Miscibility, Morphology and Crystallisation Behaviour · PDF fileMiscibility, Morphology and...

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Miscibility, Morphology and Crystallisation Behaviour of Rubber Based Polymer Blends


Zhaobin Qiu, Takayuki Ikehara and Toshio Nishi 77.1 Introduction

Rubber is one of the most widely used polymeric materials. Without rubber products, today’s transportation and engineering industries all over the world could not survive at all. The people of today’s modern society could not enjoy their lives as a result.

Rubber can be mainly classifi ed into two categories, namely natural rubber (NR) and synthetic rubber (SR). Natural rubber (or cis-1,4-polyisoprene) is one of the most well known natural polymers. Ancient Mayans and Aztecs harvested it from the Hevea tree and used it to make waterproof boots and also the balls for them to play a game similar to basketball. The NR is normally treated to give it crosslinks, which makes it an even better elastomer. Like most diene polymers, polyisoprene has a carbon-carbon double bond in its backbone chain. Polyisoprene can be harvested from the sap of the Hevea brasiliensis, but it can also be made by Ziegler-Natta polymerisation. This is a rare example of a natural polymer that we can make almost as well as nature does. Natural rubber products have very useful technical characteristics of good tensile strength, high resilience and excellent fl exibility, and resistance to impact and tear. However, on the other hand, NR is less resistant to oxidation, ozone, weathering and a wide range of chemicals and organic solvents due mainly to its unsaturated chain structure and nonpolarity. These drawbacks must cause limitations in the usage of NR. There are mainly two methods to modify the properties of NR. One is the chemical modifi cation. Chemical modifi cation can be classifi ed into three main categories: modifi cation by bond rearrangement without introducing new atoms, modifi cation by attachment of new chemical groups (like chlorine and epoxy) through addition or substitution reactions at the olefi nic double bonds, and grafting a second polymer onto the NR backbone. The examples of the fi rst category are carbon-carbon crosslinking, cyclisation, cis-trans-isomerisation, and depolymerisation. The typical examples for the second category are chlorinated NR, hydrochlorinated NR and epoxidised NR. For the third category, grafting is mostly carried out using vinyl monomers such as styrene and methacrylate. The other method to modify the properties of NR is to blend it with SR. Compared with the chemical modifi cation method, blending NR with SR is an easy and economical way to produce rubbery materials with suitable properties [1, 2]. Much attention has been paid to this method from the viewpoints of practical application and cost.

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Thermal Analysis of Rubbers and Rubbery Materials

Synthetic rubbers have also been developed to replace NR since World War II. There are three main types of synthetic rubbers. The fi rst type is rubbers with an unsaturated carbon backbone, including polybutadiene (BR), polystyrene-co-butadiene (SBR), polybutadiene-co-acylonitrile (NBR) and so on. The second type is rubbers with a saturated –C–C– main chain, including copolymers of ethylene and propylene (EPR), ethylene-propylene-diene monomer (EPDM) and so on. The third is rubbers with carbon and oxygen in the main chain, such as polyepichlorohydrin (PECH) and polyepichlorohydrin-co-ethylene oxide (PECH-co-EO). Sometimes one synthetic rubber is also blended with the other synthetic rubber in order to modify the properties and extend the application fi eld. Rubber - rubber blends are used widely in industry, for example, in tyre manufacturing.

Rubber is often blended with thermoplastics to prepare new polymeric materials. Rubber is usually an impact modifi er to thermoplastic. These blends, consisting of thermoplastics and rubber, are called as thermoplastic elastomers (TPE), which combine the processing characteristic of plastics at elevated temperature with the physical properties of conventional elastomers at service temperatures. One of the most important advantages of TPE is that the products can be reprocessed and remoulded. Therefore, TPE have started to be used instead of thermoset crosslinked rubber in many applications. TPE can play an increasingly important role in the polymer material industry. So TPE are of signifi cant commercial interest. Several typical TPE are thermoplastic polypropylene (PP), thermoplastic polyurethanes (PU), thermoplastic copolyesters, and thermoplastic polyamides (PA).

Most of the thermoplastics are semicrystalline polymers, such as PP, PA and biodegradable polymers. It is well known that the physical properties of crystalline polymeric materials depend strongly on their crystallinity and their microstructure. The studies of morphology and crystallisation of polymers are very important to get a better understanding of the relationship between structure, properties and processing. It is apparent that the crystallisation of the thermoplastics must be affected by the addition of rubber in the TPE.

In this chapter, we focus on the study of the miscibility, morphology and crystallisation behaviour of plastic/rubber polymer blends. The miscibility and crystallisation behaviour of polymer blends based on biodegradable polymers and rubber will be reviewed fi rst. Second, the morphology and crystallisation behaviour of TPE based on PA will be briefl y introduced based mainly on the authors’ research.

7.2 Miscibility and Crystallisation of Biodegradable Polymer/Rubber Polymer Blends

Miscibility plays a signifi cant role in the structure and properties of polymer blends. Miscibility describes the capability of a mixture to form a single phase over certain ranges


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of temperature, pressure and composition. Whether a single phase exists or not depends on many factors, such as the chemical structure, molar mass distribution, and molecular architecture of the components and so on. The single phase in a polymer blend may be confi rmed by many experimental methods, including light scattering, X-ray scattering and neutron scattering. However, the most convenient one is to study the glass transition temperature (Tg) of the polymer blend using thermal analysis including differential scanning calorimetry (DSC) and dynamic mechanical thermal analysis (DMTA) method. It is well known that a single Tg is the most widely and conventionally used criterion for determining the miscibility of a polymer blend. A single composition-dependent Tg indicates full miscibility with a dimension in the order of 20-40 nm. Conversely, an immiscible polymer blend exhibits more than one Tg. Some equations, such as the Fox equation – Equation 7.1 [3], Gordon-Taylor equation – Equation 7.2 [4] and Kwei equation – Equation 7.3 [5], have been proposed to predict the variation of the glass transition temperature of a miscible polymer blend as a function of composition.

