Structure, morphology and properties of epoxy networks with dangling chains cured by anhydride

Abstract

A series of epoxy networks containing side aliphatic dangling chains were prepared by the curing reaction of bisphenol F diglycidyl ether (DGEBF), the monofunctional epoxy (including n-butyl, dodecyl, cetyl) and the anhydride. In the cured epoxy networks with the flexible side aliphatic chains, micro-phase separations with different dimension were observed by atomic force microscopy (AFM). Scanning electron microscopy (SEM) measurements implied the presence of a ductile fracture on the fracture surfaces of the modified samples. The result of wide angle X-ray scattering (WAXS) characterization revealed no crystallizations in the modified epoxy resins. Significantly, the nano-sized and submicron-sized phase separations coexisting were observed in the cured E₄-20 (molar ratio of epoxy group between DGEBF and butyl glycidyl ether was 20 : 1). Subsequently, by means of adjusting the length and content of the flexible side aliphatic chains, the mechanical and thermal properties of the cured resins can be tuned in different scales. It was also observed in their dynamic mechanical properties that the storage modulus of modified samples was higher than that of the unmodified sample at glass state regions. Especially, mechanical property tests showed that the cured E₄-20 sample demonstrated excellent mechanical properties with a tensile strength of 98.5 MPa, an elongation at break of 6.1%, and a tensile modulus of 3.1 GPa.

---

Introduction

Epoxy resins are versatile thermosetting materials. Owing to their excellent engineering properties, such as their high stiffness and strength, dimensional stability and adhesive properties, epoxy resins are widely used in various fields.¹ However, their applications in advanced fields requiring high fracture strength are limited because of their high degree of chemical crosslinking structures. Consequently, a considerable amount of research has been conducted to improve the toughness of epoxy polymers. And, the most effective approach is to introduce separation phases into epoxy resins, such as rubber particles,^2–6 high performance thermoplastics,^7–10 or nano-materials.^11–18 It is well-known that the reactive liquid rubbers were first studied, frequently used and widely accepted among these strategies.^6,19 Generally, the rubber in epoxy resins was a distinct separation phase. When the rubber phases particle were formed in the epoxy matrix, the fracture toughness was improved significantly through rubber cavitation and deformation void growth, crack tip blunting and stress distribution.^19,20 Unfortunately, although toughness of epoxy resins was improved greatly by the addition of soft modifiers, the glass transition temperature (T_g) and Young's modulus were decreased obviously.¹⁹ To avoid the disadvantages, researchers have investigated rigid thermoplastics to address these issues. During the epoxy curing, the miscibility of another polymer in an epoxy network generated different dimension of separations which can improve the matrix properties. However, these thermoplastic modifiers were rather difficult to process and sometimes solvents were needed.^9,21 It has also been studied that the use of hard nanoparticles (or nano fillers) reinforced epoxy-based thermosets. However, sometimes the uniform distribution of nanoparticles (or nano fillers) was a problem.²²

Recently years, studies have focused on novel approaches to toughen epoxy resins, for instance, forming phase separations in systems through the structural design. Giuseppe R. Palmese et al.²³ introducing flexible polyetheramine into epoxy networks, provided a method to synthesize polymer networks ‘polymer network isomers’ with improved thermal and mechanical properties. Besides, X. Ramis et al.²⁴ reported the use of multiarm star polymers with a hyperbranched poly(ethyleneimine) (PEI) core and poly(ε-caprolactone) (PCL) arms end-capped with acetyl groups as modifiers in the curing of diglycidylether of bisphenol A and amine curing agent. Consequently, the decrease on T_g and glassy modulus led to a significant reduction in internal stresses. Stefano Pandini et al.²⁵ introduced the phenyl group in side location of main chains through curing reaction, and the effect of crosslink density of networks on properties of materials was investigated. Recently we have synthesized a series of epoxy resins containing side aliphatic pendant chains by the reaction between the diglycidyl ether of bisphenol F (DGEBF) and the single-ended fatty amine.²⁶ The thermosets possessed excellent mechanical and thermal properties. However, the problem of this strategy was a sharp increase of the viscosity of the blends. Taking all these strategies into account, adopting methods to get micro–nanoheterogeneity was an effective measurement to improve properties of epoxy resins.

