**Plasticization and Morphology Development in Dynamically Vulcanized Thermoplastic Elastomers**

Shant Shahbikian and Pierre J. Carreau

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/61414

## **Abstract**

Dynamically vulcanized thermoplastic elastomers constitute one of the main catego‐ ries among various types of thermoplastic elastomers (TPEs). Due to the commercial importance of this particular group of TPEs, tremendous efforts have been dedicated to improve the understanding and control the phase morphology development. The ultimate goal is to obtain materials with improved physical and mechanical proper‐ ties. As in other polymeric compounds, the parameters during the mixing stage have a significant influence on the final morphology of dynamically vulcanized blends. Furthermore, the phase morphology and, therefore, the distribution of elastomeric do‐ mains in the thermoplastic phase are also strongly dependent on the formulation. This chapter discusses the main important processing factors and, more specifically, high‐ lights the effects of plasticization and curing on the morphology development of dy‐ namically vulcanized thermoplastic elastomer blends. The following text provides fundamental information on how one should take into consideration each parameter affecting the morphology of nonreactive and reactive elastomer/thermoplastic blends.

**Keywords:** Thermoplastic elastomer, dynamic vulcanization, plasticization, morphol‐ ogy, rheological properties

## **1. Introduction**

Thermoplastic elastomers (TPEs) represent a large group of polymeric materials that are melt processable, similar to regular thermoplastics, and they exhibit rubber-like behavior identical to that of cross-linked elastomers. This special characteristic of TPEs is mainly due to the presence of thermoreversible cross-links, which are broken during the melt processing step under high shear and elevated temperature, and formed once again when melt processing is over and the compound reaches ambient temperature. The main concept behind the thermo‐

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plastic elastomers and, therefore, thermoreversible cross-links is the simultaneous presence of phase-separated hard and soft segments in the compound. Although both phases contribute to the overall physical and mechanical properties of the final product, some specific properties are usually associated with one phase or the other. Below the melting temperature of the hard phase, this phase usually provides the strength, stiffness, and chemical resistance of the material. On the other hand, the soft one acts as an elastomer providing the flexibility and the elastic nature by controlling the hardness, the compression, and the tensile sets. Furthermore, the soft phase dictates the lower service temperature limit of the product.

Commercially available TPEs can be generally classified into the following four groups presented in Table 1.


SBS: Poly(styrene-butadiene-styrene); SEBS: Poly(styrene-ethylene-butadiene-styrene); TPU: Thermoplastic polyure‐ thene; COPE: Copolyether-ester elastomers; COPA: Copolyamide elastomers; EPR: Ethylene propylene rubber; EPDM: Ethylene propylene terpolymer rubber; PP: Polypropylene; NBR: Nitrile butadiene rubber; PVC: Polyvinyl chloride; IIR: Butyl rubber (isobutylene-isoprene rubber).

**Table 1.** Classification of thermoplastic elastomers (TPEs) along with few examples.

According to the general categories of TPEs shown in Table 1, the aforementioned properties associated with TPEs could be obtained through numerous paths. In the first three groups, the material mainly consists of blocks or grafted segments of soft and hard constituents through polymerization [1,2]. In addition to these large varieties of TPEs obtained through polymeri‐ zation, mechanical blending of conventional elastomers with plastics provides another accessible route toward the production of TPEs. In this approach, the polymeric constituents are blended in a conventional melt mixing equipment, such as a twin-screw extruder. The result is often an immiscible blend, particularly due to the fact that the majority of polymeric pairs of large molecular weight and low interactive forces tend to form phase-separated microstructure. Interestingly, the phase-separated nature of these blends is what it is needed to exhibit properties associated with TPEs.