1Tg

=w1

Tg1

+w2

Tg2 (7.1)

Tg =

w1Tg1 +k(1 w1)Tg2

w1 + k(1 w1) (7.2)

Tg =

w1Tg1 +kw2Tg2

w1 +kw2

+qw1w2

(7.3)

Where Tg, Tg1, Tg2 are the glass transition temperatures of the blend, polymer 1 and polymer 2, respectively, w1 and w2 are the weight fraction of polymer 1 and 2, respectively, and k and q are fi tting parameters.

In a crystalline/amorphous polymer blend, the miscibility can also be judged from the depression of the melting point of the crystalline component with the increase of the amorphous component. It is well known that the depression of the melting point of a crystalline polymer blended with an amorphous polymer provides important information about its miscibility and its associated polymer-polymer interaction parameter. An immiscible or partially miscible blend does not typically show the depression of the melting point, which is depressed signifi cantly with increasing the content of the amorphous polymer for a miscible blend, especially one containing specifi c interaction between the components. However, the melting point of a polymer is affected not only by the thermodynamic factors but also by the morphological factors such as crystalline lamellar thickness. Therefore, the equilibrium melting point should be used to separate the morphological effect from the thermodynamic effect in discussing the melting point depression as described by the Flory-Huggins theory [6, 7].

Hoffman and Weeks [8] have shown a relationship between the apparent melting point Tm and the isothermal crystallisation temperature Tc:


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Tm = ηTc + (1-Tc) Tmo (7.4)

where Tmo is the equilibrium melting point, and η may be regarded as a measure of the

stability, i.e., the lamellar thickness, of the crystals undergoing the melting process. The equilibrium melting point can be obtained from the intersection of this line with the Tm = Tc equation [8]. The equilibrium melting point data obtained can be analysed by the Nishi-Wang equation [7] based on the Flory-Huggins theory [6]. The melting point depression is given by Equation 7.5:

1Tm

o blend( )1

Tmo pure( )

=RV2

HoV1

ln 2

m2

+1

m2

1m1

1 + 12 12

(7.5)

where Tmo (pure) and Tm

o (blend) are the equilibrium melting point of the pure crystallisable component and of the blend, respectively. Ho is the molar heat of fusion of the repeat unit for a perfectly crystallisable polymer, V is the molar volume of the repeating units of the polymers, m and are the degree of polymerisation and the volume fraction of the component in the blend, respectively. Subscripts 1 and 2 refer to the amorphous and crystalline polymer, respectively. R is the universal gas constant, and

12 is the polymer-polymer interaction parameter. When both m1 and m2 are large, as for high molecular weight polymers, these related terms in Equation 7.5 can be neglected. The interaction parameter 12 can be written as:

HoV1

RV2

1Tm

o blend( )1

Tmo pure( )

= 12 12

(7.6)

For crystalline/amorphous polymer blends, a negative value of 12 indicates that the two components are thermodynamically miscible in the melt. Furthermore, the polymer-polymer interaction becomes stronger with the decrease of the value of 12.

The development and application of biodegradable polymers has recently received more and more attention because of environment protection and the recycling of resources. Based on the difference in the preparation method, biodegradable polymers can usually be classifi ed into two types. One is the biosynthetic polymers, such as bacterial poly(3-hydroxybutyrate) (PHB) and poly(3-hydroxybutyrate-co-hydroxyvalerate) (PHBV). The other is the chemosynthetic polymers, such as the aliphatic polyesters. Polymer blending is often used to improve the properties and extend the application fi eld of biodegradable polymers [9-11]. Through the blending, the properties of biodegradable polymers can be modifi ed signifi cantly, which infl uences not only the performance of biodegradable polymers but also affects their subsequent biodegradation.

Although miscibility and crystallisation of biodegradable polymer blends have been investigated extensively, much less attention has been paid to the blending of biodegradable polymer with rubber. PHB is reported to be immiscible with rubbers,


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such as EPR, ethylene-vinyl acetate (EVA), and poly(cis-1,4-isoprene) (PIP) [12-14]. In PHB/EPR blends, the Tg values for both components do not change with composition at all, which is indicative of their complete immiscibility [12]. Furthermore, although the crystallinity of the PHB phase is only slightly infl uenced by blend composition, no change in the radial growth rate (G) of spherulites occurs with increasing EPR content. It is found that the spherulites of PHB grow in the presence of a two-phase melt consisting of molten PHB containing EPR domains as a dispersed phase in the blend. The EPR particles are fi rst rejected and then occluded in the intraspherulitic region during growth. The resulting morphology consists of PHB spherulites occluding particles of EPR in intraspherulitic regions. The miscibility of PHB with EVA copolymers containing 70 wt% vinyl acetate was investigated. The Tg, the melting temperature and the spherulite growth rate under the isothermal crystallisation conditions are independent of the blend composition, indicating that the two components are immiscible [13]. PIP was blended with PHB to improve the mechanical properties of PHB. With the change in the blend composition there is virtually no shift in the Tg of either PIP or PHB, indicating that PHB and PIP are immiscible [14].

