Modification of the Rheological Properties of a Polyamide by Adding a Copolymer Comprising an A-olefin and Maleic Anhydride

1. Introduction

It is essential that engineering plastics including polyamides are able to undergo various processing operations besides injection molding. Most engineering plastics have a narrow molecular weight distribution and no long-chain branches. Therefore, during processes in which a molten resin with a free surface is deformed, it is necessary to modify the rheological properties of the resin to provide strain hardening in transient elongational viscosity. Such processes include foaming, blow-molding, thermoforming, T-die casting, and tubular-blown film processing. Various techniques have been proposed to achieve this modification; (1) the addition of an ultra-high molecular weight fraction [1,2,3], (2) the incorporation of long-chain branches [4,5,6], (3) the addition of flexible nanofibers [7,8,9], (4) the addition of a critical gel [10,11], (5) the addition of a comb-shaped block-copolymer in an immiscible polymer blend [12,13], (6) blending with an immiscible long-chain branched polymer [14,15], and (7) making a multi-layered structure [16].

Polyamides are engineering plastics with high mechanical toughness, abrasion resistance, heat resistance, oil resistance, and gas barrier properties [17,18,19]. They are therefore potentially very useful for various products such as films, bottles, and foams. Reaction at the amide groups of polyamides is a good method of providing long-chain branches and therefore strain hardening in elongational viscosity [20]. Among various chain extenders [21,22,23,24,25,26,27,28,29], maleic anhydride is one of the most appropriate reactive compounds. Maleic anhydride graft copolymers such as polypropylene-graft-maleic anhydride (PP-g-MAH) are widely available and are used as compatibilizers for blends comprising PP and polyamides [30,31,32]. This is logical because the polymer chains of the graft copolymer are localized at the boundary between the two phases, and reduce the possibility of the coalescence of dispersed droplets [33,34]. However, such a graft copolymer is not appropriate for the rheological modification of a polyamide resin because one component of the graft copolymer, e.g., PP in PP-g-MAH, is immiscible with the polyamide resin. An alternative copolymer of 〈-olefin and maleic anhydride (O-MAH, shown in Figure 1), which reacts readily with polyamides, is a good candidate of a rheology modifier. In fact, Li et al. employed O-MAH to modify polyamide 6 (PA6) and confirmed that its foaming processability was improved [29].

However, it is not easy to characterize the rheological properties of conventional polyamides such as PA6 and polyamide 66 (PA66) especially in the low frequency region, owing to severe thermal degradation beyond their melting points [35,36,37,38]. Elongational viscosity is also difficult to evaluate [20,39,40]. Therefore, in the present study a copolyamide resin with low crystallinity was used for evaluation of the basic rheological properties at relatively low temperatures. Subsequently, the effect of the addition of O-MAH on the rheological properties of the polyamide was discussed in detail. The results obtained herein will be widely available to inform modification of the rheological properties and processabilities of various polyamide resins.

2. Results and Discussion

Figure 2 shows the IR spectra of O-MAH, PA, and their blends. Stretching vibrations attributable to the carbonyl groups of O-MAH were detected at 1780 and 1850 cm⁻¹ [29]. These peaks disappeared in PA/O-MAH (30/1), demonstrating that all the maleic anhydride in O-MAH had reacted with the amide groups in the PA during mixing and compression molding. The stretching vibrations in PA/O-MAH (90/10) weakened but did not disappear. Although some of the maleic anhydride reacted with the PA, some remained in the blend. Subsequently, therefore, we employed PA/O-MAH (30/1) as the blend sample because further reactions barely occurred during the rheology measurements.

The light transmittance of each 0.3 mm-thick PA and PA/O-MAH (30/1) film was evaluated using a UV-vis spectrometer, and the results are shown in Figure 3. Photographs of the films are also shown in the figure. The light transmittance increased with wavelength, indicating light absorption in the low wavelength range [35]. Considering that approximately 10% of the light transmittance was lost by surface reflection [41,42], both films were relatively transparent, as demonstrated by the photographs. This must have been due to the lower light scattering of the crystalline structure of PA [43,44,45]. Moreover, the O-MAH addition hardly affected the transparency. The result suggested that the extent of phase separation, if any, must have been small in the blend, and we could not detect any phase-separated structure in the blend by electron microscopy.