Inspired by these works, we introduced flexible side aliphatic chains into epoxy systems by the reaction between the monofunctional epoxy and DGEBF/anhydride blends. Because flexible aliphatic chains were incompatible with rigid backbone of benzene ring, they could self-organize to the dispersed phases in nano-size or micro-size. Meanwhile the decrease in viscosity of the blends was obvious comparing with the strategy we recently reported. In this work, three kinds of the monofunctional epoxies with alkyl chains (n-butyl, dodecyl, cetyl) were introduced in the cured epoxy networks and the effect of the length and content of the alkyl chains on the film morphology was investigated. Besides, the influence of nanostructures on the thermal and mechanical properties of the cured epoxy resins were examined respectively and the toughening mechanisms were also discussed.

Experimental section

Materials

Bisphenol F diglycidyl ether (DGEBF, equivalent epoxy weight 165.3 g mol⁻¹) was purchased from Epoxy Division of China Petrochemical Group Baling Petrochemical Co., Ltd. Butyl glycidyl ether (BGE, equivalent epoxy weight 130.18 g mol⁻¹) was purchased from Saien chemical technology Co., Ltd in China. Dodecyl glycidyl ether (DGE, equivalent epoxy weight 242.39 g mol⁻¹) was supplied by Adamas Reagent Co., Ltd in China. 1,2-Epoxyoctadecane (equivalent epoxy weight 268.48 g mol⁻¹) was purchased from Beijing HWRK Chem Co., Ltd. Hexahydrophthalic anhydride (HHPA, 99%) was purchased from the Puyang Huicheng Electronic Materials Co., Ltd. Tris-(dimethylaminomethyl)phenol (DMP-30) was purchased from Sinopharm Chemical Reagent Beijing Co., Ltd.

Preparation of thermoset epoxy resin networks with side aliphatic chains

DGEBF was mixed with BGE, DGE and 1,2-epoxyoctadecane respectively (molar ratio of epoxy group of DGEBF between the monofunctional epoxy was 20 : 1 and 40 : 1), subsequently, the stoichiometric ratio of HHPA and a calculated amount of DMP-30 were added, and then the blend was stirred for 15 min until it was homogeneous. All the blends were poured into a mold, and then heated at 90 °C for 3 h, at 120 °C for 3 h and at 150 °C for 1 h. The E₄, E₁₂ and E₁₆ were the abbreviation of cured samples of DGEBF with different monofunctional epoxy, where the subscript referred to the carbon number in the alkyl chains of the monofunctional epoxy and unmodified epoxy sample was abbreviated as E₀. The detailed formulations of the cured samples were listed in Table 1.

Table 1 Formulations of the sample (molar ratio of epoxy group)^a

Sample	E0	E4-20	E4-40	E12-20	E12-40	E16-20	E16-40
a The accelerator DMP-30 was used at a level of 0.5 wt% of HHPA.
BGE	—	2	1	—	—	—	—
DGE	—	—	—	2	1	—	—
1,2-Epoxyoctadecane	—	—	—	—	—	2	1
DGEBF	40	40	40	40	40	40	40
HHPA	40	42	41	42	41	42	41

---

Characterization

The Fourier Transform Infrared Spectroscopy (FTIR) spectra of all samples were tested on a spectrometer (Nicolet Nexus 670) with the resolution of 4 cm⁻¹ and 32 scans in the standard wave number range of 400–4000 cm⁻¹ by using KBr pellet in transmission mode.

Morphology of phase separations was characterized by tapping mode atomic force microscopy (AFM, DMFASTSCAN2-SYS, Germany). Tapping mode was employed in air using a tip fabricated from silicon (125 μm in length with 1.3 MHz resonant frequency).

The tensile fracture surfaces of the broken specimens were observed by a scanning electron microscope (SEM, Hitachi Limited, S-4700) at an accelerating voltage of 5 kV. All the surfaces of specimens were coated with gold to improve the conductivity and prevent charging.

Crystallization of these dispersed phase structures was examined by wide angle X-ray scattering (WAXS) measurements which were taken on a Rigaku D/Max 2500VB2+/PC diffract meter with Cu Kα radiation.

Tensile experiments were carried out by a tensile tester (SANS UTM5205XHD, China) at room temperature according to Chinese National Standard GB/T 1040.2-2006 and Charpy impact test was measured by a XJJD impact tester (Jinjian instrument testing Co., Ltd, China) according to Chinese National Standard GB/T 1043.1-2008, respectively.