Gassner and Owen studied the thermal and mechanical properties of PHBV/EVA blends [15], which are immiscible. The blends of PHBV and EVA show two Tg, corresponding to the two components. The mechanical properties of PHBV/EVA blends are strongly affected by blend composition. It is reported that only the PHBV component showed degradation in soil degradation tests.

However, miscible, biodegradable polymer/rubber blends are essential to modify the properties of biodegradable polymers. Until now, only a few miscible bacterial polyester/rubber blends are found, including PHB/EVA (85 wt% vinyl acetate) [13], PHB/PECH [16-19], PHB/PECH-co-EO and PHBV/PECH-co-EO blends [20].

PECH is a linear and amorphous elastomer with a low Tg of around −23 °C. PECH-co-EO is an epichlorohydrin copolymer rubber. Both PECH and PECH-co-EO are reported to be miscible with crystalline polyethylene oxide (PEO) [21]. The two polymers have also been blended with PHB and PHBV to form miscible crystalline/amorphous polymer blends. PECH is also reported to be completely miscible with PHB and PEO in the ternary blends of PHB/PECH/PEO. [22]

Goh and co-workers reported that PHB and PHBV were miscible with PECH-co-EO using DSC and polarising optical microscopy (POM) [20]. Both PHB/PECH-co-EO and PHBV(14 mol% hydroxyl valerate (HV) content)/PECH-co-EO blends were prepared by casting them from chloroform solutions. The miscibility of the blends was investigated by DSC. The Tg of the blends are composition dependent and intermediate between those of the component polymers, indicating that both PHB/PECH-co-EO and PHBV/PECH-co-EO blends are miscible. Furthermore, the Tg composition curves of both blends can be fi tted by the Kwei equation [5]. The composition dependent Tg fi ts the Kwei equation well, using k = 1.0 and q = -16.3 in PHB/PECH-co-EO blends. Similarly,


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the composition dependent Tg of PHBV/PECH-co-EO blends can also be described well by the same equation using k = 1.05 and q = 15.1. It can also be found that the melting point temperature of PHB or PHBV decreases slightly with the increase of the PECH-co-EO content. Furthermore, the cold crystallisation temperature (Tcc) of PHB is lower than that of pure PHB at lower PECH-co-EO content, and is higher than that of pure PHB at higher PECH-co-EO content in the PHB/PECH-co-EO blends. This indicates that the addition of PECH-co-EO may have some positive effects on the crystallisation of PHB when the PECH-co-EO content is lower. On the other hand, for the PHBV/PECH-co-EO blends, the Tcc of PHBV increases with the increase of the PECH-co-EO content, indicating that the cold crystallisation of PHBV is suppressed by the addition of the amorphous diluent PECH-co-EO in the blends.

For PHB/PECH-co-EO and PHBV/PECH-co-EO blends, both PHB and PHBV are biodegradable crystalline polymers, while PECH-co-EO is an amorphous component. It is well known that the depression of the melting point of a crystalline polymer blended with an amorphous polymer provides important information about its miscibility and its associated polymer-polymer interaction parameter. The melting behaviour of PHB/PECH-co-EO and PHBV/PECH-co-EO blends were studied by DSC. The equilibrium melting point is obtained for both polymer blends using the Hoffman-Weeks method. It was found that the Tm

o of both PHB and PHBV decreases with the increase of the amorphous diluent PECH-co-EO in the blends. The equilibrium melting point data obtained were analysed by the Nishi-Wang equation [7] based on the Flory-Huggins theory [6]. For PHB/PECH-co-EO blends, a negative value χ12 = −0.089 was obtained using V1 = 54.2 cm3/mol, V2 = 75 cm3/mol, ΔHo = 1.25 × 104 J/mol in Equation 7.6. As for the PHBV/PECH-co-EO blends, a negative value χ12 = −0.075 was obtained using V1

= 54.2 cm3/mol, V2 = 75 cm3/mol, ΔHo = 109 J/g in Equation 7.6. The negative polymer-polymer interaction parameters for both polymer blends, indicate that both polymer blends are thermodynamically miscible in the melt. On the other hand, as compared to a slightly more negative interaction parameter for PHB/PECH-co-EO blends, it can be concluded that the interaction between PHB and PECH-co-EO is suppressed by the presence of HV units in the case of PHBV/PECH-co-EO blends. Therefore, the miscibility of PHB/PECH-co-EO and PHBV/PECH-co-EO blends has been confi rmed not only from the single composition dependent Tg but also from the negative polymer-polymer interaction parameters.

The spherulitc morphology and growth for both PHB/PECH-co-EO and PHBV/PECH-co-EO blends were studied using POM. Typical spherulitic textures for various samples are observed when they crystallise at various crystallisation temperatures isothermally. Volume-fi lling spherulites for PHB or PHBV were obtained in the blends, indicating that the uncrystallisable component was rejected into the interlamellar or interfi brillar regions. Furthermore, the spherulitic growth rate (G) of PHB or PHBV showed a bell curve as a function of crystallisation temperature. The value of G decreased with the increase of PECH-co-EO content for a given crystallisation temperature in the blends. Meanwhile, the temperature showing the maximum values of G of PHB and PHBV


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spherulites also shifted to the low temperature range with the addition of PECH-co-EO in the blends.

Spherulitic growth kinetics can usually be analysed using the Lauritzen-Hoffman equation [23]:

logG +

U*

2.3R(Tc T )= log Go

Kg

2.3Tc( T)f (7.7)

where Go is a pre-exponential factor, U* is the transport activation energy, T∞ is a hypothetical temperature below which all viscous fl ow ceases, Kg is a nucleation parameter, T is the degree of supercooling, and f is a correction factor to account for the variation in the bulk enthalpy of fusion per unit volume with temperature, f = 2Tc /(Tm

o+ Tc). It was found that the value of Kg decreases with the increase of PECH-co-EO content for both polymer blends.