The DSC heating and cooling curves are shown in Figure 4. The heating/cooling rate was 10 °C min⁻¹. There was an endothermic melting peak at approximately 104.5 °C in the DSC curve of pure PA, and the heat of fusion of 36.7 J g⁻¹. Assuming that these crystals were composed of the main component of the polymer, i.e., PA6, the crystallinity was calculated to be 19% (the heat of fusion of perfect PA6 crystals was reported to be 230 J g⁻¹) [46]. Because PA crystallization occurred slowly, no crystallization peak was detected at this cooling rate. In the case of pure O-MAH, a main melting peak was detected at approximately 75.5 °C with a weak broad peak at approximately 55 °C. These melting peaks can be attributed to the crystallinity of the 〈-olefin [47,48,49,50]. There was a clear crystallization peak at 61.5 °C in the cooling curve. The PA/O-MAH (30/1) film produced two peaks in the heating curve at approximately 80.3 °C and 107.3 °C. The PA crystallinity was reduced by adding O-MAH. Moreover, a crystallization peak was absent from the curve generated by the blend film.

Figure 5 shows the temperature dependencies of the tensile storage modulus E' and loss modulus E" of the PA and PA/O-MAH (30/1) films. There was a peak in the E" curve at approximately -60 °C, which can be attributed to localized segmental motion [51,52]. This relaxation mode was not affected by the addition of O-MAH. The E' values decreased from approximately 20 °C owing to the glass-to-rubber transition of PA. Correspondingly, there was a peak in the E" curve. The peak temperature, i.e., the glass transition temperature T_g, of the PA film was approximately 30 °C, which was lower than that of PA6 [35]. The T_g of the PA/O-MAH (30/1) film was slightly lower. The decrease in the T_g following the addition of O-MAH was presumably attributable to its lower crystallinity, which increased segmental motion. The E' decreased gradually, indicating that the PA was not fully amorphous, as shown in Figure 4. Beyond 100 °C, E' dropped abruptly owing to the melting of the PA crystals. The E' decrease occurred at 85 °C in the blend film. This must be attributed to the melting of the O-MAH crystals.

To understand the basic characteristics of PA in the molten state, we evaluated the relationship between pressure and specific volume at 230 °C. As shown in Figure 6, a straight line was obtained. The slope provided the bulk modulus K, which was 1.47 GPa and was determined by the following relationship:

$$K={\displaystyle \frac{P}{\left(\Delta V/V_0\right)}},$$

(1)

where V₀ is the volume at P = 0.

The melt density ρ at atmospheric pressure was calculated to be 954 kg m⁻³.

The master curves of the angular frequency dependency of the oscillatory shear moduli, such as shear storage modulus G’ and loss modulus G”, of PA and PA/O-MAH (30/1) are shown in Figure 7. The reference temperature was 230 °C. Although the time–temperature superposition principle was not applicable to PA/O-MAH (30/1), as explained in detail later, the apparent flow activation energy ΔE_a was roughly calculated according to the Arrhenius equation, and was determined to be 72.4 kJ mol⁻¹ for PA/O-MAH (30/1). This value was almost the same as that of pure PA, i.e., 72.5 kJ mol⁻¹. The measurement was also performed at 130 °C in the present experiment. There were no crystals at that temperature because the time–temperature superposition principle was applicable to PA.

As shown in Figure 7, the rheological terminal region was clearly detected in PA. The slopes of G′ and G″ in the low-frequency region were 2 and 1, respectively. The zero-shear viscosity η₀ and the steady-state shear compliance J_e⁰ were calculated using the following equations [53]:

$${\eta }_0=\underset{\omega \to 0}{\lim }{\displaystyle \frac{G^{″}}{\omega }},$$

(2)

$$J_e^0=\underset{\omega \to 0}{\lim }{\displaystyle \frac{G^{\prime }}{G^{″}}},$$

(3)

The rheological terminal parameters at 230 °C were as follows: η₀ = 53.4 Pa s and J_e⁰ = 1.1 × 10⁻⁵ Pa⁻¹. For PA/O-MAH, neither parameter was obtained in the measurement range, suggesting a long-time relaxation mechanism.

