Dual-shape memory effect in radiation crosslinked thermoplastic blends: fabrication, optimization and mechanisms

Abstract

Recently, as an important class of mechanically active smart materials, thermoplastic dual-shape memory polymers (SMPs) have attracted notable attention and can be fabricated in many different manufacturing techniques. Here in this paper, we present experimental results, demonstrating a cost-effective manufacturing technique to enable thermoplastic SMPs with enhanced properties for a wide variety of applications. Thermoplastic SMPs based on low density polyethylene (LDPE) and polypropylene (PP) with various compositions were prepared by melt compounding, followed by a post-processing of e-beam irradiation at 5, 10, 15, 25, 50 and 100 kGy. SEM, DSC, tensile test, rheological properties measurement, sol/gel analysis and shape memory test were performed sequentially to investigate the relationship between the phase morphologies, the content fluctuations, the e-beam irradiation and the shape memory performances. In addition, we also optimized the fabrication process and studied mechanisms of shape memory performances. It was found that the melting point associated to the LDPE soft phase T_m,LDPE is almost independent of the content fluctuations. At the same time the mechanical properties (determined at 25 °C) and rheological properties (measured at 180 °C) can be varied systematically by controlling the structure and radiation dose. More importantly, the results also revealed that (1) irradiated LDPE-rich blends were more suitable and effective than both irradiated PP-rich and non-irradiated blends to be dual-SMPs with advantageous shape memory properties, and (2) increased radiation dose could give rise to enhanced shape recovery capacity without significantly weakening the shape fixity.

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1. Introduction

As an expanding field of high technological interest, stimuli-responsive materials currently attract a significant amount of research efforts from scientists and engineers all around the world.^1,2 Possessing advantageous properties over shape memory alloys (SMAs) and shape memory ceramics (SMCs), shape memory polymers (SMPs) are a well-known, representative class of stimuli-responsive materials and they can be programmed to “fix” a temporary shape and then recover to a “remembered” permanent shape upon exposure to an external stimulus.^3–7 Essentially, the associated behavior of a SMP is called a shape memory effect (SME). According to the nature of the molecular structure and working mechanism, it can be classified into dual-SME (one permanent shape and more than one temporary shape), triple- and multi-SME (one permanent shape and more than two temporary shape). However, due to the more excellent feasibility and reliability for extensive industrial applications, dual-SME is gradually attracting much more attention and increasingly becoming an important concept in the design of polymeric materials and composites.^8–12

Typically, these classical simple SMEs have one major drawback of missing reversibility during the shape change (“one-way SME”), and this often hampers the transfer of SMPs into commercial products. By contrast, development of polymers featuring suitable and fascinating “two-way SME” may fulfill the requirements of applications such as sensors and actuators.¹³ During the past two decades, to explore the underlying mechanism and structure of SMPs, several types of structures and models have been established to illustrate the SMEs.^14–16 Thus, from a molecular structure standpoint, we describe that dual-SMPs having a two-way shape memory capability consist of a reversible “soft phase” and a “hard phase”. The hard phase determines the permanent shape and can be made of either chemical or physical crosslinks, with an interpenetrated or interlocked supramolecular complex. Similarly, the soft phase is responsible for controlling the shape fixity e.g. by crystallization or vitrification.^17,18

Today, thermoplastic materials are a very important material class with mass production and a substantial cost advantage, which allows one to find a suitable polymer for nearly every application.^19,20 However, most traditional thermoplastic polymers do not inherently display shape memory behavior. To exhibit shape memory behavior in thermoplastic materials, three major aspects have been taken into consideration when incorporating shape memory functionality to polymer systems: (1) morphologies, (2) structures, (3) processing conditions. Usually, the former two have received much more attentions. For example, Jiang Du et al.²¹ reported that the difference in morphology caused the difference in shape memory behavior observed between layered and blend PU/PCL films with domain sizes on the micro-scale. Shannon R. Armstrong et al.²² similarly discovered that the shape memory behavior can be significantly altered by a change in the phase separated morphologies that possess domain sizes on either the nano- or micro-scale. Except concerning the morphologies, Heng Zhang et al.²³ found the structures of styrene–butadiene–styrene tri-block copolymer (SBS) and poly (3-caprolactone) (PCL) blend had an influence on its shape memory behavior. Moreover, these two factors were also studied by Jun Zhao et al.²⁴ and Cedric Samuel et al.²⁵ While for the processing conditions, nowadays the researches were just focused on manufacturing mode like chemical synthesizing,^26–28 melt co-extrusion blending²⁹ and solution co-precipitation blending.³⁰ Besides, there were precious few works appearing in the papers involving subsequently radiation crosslinking thermoplastics after polymerization, blending and remolding.