PHB/PECH blends were prepared by solution casting from dichloromethane. PHB/PECH blends showed a single Tg which fi ts the Fox equation well, indicating that both components are miscible. The polymer-polymer interaction parameter 12 = −0.054 was obtained from the Nishi-Wang equation, indicating that both components are thermodynamically miscible [16]. Therefore, the miscibility of PHB/PECH blends was confi rmed both from the single composition dependent Tg and from the negative polymer-polymer interaction parameter. Similar results were also found in PHB/PECH-co-EO and PHBV/PECH-co-EO blends [20]. As mentioned previously, the polymer-polymer interaction parameter of PHB/PECH-co-EO blends was –0.089, indicating that the introduction of the EO segment into the backbone of PECH increases the polymer-polymer interaction with PHB. The spherulitic growth rates of PHB/PECH blends were also studied using POM, which decreased with the increase of the PECH component at a given crystallisation temperature. It is also concluded that the PECH component was rejected into the interlamellar or interfi brillar regions of the PHB spherulites since the volume was space fi lling. Furthermore, the small angle X-ray scattering (SAXS) studies on the PHB/PECH blends provided more information about the localisation of the amorphous component in the spherulitic structure of the crystalline polymer [17]. The results showed that the PECH component was dispersed at the molecular level in the interfi brillar region. The addition of the PECH component to PHB also causes a reduction of the overall crystallisation rate. It is expected that the addition of PECH to PHB must also have a remarkable infl uence on the degradation of PHB in the blends. Doi and co-workers found that the biodegradability and the tensile properties of PHB were improved markedly by blending with PECH [24].


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7.3 Morphology and Crystallisation of Polyamide/Rubber Polymer Blends

Thermoplastic elastomers, which are prepared by mixing elastomers with thermoplastics, are of signifi cant interest. The essential feature of all TPE is two polymeric phases, where one is rubbery and the other is either glassy or crystalline. These materials combine the processability of a thermoplastic with the functional performance of a rubber. Such materials are tending to replace traditional rubber in a host of applications offering an easier or cheaper route for producing products which can also be easily recycled. As a result, TPE have started to be used instead of thermoset crosslinked rubber in many applications. There are seven main TPE groups found commercially – styrene block copolymers, impact modifi ed and super soft PP, polyolefi nic blends of PP and crosslinked EPDM, thermoplastic PU, melt processible rubber, thermoplastic copolyesters, and thermoplastic PA.

Among the previously mentioned TPE, elastomeric PA offer the best heat resistance in TPE and also have good chemical resistance so that its applications are cable jacketing and aeronautics. Many researchers have reported the preparation, characterisation, mechanical properties, thermal properties, morphology and crystallisation behaviour of PA (or Nylon) and rubber (elastomer) blends [25-33].

Paul and co-workers prepared blends of PA 6 and EPR grafted with maleic anhydride (EPR-g-MA) using a melt blending process [25, 26]. Two different rubbers were used. One was EPR-g-MA, which was nearly free of crystallinity. The other had a higher level of ethylene crystallinity, and was called H-EPR-g-MA. It was found that the reaction of the PA amine end groups with the grafted MA had the potential to form TPE with controlled morphology and chemical bonding between the phases. They further studied the effects of PA 6 content and crystallinity of the maleated rubber on morphological, thermal and mechanical properties of these blends. The morphology of blends of both EPR-g-MA and H-EPR-g-MA with PA 6 was studied by transmission electron microscopy (TEM) over the whole composition (100/0 to 0/100). Both blends showed similar morphology variation trends. Generally, discrete particles of the minor phase in a matrix of the major phase were observed at 20 and 80% of PA 6. Meanwhile, a tendency for co-continuity was observed for the intermediate compositions containing 40-60% of PA 6. It was found that an elongated PA 6 phase was observed at a level of 40% of PA 6, and this was more obvious at 50% PA 6. Furthermore, complete phase inversion was found at 60% of PA 6. The rubber existed as a dispersed phase within the PA 6 matrix. The morphology studies showed that the phase inversion composition was about 50% of PA 6 for both PA 6 and rubber blends. But there were some morphological differences between EPR-g-MA and H-EPR-g-MA in the blends. The PA particles were found to be smaller in the matrix of H-EPR-g-MA than those of EPR-g-MA at 20% of PA 6 because the melt viscosity of H-EPR-g-MA was higher than that of EPR-g-MA. Furthermore, the EPR-g-MA phase showed a more elongated structure than H-EPR-g-MA for blends


