degradation. This conclusion is supported by the data presented in Figures S2 and S3 of the Supplementary Materials.

In comparison with the loosely entangled amorphous structure, the highly ordered and densely packed molecular chains in the crystalline regions result in reduced impact toughness. Thus, the pronounced rise in crystallinity observed for these four samples (PE20, PE20H1, PE20H2, and PE20H4) is attributed to the significant decline in impact strength following annealing.

Notably, the initial impact strengths of PE20, PE20H1, and PE20H2 were comparable prior to annealing. A distinct differentiation was observed following the annealing process. Among these samples, PE20H1 exhibited the highest impact strength (42.4 kJ/m^2^), with a reduction rate of 44.6% relative to its pre-annealed state. In contrast, PE20 and PE20H2 showed significantly higher reduction rates of 74.3% and 67.1%, respectively. After annealing treatment, the impact strength of PE20H1 was 2.14 times that of PE20. Therefore, PE20 (showing the smallest impact strength) and PE20H1 (showing the largest impact strength) were selected as representative specimens for the SEM examination of their impact fracture surfaces.

Figure 8 shows that the unannealed PE20 (Figure 8a) and PE20H1 (Figure 8b) exhibit ductile fracture characteristics, with the PET matrix undergoing extensive deformation. Such features are absent in the annealed PE20 near the crack tip (Figure 12a), where the fracture surface is relatively smooth, lacking pronounced deformation and displaying typical brittle morphology. PET deformation becomes observable only farther from the crack tip (Figure 12c). In contrast, the annealed PE20H1 demonstrates ductile fracture both near to (Figure 12b) and far from (Figure 12d) the crack tip, although the extent of PET deformation is reduced compared with the unannealed state (Figure 8b). This disparity can be attributed to the lower crystallinity of PE20H1 (37.7%) after annealing. As indicated in Figure 3, the combined presence of HNTs and E-MA-GMA restricts PET chain relaxation more effectively than E-MA-GMA alone. During annealing at 140 °C for 160 min, the rearrangement of PET chains in PE20H1 proceeds more slowly than in PE20, resulting in a lower final crystallinity for PE20H1. Upon impact loading, the soft segments of E-MA-GMA can undergo motion and deformation, thereby inducing PET matrix deformation and dissipating the impact energy. It is well known that higher crystallinity in a given polymer reduces its capacity for deformation; thus, the higher crystallinity of annealed PE20 leads to a lower impact strength compared with PE20H1. The data in Figure 11 and Figure S1, in conjunction with these observations, demonstrate that when elastomers are employed to toughen PET, the toughening efficiency depends not only on the elastomer content, dispersion state, and interfacial adhesion [30] but also critically on PET crystallinity—higher crystallinity diminishes the toughening effect.

3.6. The Practical Applicability of the PET/E-MA-GMA/HNTs Composites

In industrial practice, injection-molded PET products are frequently subjected to post-molding heat treatments to relieve residual stresses and ensure dimensional stability and service performance; however, such thermal processing inevitably promotes crystallization, leading to a pronounced deterioration in impact strength. The present study demonstrates that the incorporation of a modest amount of HNTs can effectively mitigate this compromise, providing a viable strategy for manufacturing PET products that require post-molding heat treatment while maintaining high impact resistance.

The synergistic effect between E-MA-GMA and HNTs significantly restricts the mobility of PET chains, suppresses excessive crystallinity development, and thereby preserves high impact toughness. Based on this finding, the present work further evaluated the practical applicability of the PET/E-MA-GMA/HNTs composites in the process of injection molding. In conventional industrial practice, obtaining high-impact-strength PET components typically requires rapid quenching of the melt by maintaining the mold temperature below 25 °C. It is hypothesized that, in a PET/E-MA-GMA system containing 20 wt% E-MA-GMA, the addition of 1–2 phr HNTs may eliminate the need for such abrupt cooling. To verify this hypothesis, specimens were prepared using a mold without active cooling channels. Two formulations were processed: PE20 (20 wt% E-MA-GMA, without HNTs) and PE20H2 (20 wt% E-MA-GMA + 2 phr HNTs). The resulting samples were designated as PE20-S and PE20H2-S, respectively. Impact tests were conducted on five randomly selected specimens, and the average values are presented in Figure 13. The impact strength of PE20H2-S was 50.0 kJ/m^2^, which was 2.28 times that of PE20-S. The WAXD analysis (Figure 14) reveals that the PE20H2-S sample exhibited a crystallinity of 26%, which was significantly lower than the 40% observed for PE20-S. Examination of the fracture surfaces by SEM (Figure 15) revealed distinct failure modes. The fracture surface of PE20-S was smooth, showing minimal deformation of the PET matrix and characteristics of brittle fracture. In contrast, the fracture surface of PE20H2-S was rough, with the PET matrix undergoing pronounced plastic deformation, indicative of a ductile tough fracture. These results confirm that incorporating 2 phr HNTs into the E-MA-GMA-toughened PET system allows the omission of the rapid-quench step while effectively restraining the crystallinity increase. Thus, high-toughness PET components can be produced via a simplified processing route.