Through the analysis of shape memory mechanism, it widely accepted that the driving force for shape recovery in SMPs, occurring on heating the material to a temperature above the switching temperature, is the entropic elasticity of the hard phase which refers to networks such as covalent crosslinks, chain entanglements, or rigid crystalline phase.³¹ Thus, how to obtain a desirable network become extremely necessary and important to create a polymer that exhibits a shape memory effect. Resembling the chemical crosslinking, targeted irradiation of polymeric materials can lead to the creation of a network polymer³² and has been used to achieve the shape memory effect.^33–35 But likewise, this behavior always depends greatly on their chemical and physical structure due to manufacturing limitations.^36,37 The goal of this work is to develop a cost-effective manufacturing technique to enable SMPs with superior mechanical and dual-shape memory properties for a wide variety of applications.

In this study, we fabricated a SMP material by e-beam radiation crosslinking the thermoplastic polymer blends which were composed of low density polyethylene (LDPE) and polypropylene (PP). With respect to the dual-SME, we chose transition temperature T_trans to be within the range of T_m,LDPE to T_m,PP, while we wanted the elastic properties to be adjustable in a wide range by variation of the blend composition and the absorbed radiation dose. In this way, we intended to design an optimized thermoplastic SMP by controlling thermoplastic components and radiation doses, followed a corresponding shape memory mechanism and a detail assessment of the shape memory behavior for this type of SMP system.

2. Experimental

2.1 Materials and sample preparation

The materials used for this study were thermoplastic low-density polyethylene (LDPE) and polypropylene (PP). Both LDPE (2426H) with the density of 0.923 g cm⁻³ and PP (K7726) with the density of 0.900 g cm⁻³ were purchased from Sinopec Beijing Yanshan Company (China), having a reported melt flow index (MFI) of 1.9 g per 10 min and 30 g per 10 min, respectively. Xylene (AR) was purchased from Beijing Chemical Plant (China). Prior to melt-blending, the PP and LDPE pellets were completely dried overnight at 80 °C under vacuum to minimize water content in each component.

Binary blends comprising PP and LDPE were prepared by a Haake batch intensive mixer (Haake Rheomix 600, Karlsruhe, Germany). The melt compounding was performed at 180 °C and a screw speed of 60 rpm for 12 min. Different LDPE/PP compositions were studied (neat LDPE, 90% LDPE/10% PP, 80% LDPE/20% PP, 70% LDPE/30% PP, 50% LDPE/50% PP, 30% LDPE/70% PP, 20% LDPE/80% PP, 10% LDPE/90% PP, neat PP), the blends were abbreviated as PPx, where x was the mass fraction of PP in blends. After mixing, the blends were cut into small pieces and hot-pressed at 180 °C for 5 min, followed by cold pressing at room temperature to form about 1 mm-thick sheets.

Subsequently, the samples, enclosed into sealed polyethylene bags with limited air, were irradiated to 5, 10, 15, 25, 50 or 100 kGy using a 3.0 MeV electron accelerator which was located at the China-Kinwa High Technology Co. Ltd (Changchun, China). And then the irradiated samples were dried in an oven at 80 °C for 12 hours to minimize post-irradiation effects.

2.2 Characterizations

The morphology of non-irradiated samples were observed by a field emission scanning electron microscopy (SEM) (XL30 ESEM FEG, FEI Co., Eindhoven, The Netherlands) at an accelerating voltage of 10 kV. Before completely dried in a vacuum oven at 80 °C overnight, the non-irradiated samples were fractured in liquid nitrogen. The fracture surface was coated with a thin layer of gold before the measurement.