Although some simple nonreactive elastomer/thermoplastic blends, such as ethylene-propy‐ lene rubber/polypropylene (EPR/PP), have gained tremendous attention for their use in the automotive industry [3], the nonreactive blends have generally poor elastic recovery and poor hydrocarbon fluid resistance in comparison to their reactive counterparts [4]. Furthermore, the morphology of an immiscible simple blend is prone to change during the reprocessing and downstream operations. These issues may extremely affect the physical and mechanical properties of these types of TPEs. Consequently, simple blending of polymeric constituents is usually not sufficient to guarantee a permanent morphological feature and, therefore, stable mechanical properties regardless of the processing history of the blend. To overcome this issue, chemically cross-linking of polymeric phases is required in order to stabilize the morphology. In this case, the reactive blends of elastomer/thermoplastic blends are generally obtained through dynamic vulcanization (DV) where the elastomer phase is selectively cross-linked in the presence of a curing system. These blends are commonly known as thermoplastic vulcan‐ izates (TPVs). The origin of DV process comes from the work of Gessler and Haslett on polypropylene and chlorobutyl rubber, where the elastomer was cross-linked in the presence of zinc oxide [5]. Fischer [6-9] further pushed the boundaries of our understanding by working on dynamically and statically cross-linked elastomer blends with polyolefins. The work of Fisher on dynamic vulcanization was further extended by complementary investigations done by Monsanto [10-13]. statically cross-linked elastomer blends with polyolefins. The work of Fisher on dynamic vulcanization was further extended by complementary investigations done by Monsanto [10-13]. In both simple and dynamically vulcanized blends, the thermoplastic elastomer behavior is usually obtained for high concentration of the elastomer phase, e.g., higher than 50 wt%. This

plastic elastomers and, therefore, thermoreversible cross-links is the simultaneous presence of phase-separated hard and soft segments in the compound. Although both phases contribute to the overall physical and mechanical properties of the final product, some specific properties are usually associated with one phase or the other. Below the melting temperature of the hard phase, this phase usually provides the strength, stiffness, and chemical resistance of the material. On the other hand, the soft one acts as an elastomer providing the flexibility and the elastic nature by controlling the hardness, the compression, and the tensile sets. Furthermore,

Commercially available TPEs can be generally classified into the following four groups

**I II III IV**

SBS: Poly(styrene-butadiene-styrene); SEBS: Poly(styrene-ethylene-butadiene-styrene); TPU: Thermoplastic polyure‐ thene; COPE: Copolyether-ester elastomers; COPA: Copolyamide elastomers; EPR: Ethylene propylene rubber; EPDM: Ethylene propylene terpolymer rubber; PP: Polypropylene; NBR: Nitrile butadiene rubber; PVC: Polyvinyl chloride; IIR:

According to the general categories of TPEs shown in Table 1, the aforementioned properties associated with TPEs could be obtained through numerous paths. In the first three groups, the material mainly consists of blocks or grafted segments of soft and hard constituents through polymerization [1,2]. In addition to these large varieties of TPEs obtained through polymeri‐ zation, mechanical blending of conventional elastomers with plastics provides another accessible route toward the production of TPEs. In this approach, the polymeric constituents are blended in a conventional melt mixing equipment, such as a twin-screw extruder. The result is often an immiscible blend, particularly due to the fact that the majority of polymeric pairs of large molecular weight and low interactive forces tend to form phase-separated microstructure. Interestingly, the phase-separated nature of these blends is what it is needed

**Ionomers**

Ethylene-methacrylic acid Ethylene-acrylic acid Butadiene-acrylic acid and many more

**Blends of soft elastomer and hard thermoplastic**

*Simple (nonreactive) or dynamically vulcanized blends (DV)*

> EPDM/PP NBR/PP NBR/PVC IIR/PP and many more

the soft phase dictates the lower service temperature limit of the product.

**Random copolymers**

EPR Ethylene-α Olefin Propylene-α Olefin and many more

**Table 1.** Classification of thermoplastic elastomers (TPEs) along with few examples.

presented in Table 1.

50 Thermoplastic Elastomers - Synthesis and Applications

**Block copolymers**

*Triblock or segmented block copolymers:*

> SBS SEBS TPU COPE COPA and many more

Butyl rubber (isobutylene-isoprene rubber).

to exhibit properties associated with TPEs.

In both simple and dynamically vulcanized blends, the thermoplastic elastomer behavior is usually obtained for high concentration of the elastomer phase, e.g., higher than 50 wt%. This single argument on the concentration of the elastomeric component has a huge consequence on the morphology. The simple blends with high elastomer content tend to form a cocontinuous morphology, where both elastomer and thermoplastic constituents are intercon‐ nected throughout the whole bulk of the material (Fig. 1a). However, in dynamically vulcanized blends, the elastomer (usually the major and initially part of the co-continuous phase) becomes discontinuous and dispersed during dynamic vulcanization (Fig. 1b). Eventually, the thermoplastic becomes the continuous phase surrounding the cross-linked elastomer particles. This morphology transformation is known as *phase inversion*. single argument on the concentration of the elastomeric component has a huge consequence on the morphology. The simple blends with high elastomer content tend to form a co-continuous morphology, where both elastomer and thermoplastic constituents are interconnected throughout the whole bulk of the material (Fig. 1a). However, in dynamically vulcanized blends, the elastomer (usually the major and initially part of the co-continuous phase) becomes discontinuous and dispersed during dynamic vulcanization (Fig. 1b). Eventually, the thermoplastic becomes the continuous phase surrounding the cross-linked elastomer particles. This morphology transformation is known as *phase inversion*.