Differential scanning calorimetry (DSC) measurements were used to study the thermal behavior of epoxy system by TA Instruments Q20 DSC equipped with an RCS 90 cooling system. The 5–10 mg sample was sealed in aluminum crucibles with the following thermal program (dry nitrogen was used as the protective atmosphere, 50 mL min⁻¹): heating from 40 °C to 180 °C at 20 °C min⁻¹, cooling to 40 °C at the maximum cooling rate, and then reheating to 180 °C at 10 °C min⁻¹. The values of the glass transition temperature (T_g,DSC) were measured from the second heating trace.

Dynamic mechanical properties were studied by dynamic mechanical analyzer (Q800, TA Instruments). Samples were tested with a double cantilever mode at heating rate of 5 °C min⁻¹ and an oscillatory frequency of 1 Hz. The glass transition temperature (T_g,DMA) defined by the tan δ peaks were obtained from the DMA spectra. Specimen dimensions were 60 mm × 5.9 mm × 2 mm.

Results and discussion

FTIR characterization

Fig. 1 shows the FTIR spectra of E₀ and E₄-20 samples. E₄-20 sample was as an example of all the modified epoxy. It can be seen from Fig. 1 that the absorption peaks (anhydride C=O stretching) near 1790 cm⁻¹ and 1864 cm⁻¹ disappeared in the cured samples. Besides, the absorption peak (ester C=O stretching) near 1730 cm⁻¹ appeared, and epoxide absorption peaks (913 cm⁻¹) disappeared. Based on the results, these are clear signs of the epoxy–anhydride curing, which indicates that the monofunctional epoxy is carried out reaction with HHPA completely, namely, the thermoset epoxy resin networks with side aliphatic chains are prepared successfully.

Fig. 1 FT-IR spectra of uncured E₀, E₄-20 and cured E₀, E₄-20.

Morphological behaviour

Fig. 2 shows the phase and 3D height images of E₀, E₄, E₁₂ and E₁₆. As shown in the AFM images, obviously, the morphology of E₀ was continuous phase, as expected for monophasic materials (Fig. 2a and a′). In contrast, all modified samples clearly showed different shape and size of dispersed phases which confirmed the idea proposed above. The different flexible alkyl chains were found to cause the variations in the size of the phase separations. For the E₄ system, numerous 23–30 nm nano-sized bump structures and 200–400 nm submicron-sized bump structures were coexisted in E₄-20 (Fig. 2b and b′). With the decrease in content of the modifier, the size of dispersed phases became small (Fig. 2c and c′). Similarly, multidimensional phase structures were also found in E₁₂-40 (Fig. 2e and e′). Besides, with longer alky chains in E₁₆ system, the relatively large phase separation structures were found, and more content of modifier contributed to the coalescence of dispersed phase in E₁₆-20 (Fig. 2f and f′). The phenomenon is ascribed to longer and more content flexible alkyl chains tending to form larger scale phase separation. Hence, under the comprehensive effect of the length and the content of flexible alkyl chains, appropriate length and content of alky chains in the epoxy networks can form multidimensional dispersed phases coexisted in AFM images. The formation mechanism of phase separations is speculated on that due to the incompatibility of side aliphatic dangling chains with rigid backbone of benzene ring in the epoxy system, those aliphatic chains can gather together to form physical aggregations in the process of curing. Over time, such aggregations become larger and larger and form nano- and submicro-dispersed phase structures in the end. As mentioned in the introduction, such nano-sized and submicron-sized phase separations can toughen epoxy resins effectively.

Fig. 2 AFM images of E₀ (a, phase; a′, 3D height), E₄-20 (b, phase; b′, 3D height), E₄-40 (c, phase; c′, 3D height), E₁₂-20 (d, phase; d′, 3D height), E₁₂-40 (e, phase; e′, 3D height), E₁₆-20 (f, phase; f′, 3D height) and E₁₆-40 (g, phase; g′, 3D height).

WAXS characterization

Fig. 3 shows the WAXS curves of E₀, E₄-20, E₁₂-20 and E₁₆-20. It can be seen from Fig. 3 that E₀ was the typical amorphous polymer, which demonstrated a un-diffraction curve in the WAXS characterization.²⁷ Similarly, diffraction peaks were not observed in E₄-20, E₁₂-20 and E₁₆-20, which contained maximum content of modifiers. The result indicates that these bump structures are self-assembled structures of flexible aliphatic chains instead of similar polyethylene crystallizations.

Fig. 3 X-ray diffraction diagrams of cured E₀, E₄-20, E₁₂-20 and E₁₆-20.