Differential scanning calorimetry (DSC) was used to determine the shape memory transition temperature of non-irradiated and irradiated LDPE/PP blend films. Thermal analysis was performed under nitrogen atmosphere using a TA DSC Q20 (USA). A “heat/cool/heat” cycle was utilized, the heating scans ranged from 30 °C to 200 °C at a cooling and heating rate of 10 K min⁻¹ with sample weights of 5–8 mg. The second heating cycle was of interest in this study.

The static mechanical properties of the irradiated samples were determined by an Instron 1211 testing machine (USA). All tests were conducted at a crosshead speed of 10 mm min⁻¹ at room temperature according to ASTM D638-2008. At least five specimens for each sample were measured to get an average value. Strain–stress plots were recorded during the experiment.

Rheological measurements were carried out on a rheometer (AR2000EX, TA Instruments-Waters LLC, USA) equipped with a parallel plate geometry using plates of 25 mm diameter. Sheet samples of about 1.0 mm were melted at 180 °C under an N₂ atmosphere. The oscillatory frequency sweep ranged from 100 to 0.1, with a fixed strain of 1.25%.

Gel fractions were measured by Soxhlet extractor with a relatively large volume of boiling xylene (refluxing 36–48 h). Then the extracted samples were dried completely in a vacuum oven at 80 °C and about 0.8 atm for 24 h, and the weight were massed again. The gel fraction was calculated as follows

1 − gel fraction(%) = sol fraction(%)

where m₀ and m₁ were the mass of the sample before and after the solvent extraction, respectively. For the reproducibility, all the reported values were an average of the three different tests.

To examine the shape memory behavior, samples were first cut into rectangular specimens of approximately 20 mm × 5 mm × 1.0 mm. Then shape memory test was carried out according to the following program: (1) heating up the sample to 130 °C and holding for 5 min. (2) Stretching the sample to 100% (ε_m). (3) Cooling down the samples to 25 °C under the stress of ε_m. (4) Unloading the sample to zero stress and then recording the strain (ε_u). (5) Heating up the sample to 130 °C without load and keeping isothermal for 5 min, and then recording the strain (ε_n). For a better illustration, the schematic description was depicted in Fig. 1.

Fig. 1 Schematic description of a dual-shape memory testing.

In this experiment, all thermomechanical cycles were repeated at least three times for the reproducibility of dual-SME. The shape fixity ratio (R_f) and the shape recovery ratio (R_r)^3,26 could be calculated from relations 1 and 2, respectively.

3. Results and discussion

3.1 Architecture/morphology

According to the previous work,¹⁴ phase architecture and morphology are known to play a decisive role in determining the physical properties and the shape memory properties of polymer blend. Fig. 2 shows SEM images of the cryo-fractured surface of non-irradiated blends with various content of PP ω^PP. It can be clearly seen that the blend films presented a composition dependent structure, as would be expected based on basic knowledge of polymer blend morphology.^38,39 With the increasing content of PP, the LDPE/PP blends showed a shift in morphology from a droplet structure to co-continuous and back to droplet. More specifically, the PP10 and PP80 exhibited typical sea-island architectures, while the PP30 and PP50 showed co-continuous phase morphologies. Moreover, phase separations also could be observed in their blends, which implied the immiscibility between LDPE and PP. Based on the above discussions, we can conclude that the co-continuous window for non-irradiated LDPE/PP blends is the ω^PP of 30–70 wt%.

Fig. 2 The SEM micrographs of cryo-fractured surfaces of non-irradiated LDPE/PP samples. (a) PP10. (b) PP30. (c) PP50. (d) PP80.

3.2 Thermal, mechanical and rheological properties

Fig. 3 showed the DSC plots of non-irradiated and 10 kGy irradiated LDPE/PP blends with various PP weight content ω^PP. From the DSC thermograms, a consistent LDPE melting temperature was observed between all compositions, thus displaying that the switching temperature is independent of composition and allows for the tuning of mechanical properties. In addition, there were a slight decrease of T_m,LDPE (111 to 110 °C after irradiation) and an another relatively large decline of T_m,PP (162 to 152 °C when ω^PP exceeded 30 wt% and 166 to 162 °C when ω^PP was less than 30 wt%), indicating that 10 kGy irradiation may cause lightly chain crosslinking of the blends which restricted polymer chain mobility and hindered its crystallization.