The PET/E-MA-GMA/HNTs composite system developed in this study exhibits significant potential for industrial applications. Its key advantage lies in the synergistic effect between HNTs and E-MA-GMA, which enables the fabrication of products with low crystallinity and high toughness without relying on rapid quenching. This provides a novel and industrially feasible strategy to address the brittleness of PET—especially recycled PET—caused by increased crystallinity during processing. For practical applications, this material is suitable for fields demanding high toughness and thermal resistance, including non-structural automotive interior parts, electrical and electronic enclosures, and high-performance packaging containers. Compared with conventional processes for fabricating high-toughness PET (rapid quenching), the proposed composite allows injection molding at higher mold temperatures, which effectively reduces the internal stress, improves the dimensional stability, and eliminates the need for dedicated cold molds or quenching equipment. Consequently, the processing flow is greatly simplified, with reduced energy consumption and production costs. Furthermore, the synergistic interaction between E-MA-GMA and HNTs effectively elevates the upper performance limit of recycled PET, facilitating its upgrading from low-value recycling to high-value engineering applications. As a fully thermoplastic system, the composite retains the essential characteristic of melt-processability, which is fundamental to polymer recycling. Therefore, it can be mechanically recycled through conventional steps such as grinding, re-melting, and re-processing, similar to virgin PET. This aligns with the strategic goals of sustainable development and resource circularity. Future work may focus on optimizing formulation ratios for industrial production, evaluating the long-term thermal stability and multi-cycle recyclability of the material to accelerate its industrial scale-up.

4. Conclusions

This study employed a melt blending method to synergistically modify recycled PET with HNTs and E-MA-GMA, focusing on the synergistic effect of HNTs and E-MA-GMA on the properties of PET. The SEM results indicated that HNTs were mainly located at the E-MA-GMA/PET interface. At 1 phr, they were individually dispersed and connected to PET. Increasing the content led to some individual dispersion within the PET matrix, but interfacial dispersion remained dominant. Agglomeration became evident at 4 phr within the matrix. Rheological analysis indicated that the combined addition of E-MA-GMA and HNTs imposed a significantly stronger restriction on the relaxation of PET molecular chains than the addition of E-MA-GMA alone, thereby notably reducing the crystallization rate of PET during processing. The DSC results demonstrated that HNTs can act as a heterogeneous nucleating agent for PET, effectively enhancing its crystallization temperature and rate. In contrast, E-MA-GMA exhibited a negligible nucleating effect. Furthermore, its reaction with PET significantly reduced the crystallization rate of PET. When HNTs and E-MA-GMA were incorporated together, HNTs still functioned as a nucleating agent for PET, but their nucleation efficiency was significantly compromised. The impact performance of unannealed PET blends and composites containing E-MA-GMA exceeded 68 kJ/m^2^ due to their low crystallinity, resulting from the low injection molding temperature (25–30 °C). Annealing significantly increased the crystallinity, leading to a drastic drop in the impact strength. This confirms that in elastomer-toughened PET systems, higher crystallinity severely reduces the impact resistance. The synergy between HNTs and E-MA-GMA restricted the chain mobility more effectively than E-MA-GMA alone, resulting in a lower crystallinity and higher impact strength in the annealed PE20H1 compared with the PE20. Consequently, HNTs-containing composites maintained low crystallinity even at higher mold temperatures, achieving both high toughness and easy processability. The developed PET/E-MA-GMA/HNTs composites demonstrate potential for replacing virgin plastics with recycled PET in demanding applications such as automotive and electronics, offering environmental and resource benefits.