Mechanical properties

Tests were conducted to study the effect of nanostructures on mechanical properties and the results were listed in Table 2. Significantly, due to introducing the phase separations into epoxy networks, excellent improvements in the mechanical properties of the cured epoxy samples were achieved. By means of adjusting the content and length of the aliphatic chains, the tensile strength, the tensile elongation at break and the tensile modulus can reach to the maximum of 98.5 MPa, 6.1% and 3.1 GPa respectively for the E₄-20 sample. Meanwhile, the tensile strength of E₁₂-40 sample was also tested with obvious improvement by 91.7 MPa. However, these properties of samples with single size of phase separation had no obvious improvement compared with that of E₄-20 and E₁₂-40, these results of mechanical tests coincided with the morphologies, showed in AFM phase images, that nano-sized and submicron-sized phase separations were coexisted in the E₄-20 and E₁₂-40 samples respectively. This phenomenon indicates the coexistence of these two different toughening mechanisms in E₄-20 and E₁₂-40 samples, which may explain why the modified systems possessed much higher mechanical properties. As Evans et al.²⁰ reported a multiplicative effect of bimodal rubber-particle size distribution toughening epoxy resins.

Table 2 Mechanical properties of E₀, E₄, E₁₂ and E₁₆

Cured samples	E0	E4-20	E4-40	E12-20	E12-40	E16-20	E16-40
Impact strength (kJ m^−2)	14.5 ± 1.4	22.5 ± 2.0	14.5 ± 2.0	19.7 ± 1.2	21.4 ± 1.4	13.1 ± 1.8	9.8 ± 2.1
Tensile strength (MPa)	62.3 ± 2.1	98.5 ± 0.9	63.7 ± 1.6	66.6 ± 2.2	91.7 ± 2.7	70.1 ± 2.3	66.7 ± 2.3
Strain elongation (%)	2.7 ± 0.1	6.1 ± 0.1	2.8 ± 0.1	3.0 ± 0.3	4.9 ± 0.9	3.2 ± 0.2	3.1 ± 0.2
Tensile modulus (GPa)	2.9 ± 0.1	3.1 ± 0.1	3.1 ± 0.1	3.1 ± 0.1	3.1 ± 0.1	3.1 ± 0.1	3.1 ± 0.1

---

Fig. 4 shows the morphology of the tensile fracture surface of epoxy systems measured by SEM. The unmodified sample was a homogeneous material, with a relatively smooth surface, and showed no sign of plastic deformation, representing typically brittle fracture.²⁸ In contrast, the modified samples showed ductile fracture surfaces with yielding which was covered massive crazing and stress-whitening. Generally, shear yielding induced by cavitations was considered the major energy absorption mechanism in neat epoxies.^29–31 Similar to rubber toughening epoxy systems, submicron-sized phase separation structures enhanced shear localization by acting as stress concentrators. The process included matrix deformation and cavitations of the structures in response to the stresses near the crack tip. Subsequently, shear yielding occurred between the crazing formed by the cavitated submicron-sized phase structures. Plastic deformation blunted the crack tip, which reduced the local stress concentration and allowed the material to support higher loads before failure occurs.³² In addition, the role of the nano-sized phase structures was to restrict the microcrack to propagate within the ligament, and made the crack tip region sustain even higher fracture loading by maintaining a higher critical stress level. This higher critical stress level may even generate greater triaxial stress ahead of the crack tip. This produced higher dilatation and caused a higher degree of cavitation in the large nanostructure phases. Finally, under the synergic effects of various sizes of phase separations toughening epoxy resins, the mechanical properties of thermosets was improved dramatically.

Fig. 4 SEM micro-morphology of the fracture surfaces of E₀ (a), E₄-20 (b), E₄-40 (c), E₁₂-20 (d) E₁₂-40 (e), E₁₆-20 (f) and E₁₆-40 (g).