Fig. 3 The DSC plots of LDPE/PP blends. (a) T_m of non-irradiated samples. (b) T_m of 10 KGy irradiated samples.

The mechanical properties of the binary polymer blends were determined by tensile tests at ambient temperature (about 25 °C), below T_m,LDPE and T_m,PP, which are presented in Fig. 4. At 25 °C, the materials were elastic with ε_b between around 30% and almost 673% when the PP content was less than 30 wt%. However, the samples were transformed into brittle behavior with a continuous increase of ω^PP. Meanwhile, the Young's modulus E_t increased from 162 MPa to 1026 MPa and the fracture energy at break U_b decreased from 5851 mJ to 61 mJ with increasing weight content of PP. In other words, the crystalline domains act as physical crosslinks and promote the stiffness of the blends. After irradiation at 10 kGy, the changes and tendencies kept the nearly same except for an increase of σ_b and a decrease of ε_b. This may be attributed to formation of the crosslinking network structure in the LDPE/PP blends.

Fig. 4 Stress–strain plot of (a) non-irradiated and (b) 10 KGy irradiated blends; the dependence of Young's modulus (E_t), strain at break (ε_b), stress at break (σ_b) and fracture energy at break (U_b) of (c) non-irradiated and (d) 10 KGy irradiated blends as a function of PP weight content.

Fig. 5 shows the complex storage modulus G′, loss modulus G′′and viscosity |η*| versus frequency for non-irradiated and 10 kGy irradiated LDPE/PP blends at 180 °C, above T_m,LDPE and T_m,PP. Obviously, the difference mainly appeared in the terminal (low frequency) zone. For the non-irradiated PP-rich samples, a typical Newtonian fluid was showed in a wide low-frequency range and pseudoplastic shear thinning behavior when the frequency increased. While for the non-irradiated LDPE-rich samples, the low-frequency viscosity plateau became unclear and finally disappeared with increasing content of LDPE. After irradiated at 10 kGy, the LDPE-rich samples exhibited much higher G′, G′′ and |η*| values, which unlike the PP-rich samples. These differences could be explained by the presence of the formed crosslinking network acting like obstacles to resist the movement of the polymer molecular chains, resulting in an enhancement of the G′, G′′ and |η*|.

Fig. 5 Rheological properties (storage modulus G′, loss modulus G′′ and complex viscosity |η*|) of non-irradiated (a, b, c) and irradiated (a*, b*, c*) LDPE/PP blends versus frequency at 180 °C. (Data for irradiated neat PP were not presented due to its large brittleness which made sample preparation difficult).

Meanwhile, Han plots for non-irradiated and irradiated LDPE/PP blends were measured at 180 °C and the results are presented in Fig. 6. As can be seen from the graph, only the irradiated LDPE-rich samples showed a solid-like (G′ > G′′) melt behavior over the entire angular frequency range. Furthermore, the elasticity of the samples gradually decreased with increasing PP content, this effect was more significant for the irradiated blends, which can be ascribed to the formation of crosslinks upon irradiation.

Fig. 6 Han plots of G′ versus G′′ for LDPE/PP blends at 180 °C: (a) non-irradiated; (b) irradiated at 10 kGy. (Data for irradiated neat PP were not presented due to its large brittleness which made sample preparation difficult).

According to the above discussions, it could be concluded that the mechanical properties (below T_trans) and rheological properties (above T_trans) can be tuned by content fluctuations and e-beam irradiation without apparent great changes of the switching temperature. In addition, the results simultaneously revealed that irradiated LDPE-rich blends may be more suitable than irradiated PP-rich blends for shape recovery due to their relatively higher ductility and elasticity.