261


with intermediate compositions. Thermal and mechanical properties of PA 6 and two different rubber blends were also investigated using DSC and DMTA. Both polymer blends showed a melting peak at around 217 °C, corresponding to the melting peak of PA 6. However, they showed very different melting behaviour at a low temperature range. A melting peak at 45 °C was found for H-EPR-g-MA phase with a high heat of fusion, while a melting peak at 125 °C was found for EPR-g-MA with a low heat of fusion. The previously described, different melting behaviour was believed to result from the difference in the crystallinity formed from the sequences of ethylene units in the rubber. The melting peaks for PA 6 and the rubber did not change signifi cantly in the blends. The mechanical properties of both blends were studied by DMTA. The Tg of both the rubber and the PA 6 phase were also obtained from the locations of the tan δ peaks with DMTA. Both blends showed two Tg corresponding to the Tg of rubber and that of PA 6. Both polymer blends showed similar trends. The Tg of the EPR-g-MA rubber phase increased from −38.5 to −34.3 °C when the PA content increased from 0 to 40%. For H-EPR-g-MA, the Tg of the rubber phase increased from −23 to −18.1 °C when the PA content increased from 0 to 40%. However, for both blends, the Tg of the dispersed rubber phase decreases below that of neat rubber when the PA content increases further from 50 to 80% and becomes the major matrix phase. It should also be noted that the values of Tg for the rubber phase of H-EPR-g-MA blends were higher than those of EPR-g-MA blends due to the higher crystallinity of H-EPR-g-MA. On the other hand, the Tg of PA 6 decreased from 65.8 to 52.3 °C in the blends with EPR-g-MA when the rubber phase increased from 0 to 60%, while for PA 6 and H-EPR-g-MA blends, the Tg of PA 6 decreased from 65.8 to only 61.6 °C when the rubber content increased from 0 to 60%. In the blends with 80% rubber, the Tg of PA 6 could not be observed. Moreover, strain-hardening and cold-drawing were observed for both blend systems in the intermediate and PA-rich composition range. Modulus values from stress-strain measurements and dynamic, mechanical, thermal measurements were compared with predictions using a model by Hill for composite materials. Blends based on rubber with high ethylene crystallinity gave better agreement with the model than those based on amorphous rubber.

Groeninckx and co-workers studied the melt rheology and morphology of PA-6/EPR blends as a function of composition, temperature, and compatibiliser loading [27]. It was found that incompatibilised blends with a higher PA 6 content (or rubber content) had viscosities approximately intermediate between those of the component polymers. EPR was found to be dispersed as spherical inclusions in the PA matrix up to 30 wt% of its concentration. A kind of co-continuous morphology was observed between 30 and 50 wt% PA and a phase inversion beyond 70 wt% PA. Experiments were also carried out on in situ compatibilisation using MA-modifi ed EPR (EPR-g-MA), which reacted with the amino end groups of PA. This reaction produced a graft copolymer at the blend interface as a compatibiliser. When a few percent of modifi ed EPR were added, the viscosity of the blend was found to increase, however, the viscosity levelled off at higher concentrations, indicating a high level of interaction at the interface.


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Further morphological investigations showed that the size of the dispersed phase initially decreased when a few percent of the graft copolymer was added followed by a clear levelling off at higher concentration.

Zhang and co-workers studied the effect of several compatibilisers on mechanical properties and morphology of EPDM/PA copolymer (PA) blends [28]. A signifi cant reduction of the dispersed phase dimension was observed when chlorinated polyethylene (CPE) was added to an EPDM/PA blend, due to the interaction that exists between CPE and PA. A speculative description of confi guration was proposed to interpret the morphological investigation made on these blends based on DSC, DMTA, scanning electron microscopy and TEM characterisation. The studies of mechanical properties showed that the materials obtained possess useful strength and excellent heat resistance.

It should be noted that the previously mentioned PA/rubber blends are also called thermoplastic vulcanisates (TPV), which combine the excellent processing characteristics of thermoplastics with the elastic properties of elastomers. However, the elastomer is prepared by static vulcanisation. Recently, dynamic vulcanisation, the process of vulcanising elastomer during melt-mixing with molten thermoplastic, has become the best way to prepare TPV combining the excellent processing characteristics of thermoplastics with the elastic properties of elastomers. The elastic properties of TPV are similar to the more conventional class of thermoplastic elastomers based on hard segment and soft segment block copolymers. This technology has led to a signifi cant number of new thermoplastic elastomeric products commercialised during the mid-to-late 1980s. Some thermoplastic elastomers, through dynamic vulcanization, have been commercialised. A well-known commercial example of dynamically vulcanised thermoplastic elastomer compositions are blends of PP and EPDM. These TPV are prepared by fi rst melt-mixing PP with EPDM where a kind of co-continuous phase morphology is formed. Subsequently, a vulcanising agent is added to crosslink the EPDM rubber phase, and the crosslinked rubber phase will be sheared into small particles and fi nally fi nely dispersed in the thermoplastic matrix. Researches on the dynamic vulcanisation of blends of nitrile elastomer with PA and blends of CPE rubber with PA have also been reported [29, 30].

However, there are few publications available on dynamic vulcanisation of EPDM/PA blends. The diffi culty in preparing the EPDM/PA TPV is due to the high interfacial energy between the two components. For PA 6/rubber blends with 50 or 60 wt% of an EPDM rubber, Groeninckx and co-workers found that the blends exhibited co-continuous morphologies and thereby relatively poor mechanical properties [31]. Therefore, in order to improve the mechanical properties, they used a suitable compatibiliser and slightly crosslinked the rubber phase during melt-mixing. Through this process, it was possible to disperse up to 60 wt% rubber in the PA matrix and to improve the mechanical properties markedly. Such materials exhibited good elastic properties with


263


a thermoplastic processability. They investigated the infl uence of the compatibiliser, the crosslinking agent and the viscosity ratio of rubber/thermoplastic on the blend phase morphology using TEM. The viscosity of the PA phase should be low enough to shift the phase inversion towards higher rubber content. On the other hand, if the viscosity of the PA was too low, a kind of coarse blend morphology was achieved resulting in poor mechanical properties. Therefore, the viscosity ratio of rubber/PA played a crucial role in order to achieve a PA6/rubber TPV with fi ne rubber dispersion.