Dynamic mechanical properties

The thermal mechanical properties of the epoxy resins were investigated by DMA and the results were shown in Fig. 5. With the temperature increasing, all samples demonstrated a distinctive glass transition with the storage modulus (E′) sharply decreasing from a glass state to a rubber state. As shown in the Fig. 5, when the temperature rose, sharp decreases in E′ of E₁₂ and E₁₆ were observed, comparing with E₀ in the glass state. However, different from E₁₂ and E₁₆, the downtrend in E′ of E₄ was similar to that of E₀. Besides, the E′ of E₄-20 was higher than that of E₀ in the glass state. It was noteworthy that all the modified systems demonstrated a higher storage modulus than that of the E₀ at the temperature below 60–75 °C, and then the storage modulus were lower than that of E₀ at the temperature over 75 °C. On one hand, the flexible chains introduced into the epoxy matrix, in the theory,³² can decrease the rigidness of the system, and the decline was more obvious with increasing length and content of the flexible chains. However, in this work, the side aliphatic dangling chains self-assemble to form nano-sized and submicron-sized phase separations mentioned in above, furthermore, due to the chemical bond joints between the side aliphatic chains and backbone, the aggregations of aliphatic chains inevitably lead main chain closer, namely, the matrix networks stack up more closely. Therefore, much closer networks of the system can demonstrate higher storage modulus at a room temperature. On the other hand, with the temperature increasing, the self-assembled physical aggregations are un-jointed gradually. The accumulational network has more free volume which facilitates the network of the system wriggle. As a result, the storage modulus of matrix resins decreases gradually with the plasticizing effect of the flexible alkyl chains and it is lower than that of the E₀ at the temperature over 75 °C in the end (Fig. 6).

Fig. 5 The storage modulus curves of E₀, E₄, E₁₂ and E₁₆.

Fig. 6 The tan δ curves of E₀, E₄, E₁₂ and E₁₆.

The thermal properties

The tan δ peak demonstrates the T_g, whose shape and intensity represent the degree of flexibility and order of a material.³³ The peak of tan δ was observed in the glass-transition region of the cured resin. In the E₄ systems, just slight decrease (2–3 °C) of the T_g,DMA was observed in the region which means that the cross-linking density of the cured resin toughened by BGE was likely to remain at about the same level in the network. However a decrease about 10–20 °C of T_g,DMA was observed in E₁₂ and E₁₆. The thermal properties of the epoxy resins were further investigated by the DSC. Such similarity was also observed in Fig. 7. With increasing content and longer length chain of the monofunctional epoxy being introduced in the matrix, the T_g,DSC of the epoxy resins shifts to lower temperature. Those results are due to the decrease in cross-linking density of the network by introducing flexible aliphatic chains to systems,³⁴ and the chains of network move more easily from a glass state to a rubber state. In addition, more content and longer aliphatic chains make the effect more apparent. Significantly, corresponding with above dynamic mechanical properties tests, the epoxy resin toughened by BGE also has a slight decrease of T_g, which indicates that the E₄-20 sample with improved mechanical properties is obtained without sacrificing essential thermal properties.

Fig. 7 DSC thermograms of the cured epoxy.

Conclusion

A series of epoxy systems containing side aliphatic dangling chains with different length were designed and prepared respectively. The aliphatic chains self-organized into nano-sized and submicron-sized phase separations, and diffraction peaks were not observed in the WAXS characterization. The effect of the length and content of flexible chains on mechanical and thermal properties of epoxy networks were investigated. SEM micrographs from fracture surfaces of the modified samples demonstrated massive crazing and stress-whitening, which confirmed that the self-assembled structures resulted in the plastic deformation of the epoxy matrix. The E₄-20 sample showed excellent mechanical properties without sacrificing essential thermal properties. The tensile strength, elongation at break and tensile modulus were increased to 98.5 MPa, 6.1% and 3.1 GPa, respectively. As expected, higher modulus of modified epoxy resins were observed compared with that of E₀ at the temperature blow 60–70 °C. In summary, side aliphatic dangling chains self-assemble into phase-separated structures which improved thermal and mechanical properties of cured epoxy resins.

Acknowledgements

The authors gratefully acknowledge the financial support from the National Key Research Program of China under Grant No. 2016YFB0302000.