3.3 Shape memory properties of non-irradiated and irradiated blends

In order to confirm the influence of e-beam radiation on the LDPE/PP blends, gel fraction percent of the 10 kGy irradiated samples is exhibited in Fig. 7. As the plot clearly displayed, the gel fraction percent was extremely low when PP content ω^PP exceeded 50 wt%, suggesting that the effect of radiation degradation predominates over the crosslinking process. However, for the LDPE-rich samples, it was first slightly reduced with increasing ω^PP to 30 wt%, and then the decrease became phenomenal up to 50 wt% PP. This phenomenon implied that e-beam irradiation led to a light but predominant radiation crosslinking of the LDPE-rich samples. Besides, the two-phase interface area may also affect irradiation crosslinking effect, probably because the irradiated PP30 sample did not show a sudden large drop of gel fraction percent.

Fig. 7 Gel fraction and sol fraction of irradiated LDPE/PP blends with different ω^PP at 10 kGy.

Then, shape memory behavior was measured quantitatively and qualitatively, respectively. Comparison of the macroscopic shape memory effect between the non-irradiated and irradiated samples is presented in Fig. 8. After exposed to a similar thermomechanical treatment, the programmed samples were heated to 130 °C to induce the shape memory effect. Both the non-irradiated and irradiated PP-rich samples (PP70) nearly did not change shape due to their large brittleness, while the non-irradiated LDPE-rich sample (PP30) was almost melting and became extremely soft, indicating that the physical crosslinks of PP crystalline domains were not stable enough to quantitatively sustain the permanent shape in the thermomechanical cycle when these materials in their deformed temporary shape were heated above T_m,LDPE. In addition, only the irradiated LDPE-rich sample (PP30) showed relatively satisfying shape memory effect. In this way, it has been demonstrated that the formation of a radiation crosslinking network in the LDPE/PP blends could be beneficial for the pronounced shape memory capability.

Fig. 8 The series of photographs demonstrates the macroscopic shape-memory effect of irradiated LDPE-rich sample c (PP30) with non-irradiated LDPE-rich sample a (PP30), non-irradiated PP-rich sample b (PP70) and irradiated PP-rich sample d (PP70) as reference samples. (1) Permanent shape at room temperature. (2) Temporary shape of these materials at room temperature obtained by deforming at 130 °C. (3) Shape recovery or change when heated to 130 °C.

On the other hand, the shape memory properties of the polymer blends were quantified by two parameters, shape fixity and shape recovery ratios. Both ratios are determined with values obtained from the stress–strain curves during thermo-mechanical cycling (see Fig. 1) between T_low = 25 °C and T_high = 130 °C. The shape fixity is a measure of how well a temporary shape is maintained after forming while the shape recovery describes the dimensional difference between the original permanent shape created during initial processing and the recovered shape achieved after thermomechanical cycling. In Fig. 9a, shape fixity and recovery ratios for the 10 kGy irradiated LDPE/PP blends showed an initial just consistent trend, particularly when the ω^PP did not exceed 30 wt%. After the point, both the two parameters appeared to a tremendous decrease. Ultimately, the values even became invalid for samples with ω^PP over 70 wt%.

Fig. 9 (a) Shape fixity ratio (R_f) and recovery ratio (R_r) of each irradiated LDPE/PP sample (data for irradiated samples with ω^PP above 70% were not presented due to its large brittleness which made shape memory test difficult); (b) comparison and illustration showing the differences and changes between non-irradiated and irradiated samples.

To explicate these differences and changes between non-irradiated and irradiated samples, we finally gave the corresponding comparison and illustration, Fig. 9b. Here, according to some previous work 5,¹¹ a schematic model of the thermoplastic SMP was proposed. The SMP consist of a “hard phase”, which determines the shape and a “soft phase”, which can be triggered by external stimuli. Generally, the hard phase can be physical or chemical crosslinks while the soft phase can be either crystalline or amorphous domains. Here, for all samples, LDPE crystal served as soft phases due to the selected transition temperature over T_m,LDPE but below T_m,PP. However, hard phases were significantly different between non-irradiated and irradiated blends because the e-beam irradiation caused crosslinking for LDPE and degradations for PP. Based on the above results, comprehensive explanations and assessments were thus established to help comprehend the dual shape memory effect.