Huang and co-workers prepared EPDM/PA copolymer TPV using suitable compatibilisers and dynamic vulcanisation [32, 33]. In EPDM/PA copolymer TPV, the rubber/thermoplastics ratio was chosen to be 65/35. Therefore, PA was the minor component and formed the dispersed phase in an EPDM matrix for a 65/35 EPDM/PA blend. During dynamic vulcanisation of such a blend, EPDM and PA had to undergo a phase inversion to maintain the thermoplasticity of the blend. The general way for preparing TPV was to blend rubber with a thermoplastic through melt-mixing. Thus, a co-continuous phase was generated. A vulcanisation agent was then added to crosslink the rubber phase, which was no longer able to coalesce into a continuous phase. As the degree of crosslinking advanced during mixing, the continuous rubber phase became elongated further and then broke up into polymer droplets. Therefore, the thermoplastic became the continuous phase, while the rubber became the dispersed phase. Thus TPV were obtained with fi nely dispersed rubber particles in the thermoplastics matrix. However, Groeninckx and co-workers reported in their studies of the dynamic vulcanisation of PA 6/EPDM blends, that the phase inversion region could not be shifted due to the compatibilisation reaction between EPDM-g-MA and PA 6 [31]. In their study, peroxide was premixed with the rubber before the PA 6/EPDM blend was compounded. Huang and co-workers [32] prepared the EPDM/PA copolymers TPV in a different way. EPDM/PA copolymers were prepared as follows: a PA copolymer (a copolymer of 70% PA 1010, 20% PA 66 and 10% PA 6) with a compatibiliser (EPDM-g-MA or EPR-g-MA) was fi rst melt-mixed, and then EPDM was melt-mixed. Later, a sulfur vulcanizing agent and a coagent were added to crosslink the rubber phase. There were different variables that may affect the properties of EPDM/PA blends: the ratio of EPDM to PA, the volume fraction of compatibiliser, the molecular weight and the viscosity of EPDM and PA, composition and functionality of the compatibiliser and the crystalline structure of these systems. Huang and co-workers [33] studied the effect of dynamic vulcanisation on the morphology and crystallisation behaviour of EPDM/PA copolymer TPV in more details.

It should be noted that the morphology study for TPV is very important to confi rm the phase inversion. For TPV, the phase morphology of the blends is usually observed with TEM. Recently, atomic force microscopy (AFM) has been recognised as a powerful surface characterisation technique and has been widely used to study surface morphology of homopolymers, block copolymers, and polymer blends. AFM does not require the staining of the sample as TEM does. Therefore, it has become a useful and convenient tool to study the phase morphology for TPV.


264


Figure 7.1 shows the phase morphology for the dynamically vulcanised EPDM/PA copolymers sample, which was obtained by tapping mode AFM. The morphology of large crosslinked white EPDM particles was observed to be dispersed in the dark PA matrix. The large rubber particles had a tendency to be co-continuous since no compatibiliser was used. EPDM-g-MA and EPR-g-MA were used as compatibilisers in EPDM/PA copolymer TPV. The effects of the variety and the content of the compatibiliser on the phase morphology of the EPDM/PA copolymer TPV had also been investigated by AFM.

Figure 7.2 shows the phase morphology of EPDM/EPDM-g-MA/PA copolymer TPV with different amounts of EPDM-g-MA as a compatibiliser. The bright area represented the crosslinked rubber phase, and the dark area was the PA phase. It could be observed from the AFM images shown in Figure 7.2 that the rubber phase could be fi nely dispersed in the PA phase by increasing the content of the EPDM-g-MA, indicating that phase inversion existed due to the introduction of EPDM-g-MA as a compatibiliser. These results indicated that the phase inversion region could be shifted although there was a compatibilisation reaction between EPDM-g-MA and PA. It can be explained as follows: PA kept its fl ow mobility in the molten state. Some PA droplets may be encapsulated by the rubber phase during the dynamic vulcanisation process because of the steric stabilisation effect of the EPDM-g-PA formed. The sub-inclusions of PA inside the rubber particles could be observed in the AFM images. However, the PA phase tended to coalesce and could form the continuous phase because the degree of crosslinking of the rubber phase was not high enough to form a network and the PA phase still kept the fl ow mobility in the molten state. Furthermore, the size of the crosslinked rubber particles decreased with increasing amount of EPDM-g-MA content due to the compatibilisation reaction. When EPDM was totally replaced by EPDM-g-MA, the size of the rubber particles were at their smallest.

Figure 7.1 AFM image of dynamically vulcanised EPDM/PA (65/35)


265


Figure 7.2 AFM images of dynamically vulcanised EPDM/EPDM-g-MA/PA TPV. (a) 58.5/6.5/35; (b) 52/13/35; (c) 39/26/35; (d) 13/52/35; (e) 0/65/35

(a)

(c)

(e)

(b)

(d)


266


The effect of EPR-g-MA as a compatibiliser on the phase morphology of EPDM/PA copolymer TPV has also been studied using AFM. Figure 7.3 shows the phase morphology of EPDM/EPR-g-MA/PA copolymer TPV with different contents of EPR-g-MA as a compatibiliser. Since EPR-g-MA could not be crosslinked by sulfur, three phases existed in these TPV. It is observed in Figure 7.3a and 7.3b that the bright crosslinked EPDM particles were dispersed in the dark matrix of PA and EPR-g-MA. It should be noted that the morphology of TPV using EPR-g-MA as the compatibiliser showed much smaller dispersed rubber particles compared with morphology of TPV using the same content of EPDM-g-MA as the compatibiliser. The small dispersed rubber particles were expected to have a positive effect on the mechanical properties of TPV. When the EPR-g-MA content was increased to 26 wt%, it could be seen from Figure 7.3c that the size of the crosslinked EPDM particles decreased and the larger bright regions seemed to be composed of small particles, which was more obvious in the enlarged image inside Figure 7.3c. The uncrosslinked EPR-g-MA at the interface between the crosslinked EPDM and PA tended to coalesce with increasing compatibiliser content of EPR-g-MA. When the EPR-g-MA content was 52 wt%, the uncrosslinked EPR-g-MA became the dark continuous phase, and PA and crosslinked EPDM became the bright dispersed phase as shown in Figure 7.3d. When EPDM was totally replaced by EPR-g-MA, the blend was composed of PA and EPR-g-MA only. Since EPR-g-MA could not be vulcanised by sulfur, the bright PA particles were found to be dispersed into the dark matrix of EPR-g-MA.