Notes and references

1. H. Gu, C. Ma, J. Gu, J. Guo, X. Yan, J. Huang, Q. Zhang and Z. Guo, J. Mater. Chem. C, 2016, 4, 5890–5906.
2. Y. Li, N. Hu, T. Kojima, T. Itoi, T. Watanabe, T. Nakamura, N. Takizawa, T. Inoue, H. Cui, S. Atobe and H. Fukunaga, J. Nanotechnol. Eng. Med., 2012, 3, 011004.
3. D. Quan and A. Ivankovic, Polymer, 2015, 66, 16–28.
4. F. Xu, X. S. Du, H. Y. Liu, W. G. Guo and Y. W. Mai, Composites, Part B, 2016, 95, 423–432.
5. H. Zhou and S. Xu, Mater. Lett., 2014, 121, 238–240.
6. P. P. Vijayan, D. Puglia, H. J. Maria, J. M. Kenny and S. Thomas, RSC Adv., 2013, 3, 24634.
7. R. Francis and D. K. Baby, Ind. Eng. Chem. Res., 2014, 53, 17945–17951.
8. H. Garate, I. Mondragon, N. B. D'Accorso and S. Goyanes, Macromolecules, 2013, 46, 2182–2187.
9. M. Rico, J. López, B. Montero and R. Bellas, Eur. Polym. J., 2012, 48, 1660–1673.
10. X. Zhang, X. Yan, J. Guo, Z. Liu, D. Jiang, Q. He, H. Wei, H. Gu, H. a. Colorado, X. Zhang, S. Wei and Z. Guo, J. Mater. Chem. C, 2015, 3, 162–176.
11. P. He, M. Huang, B. Yu, S. Sprenger and J. Yang, Compos. Sci. Technol., 2016, 129, 46–52.
12. R. Atif, I. Shyha and F. Inam, RSC Adv., 2016, 6, 34188–34197.
13. R. Atif, J. Wei, I. Shyha and F. Inam, RSC Adv., 2016, 6, 1351–1359.
14. J. S. Monfared Zanjani, B. S. Okan, Y. Z. Menceloglu and M. Yildiz, RSC Adv., 2016, 6, 9495–9506.
15. Y. Zhou, L. Li, Y. Chen, H. Zou and M. Liang, RSC Adv., 2015, 5, 98472–98481.
16. O. Eksik, A. Maiorana, S. Spinella, A. Krishnamurthy, S. Weiss, R. A. Gross and N. Koratkar, ACS Sustainable Chem. Eng., 2016, 4, 1715–1721.
17. L. Ruiz-Prez, G. J. Royston, J. P. A. Fairclough and A. J. Ryan, Polymer, 2008, 49, 4475–4488.
18. G. Šebenik, M. Huskić, D. Vengust and M. Žigon, Appl. Clay Sci., 2015, 109–110, 143–150.
19. H.-J. Sue, Polym. Eng. Sci., 1991, 31, 275–288.
20. A. G. Evans, T. K. Chen and Y. H. Jan, J. Mater. Sci., 1992, 27, 111–121.
21. I. A. Zucchi, M. J. Galante and R. J. J. Williams, Polymer, 2005, 46, 2603–2609.
22. F. H. Gojny, M. H. G. Wichmann, B. Fiedler and K. Schulte, Compos. Sci. Technol., 2005, 65(15–16), 2300–2313.
23. M. Shari, C. Jang, C. F. Abrams and G. R. Palmese, Macromolecules, 2015, 48, 7495–7502.
24. C. Acebo, M. Alorda, F. Ferrando, A. Serra, J. M. Morancho, J. M. Salla and X. Ramis, eXPRESS Polym. Lett., 2015, 9(9), 809–823.
25. S. Pandini, F. Bignotti, F. Baldi and L. Sartore, AIP Conf. Proc., 2016, 1736, 020173.
26. T. Yang, R. Wang, X. Hou, J. Cheng and J. Zhang, Mater. Lett., 2016, 166, 150–153.
27. A. Schiraldi and P. Baldini, Epoxy polymers, 1983, vol. 28.
28. W. Liu, R. Zhou, H. L. S. Goh, S. Huang and X. Lu, ACS Appl. Mater. Interfaces, 2014, 6, 5810–5817.
29. C. K. Riew and A. J. Kinloch, Adv. Chem., 1993, vol. 233.
30. R. A. Pearson and A. F. Yee, J. Mater. Sci., 1989, 21, 2462–2474.
31. R. S. Raghava, J. Polym. Sci., Part B: Polym. Phys., 1987, 25(5), 1017–1031.
32. H. C. Erythropel, S. Shipley, A. Börmann, J. A. Nicell, M. Maric and R. L. Leask, Polymer, 2016, 89, 18–27.
33. Y. He, D. Xie and X. Zhang, J. Mater. Sci., 2014, 49, 7339–7352.
34. R. Sengupta, S. Chakraborty, S. Bandyopadhyay, S. Dasgupta, R. Mukhopadhyay, K. Auddy and A. S. Deuri, Engineering, 2007, 47, 21–25.

---