Therefore, for irradiated blends, with decreasing weight ratio of the soft phase LDPE, the dominant part of the hard phase underwent a transition from more efficient radiation crosslinking networks to relatively inefficient physical crosslinks of crystalline PP domains to determine the temporary shape, resulting in a decline of the shape recovery ratio. Simultaneously, a similar tendency appeared for the values of the shape fixity, not only due to the decrease of LDPE phase but also because the formed network hinders the crystallization of LDPE. Clearly, the influence of content fluctuations predominates over the e-beam irradiation (see Fig. 9a), which could be attributed to the formation of a light network caused by 10 kGy irradiation.

3.4 Effects of radiation dose on shape memory efficiency

The method of e-beam irradiation is widely used for the improvement of mechanical properties of polymer and functional features such as shape memory.⁴⁰ Usually, the changes in polymer properties, including viscosity and crosslink density are increasing with radiation dose if the crosslinking processes prevail over the degradation.^41,42 Fig. 10 shows the effects of network formation (gel fraction) as a function of radiation dose for the 70% LDPE/30% PP blend. Apparently, the gel fraction was dependent on radiation dose, indicating that radiation crosslinking proceeded and this process exceeded the main chain scission.

Fig. 10 Gel fraction as a function of radiation dose for 70% LDPE/30% PP blends.

Measurements of rheological properties at 180 °C also confirmed this observation indicating an increase of complex viscosity and a tendency to more elastic behavior (G′ > G′′) with increasing radiation dose, as presented in Fig. 11.

Fig. 11 Complex viscosity |η*| versus frequency (a) and Han plots of G′ versus G′′ (b) for 70% LDPE/30% PP blends exposed to different radiation dose.

In Fig. 12a the shape memory performance of the 70% LDPE/30% PP blend as a function of radiation dose is depicted. The recovery ratio significantly increased from 17.5% to 93.8% with the absorbed dose increasing from 5 kGy to 100 kGy, while values determined for fixity ratio follow an inverse trend, the values showed a relatively slight decrease and just fluctuated between 85.9% and 78.6%. Similarly, in order to have a better understanding of the effects of radiation dose on shape memory efficiency, the schematic corresponding illustrations and summaries are showed in Fig. 12b. From the above discussion, increased radiation dose could definitely lead to increase the extent of crosslinking network formation, which gave rise to enhanced rubbery properties. Because shape recovery ratio was dominated by the elastic strength and the amount of strain energy stored in the hard phase, the samples also receive much greater recoverability with radiation dose increasing. While the fixity ratio possessed a slight decrease, this might be caused by the occurrence of continuously enhanced network which restrained crystallization of LDPE domains.

Fig. 12 (a) Shape fixity ratio (R_f) and recovery ratio (R_r) as a function of radiation dose for 70% LDPE/30% PP blends; (b) corresponding illustrations and summaries showing the changes.

4. Conclusion

A radiation crosslinked dual-SMP blend system based on two immiscible thermoplastic polymers, LDPE and PP, was designed and fabricated. The “soft phase” was provided by LDPE crystalline domains, while physical crosslinks of PP crystalline domains and radiation crosslinking networks dominantly served as “hard phase” before and after the e-beam irradiation respectively. In order to illuminate the underlying effects of phase morphologies, content fluctuations and e-beam irradiation on the shape memory performances, SEM, DSC, tensile test, rheological properties measurement, sol/gel weight fraction analysis and shape memory test were performed sequentially. On basis of the results, the following conclusions can be drawn:

(1) A close correlation was observed between the content fluctuations, phase morphologies and the shape memory properties. The content fluctuations of soft phase and hard phase caused an alteration of the architecture in the thermoplastic blends, which finally, resulted in numerical changes of the shape fixity and recovery ratio.

(2) The shape memory test data showed that irradiated LDPE-rich blends were more suitable and effective than both irradiated PP-rich and non-irradiated blends to be dual-SMPs with advantageous shape memory properties. Furthermore, increased radiation dose could also give rise to enhanced shape recovery capacity without significantly weakening the shape fixity.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c5ra11044k

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