The crystallisation behaviour of EPDM/PA copolymer TPV had also been studied using DSC. Five different samples were used including neat PA copolymer (PA), unvulcanised EPDM/PA (65/35), vulcanised EPDM/PA (65/35), vulcanised EPDM/EPDM-g-MA/PA (52/13/35), and vulcanised EPDM/EPR-g-MA/PA (52/13/35). There were two reasons for selecting the fi ve different samples. One was to study the effect of vulcanisation on the crystallisation behaviour of PA of binary EPDM/PA blends. The other was to investigate the effect of different compatibilisers (EPDM-g-MA and EPR-g-MA) on the crystallisation behaviour of PA in the ternary TPV. Both nonisothermal crystallisation and isothermal crystallisation processes were used to fulfi ll these two purposes.

For the nonisothermal crystallisation process, all the samples were cooled from the crystal-free melt at a cooling rate of 5 °C/minute. Figure 7.4 shows the DSC traces for the fi ve different samples. Neat PA copolymer exhibited a broad crystallisation exotherm with the peak at around 100 °C. The crystallisation peak temperature of PA copolymer almost did not change in the binary blends of EPDM/PA, irrespective of the vulcanisation. But it was found that the crystallisation enthalpy of the PA phase decreased in the case of vulcanised EPDM/PA compared with that of unvulcanised EPDM/PA blend. However, the crystallisation peak temperature shifted to 110.5 and 107.5 °C, for EPDM/EPR-g-MA/PA and EPDM/EPDM-g-MA/PA TPV, respectively, indicating that the introduction of the compatibilisers EPR-g-MA and EPDM-g-MA had had a positive effect on the crystallisation of the PA phase. This phenomenon could be ascribed to the


267


Figure 7.3 AFM images of dynamically vulcanised EPDM/EPR-g-MA/PA TPV. (a) 58.5/6.5/35; (b) 52/13/35; (c) 39/26/35; (d) 13/52/35; (e) 0/65/35

(a)

(c)

(e)

(b)

(d)


268


fact that addition of compatibiliser would help to form fi ne dispersion of crosslinked rubber particles, which may subsequently act as heterogeneous nucleating centres for the crystallisation of the PA phase. The Tc and the crystallisation enthalpy ( H) of the PA phase for the fi ve different samples are summarised in Table 7.1.

Furthermore, the effect of the content of the compatibilisers EPR-g-MA and EPDM-g-MA on the crystallisation peak temperature of the PA phase in the TPV has also been studied. For both TPV with different compatibiliser, the Tc increased and fi nally arrived at a constant value with increasing compatibiliser content as shown in Figure 7.5. The values of H of the PA phase in the TPV were less than that of neat PA. The decrease of crystallinity could be ascribed to the crystallisation in confi ned spaces between rubber

Table 7.1 Crystallisation peak temperature and enthalpy of crystallisation of PA phase for the fi ve different samples during nonisothermal

crystallisation at a cooling rate of 5 oC/min Tc (°C) H (J/g)

PA 41.2 41.2EPDM/PA (unvulcanised) (65/35) 39.4 39.4EPDM/PA (65/35) 31.1 31.1EPDM/EPR-g-MAH/PA (52/13/35) 39.1 39.1EPDM/EPDM-g-MAH/PA (52/13/35) 39.1 39.1

Figure 7.4 DSC cooling traces at a cooling rate of −5 oC/minutes: (a) PA; (b) EPDM/PA (unvulcanised) (65/35); (c) EPDM/PA (65/35); (d) EPDM/EPR-g-MAH/PA (52/13/35); (e)

EPDM/EPDM-g-MAH/PA (52/13/35)


269


particles. Figure 7.5 shows that the value of H of the PA phase went through a maximum value at a compatibiliser content of 20% in both TPV. The addition of compatibiliser may reduce the particle size of the dispersed rubber phase which was benefi cial for nucleation of PA, but at the same time introduced the compatibilisation reaction between compatibiliser and PA which would restrict the growth of PA crystallinity. Therefore, the value of H of the PA phase went through a maximum with increasing compatibiliser content.