As is well known, the rubbery plateau modulus G_N⁰ of a simple polymer melt can be calculated using the following equation [53]:

$$G_N^0={\displaystyle \frac{2}{\pi }}{\displaystyle {\int }_{-\infty }^a{G}^{″}d\ln \omega },$$

(4)

where a is the angular frequency at the upper limit of the terminal region. The value of G_N⁰ is usually obtained by doubling the numerical integration of G″ from ln ⍵ = -∞ to the maximum of G″. The G_N⁰ of PA is 2.1 MPa. Because the melt density ρ was known, the average molecular weight between entanglement coupling by points, i.e., M_e, can be calculated from G_N⁰ [53]:

$$M_e={\displaystyle \frac{\rho RT}{G_N^0}},$$

(5)

where R is the gas constant.

The M_e value of the PA used in the present study was 1900. This value is close to those of other aliphatic polyamide resins, such as PA6 (M_e = 2490) and PA66 (M_e = 2000) [54].

Considering the thermo-rheological complexity detected in PA/O-MAH (30/1), the van Gurp–Palmen plot, i.e., loss angle δ versus absolute value of complex shear modulus |G^*|, is shown in Figure 8. A curve that was typical of most simple polymer liquids was obtained for PA without deviation at any temperature. In contrast, PA/O-MAH (30/1) exhibited a significantly different trend. Although it was possible to superimpose the data obtained at various temperatures onto each other in the high |G^*| region, the deviation was obvious in the low |G^*| region. The δ value decreased as the measurement temperature decreased. This phenomenon has often been detected in long-chain branched polymers owing to the large activation energy in the low modulus region [55,56]. Therefore, the present results demonstrated that the reactions between PA and O-MAH generated long-chain branches.

Figure 9 shows the shear rate $\dot{\gamma }$ dependence of shear stress σ and primary normal stress difference N₁ at 130 °C. Both σ and N₁ increased with increasing $\dot{\gamma }$ in the samples. In addition, the values of N₁ for PA/O-MAH (30/1) were much higher than those for pure PA over the whole shear-rate region. Furthermore, it should be noted that the order of σ and N₁ was reversed, i.e., σ > N₁ for PA and σ < N₁ for the blend, demonstrating that the addition of O-MAH greatly enhanced elasticity over viscosity. This behavior was clearly confirmed by the right figure.

In the high shear rate region, a capillary rheometer was employed to evaluate the steady-state shear viscosity η and the appearance of the extruded strands. In the present study, neither Bagley nor Rabinowitsch corrections were performed. As shown in Figure 10, the steady-state shear viscosity was enhanced to some degree by the addition of O-MAH. However, flow instabilities such as shark-skin failure and volumetric gross melt fracture were absent over the whole shear rate range. This result suggested that the applied shear stress at die exit and elongational stress at die entry were still lower than their critical onset values of flow instabilities [6,57,58,59].

As revealed by the photographs, the diameters of the extruded strands of PA/O-MAH (30/1) were much larger than those of PA, demonstrating that the melt elasticity was greatly increased by the addition of O-MAH. These results corresponded with those in Figure 9.

Figure 11 shows the growth curves of the uniaxial elongational viscosity η_E⁺ at 130 °C of PA/O-MAH (30/1). It was impossible to evaluate the η_E⁺ of pure PA owing to its severe downward deformation due to gravitational force. The numerals in the figure represent the elongational strain rates $\dot{\epsilon }$. There was marked strain hardening at all strain rates. This is typical for a polymer melt with a well-developed long-chain branch structure [4,5,6,30,57,58]. As mentioned in the introduction, such rheological properties are essential for various processing operations.

3. Materials and Methods

A polyamide resin with low crystallinity (PA) was used. Its chemical composition was as follows: PA6 = 55 mol%, PA66 = 13 mol%, and PA610 = 32 mol%. Its number- and weight-average molecular weights were characterized by size exclusion chromatography using poly(methyl methacrylate) as a standard, and were 1.14 × 10⁴ Da and 2.46 × 10⁴ Da, respectively. An alternative 〈-olefin–maleic anhydride copolymer (O-MAH; DIACARNA 30M; Mitsubishi Chemical Corp., Tokyo, Japan) was employed and is depicted in Figure 1 [29].