Figure 7.5 Effect of compatibiliser content on Tc and H in EPDM/PA TPV (EPDM+compatibiliser/PA: 65/35)


270


Apart from the nonisothermal crystallisation behaviour, the isothermal crystallisation of the fi ve different samples was also been investigated by DSC. The isothermal crystallisation temperature was chosen in the range of 100-108 °C. The well-known Avrami equation is often used to analyse the isothermal crystallisation kinetics – it assumes that the relative degree of crystallinity develops with crystallisation time t as:

1 − Xt = exp (− k t n) (7.8)

where n is the Avrami exponent depending on the nature of nucleation and growth geometry of the crystals, and k is a composite rate constant involving both nucleation and growth rate parameters [34]. The Avrami parameters n and k can be obtained from the plots of log(-ln(1-Xt)) versus log t. The Avrami exponents n and crystallisation rate constants k of the fi ve different samples at 100-108 °C are listed in Table 7.2. It was found

Table 7.2 Avrami parameters for the fi ve samples at various crystallisation temperature

Sample Tc (oC) n k (s-n)

Neat PA100

2.27.58 × 10

102 6.76 × 10104 5.88 × 10106 5.24 × 10108 4.36 × 10

EPDM/PA (unvulcanised) (65/35) 100 2.3 1.73 × 10102 1.31 × 10104 1.09 × 10106 8.51 × 10108 6.76 × 10

EPDM/PA(65/35)

100 2.5 3.71 × 10102 2.63 × 10104 2.18 × 10106 1.44 × 10108 1.07 × 10

EPDM/EPDM-g-MAH/PA (52/13/35) 100 2.5 6.02 × 10102 5.24 × 10104 4.46 × 10106 3.46 × 10108 2.57 × 10

EPDM/EPR-g-MAH/PA (52/13/35) 100 2.4 3.16 × 10-5

102 1.94 × 10-5

104 1.51 × 10-5

106 8.91 × 10-6

108 6.16 × 10-6


271


that the value of the average n did not change signifi cantly despite the sample and the crystallisation temperature. For the neat PA copolymer, the value was around 2.2, which increased a little to 2.3-2.5 in the blends with EPDM despite the vulcanisation and the compatibiliser, indicating that the crystallisation mechanism of the PA phase remained the same. On the other hand, the overall crystallisation rate k was found to be dependent on the crystallisation temperature. The value of k decreased with increasing the crystallisation temperature for the same sample. Furthermore, it was also found that the value of k changed with the different samples used in this study, which is shown in Figure 7.6.

Figure 7.6 Plots of half crystallisation time versus Tc (a) and plots of crystallisation rate t0.5

-1 versus Tc (b) for various PA samples


272


The half-life crystallisation time t0.5, the time required to achieve 50% of the fi nal crystallinity of the samples, is an important parameter for the discussion of crystallisation kinetics. The value of t0.5 can be calculated by the following equation based on the Avrami equation:

t0.5 = (ln2/k) 1/n (7.9)

where k and n are the same as in the Avrami equation. Figure 7.6a shows the plots of the values of t0.5 for the fi ve different samples as a function of crystallisation temperatures. It could be concluded from Figure 7.6a that the crystallisation time increased with increasing the crystallisation temperature. Meanwhile, the variety of the samples also had a signifi cant infl uence on the crystallisation time. For a given crystallisation temperature, neat PA copolymer needed the longest crystallisation time, and EPDM/EPR-g-MA/PA required the shortest crystallisation time. Usually, the crystallisation rate was described as the reciprocal of t0.5. Therefore, the values of t0.5

-1 were plotted as a function of crystallisation temperature in Figure 7.6b for all the samples to compare the effect of the variety of the sample on the crystallisation rate. From Figure 7.6b, the crystallisation rate was found to decrease with increasing the crystallisation temperature for each sample. Meanwhile, the crystallisation rate showed the following order among the fi ve samples at a given crystallisation temperature, PA < unvulcanised EPDM/PA < vulcanised EPDM/PA < EPDM/EPDM-g-MA/PA < EPDM/EPR-g-MA/PA. The crystallisation rate was highest in compatibilised TPV and lowest in neat PA while intermediate in uncompatibilised TPV and unvulcanised EPDM/PA blends. Comparing with unvulcanised EPDM/PA blends with the vulcanised EPDM/PA blends, it seemed that the dynamic vulcanisation process could increase the crystallisation rate for the PA copolymer obviously, especially when a suitable compatibiliser is used. The dynamic vulcanisation introduced fi ne crosslinked rubber particles which could act as heterogeneous nucleating centres, which would enter the PA matrix and cause an increase in nucleating rate. For unvulcanised EPDM/PA blends, PA copolymer was the island phase dispersed in the unvulcanised rubber matrix. In this case, the crystallisation of PA would be restricted. While for TPV, the immiscibility was improved and the crystallisation of PA was increased accordingly. Without any compatibiliser, it would be diffi cult to disperse a rubber phase in a PA matrix during dynamic vulcanisation and the crosslinked rubber particles were large (as shown in Figure 7.1). Using EPDM-g-maleic anhydride (MAH) or EPR-g-MAH as compatibiliser, the corresponding succinic anhydride groups could react readily with the terminal amine group of the PA copolymer leading to a graft copolymer formed in situ. The copolymer acted as a macromolecular surfactant and its presence, during mixing, permits the formation of very small droplets of the elastomer which later, during dynamic vulcanisation, became very small particles of vulcanised rubber. These fi nely dispersed vulcanised rubber particles increased the density of nucleating centres and thus increased the nucleating rate. On the other hand, the compatibilisation reaction restricted the mobility of PA copolymer chains and decreased the crystal growing rate. However, Figure 7.6 indicates that the nucleating rate was the main infl uencing factor and the crystallisation rate increased by adding compatibiliser. Compared


273


with the morphology of TPV with the same dosage of EPDM-g-MAH compatibiliser (Figure 7.2b), the morphology of TPV using EPR-g-MAH as a compatibiliser showed much smaller dispersed rubber particles (Figure 7.3b), which may contribute to the higher crystallisation rate. Therefore, the preparation of TPV through the dynamic vulcanisation was an effi cient route to produce plastic and rubber blends with suitable properties. The development of plastic/rubber blends will fi nd more and more application in the near future in industry and improve the quality of modern lives.

Acknowledgement

The authors thank Dr. Hua Huang for her kind contributions to this chapter.

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