Before melt-mixing, the PA and O-MAH were dried under vacuum for 4 h at 65 °C. They were then melt-mixed using an internal batch mixer (Labo Plastomill; Toyo Seiki Seisaku-sho Ltd., Tokyo, Japan) at 140 °C for 3 min. The blade rotation speed was 30 rpm. The PA/O-MAH weight fractions were 100/0, 30/1, and 90/10. The obtained blends were compressed into flat films using a compression-molding machine at 140 °C for 1 min, and subsequently cooled at 30 °C for 3 min.

The reaction of O-MAH was evaluated by attenuated total reflectance Fourier-transform infrared (ATR-IR) spectroscopy (Spectrum100; PerkinElmer Co., Ltd., Waltham, MA, USA). KRS-5, i.e., thallium bromoiodide, was used as an ATR crystal.

The light transmittance of each film (0.3 mm thick) was measured at 25 °C using a UV-vis spectrometer (Lambda 25; PerkinElmer Co., Ltd.) in the wavelength range of 200 to 800 nm.

Thermal analysis was conducted using a differential scanning calorimeter (DSC) (DSC8500; PerkinElmer Co., Ltd.) under a nitrogen atmosphere. Each sample was placed in an aluminum pan and heated from 25 to 140 °C at a rate of 10 °C min⁻¹, then cooled at 10 °C min⁻¹. Each sample weighed approximately 5 mg.

The temperature dependencies of the tensile storage and loss moduli (E′ and E″, respectively) were evaluated from -80 to 140 °C using a dynamic mechanical analyzer (Rheogel-E4000; UBM Co., Ltd., Muko, Japan). The measurements were performed at a frequency of 10 Hz and a constant heating rate of 2 °C min⁻¹. Rectangular samples cut from the compression-molded films (5 mm wide, 7 mm long, and 1 mm thick) were employed for the measurements.

The melt density of the PA was measured under various pressures using a pressure-volume-temperature (PVT) measuring machine (PVT Test system; Toyo Seiki Seisaku-sho Ltd.) at 230 °C.

The angular frequency dependencies of the shear storage and loss moduli (G′ and G″, respectively) were evaluated at various temperatures using a rotational rheometer (AR2000ex; TA Instruments, Inc., New Castle, DE, USA) with a cone-and-plate geometry. The diameter of the cone was 25 mm and the cone angle was 4°. The shear stress and primary normal stress difference under steady shear were also measured using this machine at 130 °C. The transient uniaxial elongational viscosity was evaluated using the rheometer equipped with a universal testing platform (SEG2-G; Xpansion Instruments LLC, Tallmadge, OH, USA) at 130 °C. Rectangular specimens (10 mm wide, 15 mm long, and 1 mm thick) were used for the measurements.

Capillary extrusion was performed using a pressure-driven capillary rheometer (140-SAS-2002; Yasuda-Seiki Seisaku-sho Ltd., Nishinomiya, Japan) at 130 °C to evaluate the steady-state shear viscosity and the extrusion processability. A circular die with a length of 10 mm and a diameter of 1 mm was employed. The entrance angle of the die was 180°.

4. Conclusions

Polyamide resins generally have a narrow molecular weight distribution and no long-chain branches. These properties cause difficulties in various processing operations due to poor melt elasticity including an absence of strain hardening in the transient elongational viscosity. In the present study, the rheological properties were evaluated over a wide temperature range using a polyamide resin with low crystallinity. The basic viscoelastic properties of the pure PA, such as the G_N⁰ and M_e, were also determined. The present study verified that the addition of O-MAH successfully modified the rheological properties of the resin through chemical reactions between the amide groups and the maleic anhydride groups. Because O-MAH is not a block copolymer, macroscopic phase separation, which causes light scattering, was not detected. The addition of O-MAH provided the PA with long-chain branches. Consequently, the melt elasticity, i.e., the primary normal stress difference and strain hardening in the transient elongational viscosity, was greatly enhanced. Although the shear viscosity increased to some degree, the extrusion performance was good and there were no flow instabilities. O-MAH is a solid at room temperature and is easy to handle. Therefore, the technique described herein can be applied to various processing operations without any difficulty.