Fig. 1. AFM phase micrographs of: (a) nonreactive simple, and (b) dynamically vulcanized EPDM/PP 50/50 (wt/wt%) blends (from [14]) **Figure 1.** AFM phase micrographs of: (a) nonreactive simple, and (b) dynamically vulcanized EPDM/PP 50/50 (wt/wt %) blends (from [14])

By far, the importance of the phase-separated structure and the role of morphology have been clearly emphasized. For one familiar with polymer blending technology, the fine control of phase morphology even in a simple nonreactive blend composed of merely two components represents a huge challenge. Moreover, the presence of a complex flow field in industrial mixing equipment with simultaneous breakup and coalescence of the dispersed phase creates a far more complicated environment for comprehending the phase morphology development. The level of complexity may even increase when additional components such as processing aids, fillers, plasticizers, and curing system are added. This is commonly the case in the TPV industry where By far, the importance of the phase-separated structure and the role of morphology have been clearly emphasized. For one familiar with polymer blending technology, the fine control of phase morphology even in a simple nonreactive blend composed of merely two components represents a huge challenge. Moreover, the presence of a complex flow field in industrial mixing equipment with simultaneous breakup and coalescence of the dispersed phase creates a far more complicated environment for comprehending the phase morphology development. The level of complexity may even increase when additional components such as processing aids, fillers, plasticizers, and curing system are added. This is commonly the case in the TPV industry where the fine-tuning of the final properties is achieved by combining several different reactive and nonreactive additives.

## **2. Plasticization**

The use of plasticizers in both rubber and thermoplastic elastomer industries is a wellestablished technology [15]. Despite the overall processing cost reduction of the final product in the presence of a plasticizer, some technical aspects of a plasticized compound can also be improved. For instance, plasticizers have been used to improve the low temperature mechan‐ ical properties to reduce the hardness and acts as a dispersion aid for fillers and additives [16]. It further improves the resistance to oil swell, heat stability, hysteresis, permanent set, elastic recovery, as well as the melt processability and the final appearance of the compound [13].

Several theories such as gel [17,18], lubricity [19-21], and free volume [22] theories have so far been developed and further extended to explain different mechanisms involved during plasticization. In a rather general way, the lubricity theory considers that a plasticizer reduces the intermolecular friction between polymer chains, which is originally considered as its source of rigidity. It acts as lubricant and reduces the resistance to sliding between molecules. In the gel theory, the polymer molecules are considered to form a tridimensional structure held by loose attachments along their chains. According to this theory, the stiffness of the polymer is mainly due to the presence of this tridimensional structure. In such systems, plasticizers act in favor of reducing the number of attachments between polymer molecules and, therefore, enabling the molecules to change their conformation. In the free volume theory, the friction between polymer chains is attributed to the volume between molecules. The free volume in polymers is essentially considered as the required space for chain mobility. By increasing the temperature in a nonplasticized system, the chain mobility increases as a consequence of an increase in free volume. On the other hand, the shrinkage of the free volume with decreasing temperature may reach to critical level where only limited free space is available for polymer chains to have large segmental motion. This critical temperature is known as the glass transition temperature (*Tg*). Therefore, polymers below *Tg* behave as solid glassy materials, whereas at temperatures above *Tg* they possess rubber-like properties. According to this theory, the free volume in a polymer may be increased through different paths [23]: (1) by increasing the temperature (as mentioned earlier), (2) by lowering the molecular weight of the polymer resulting in an increased concentration of end groups, (3) by incorporation and/or increasing the length of side chains in the polymer, (4) by incorporating segments with low steric hindrance and low intermolecular interaction along the polymer chains, (5) by adding lower molecular weight compounds with lower *Tg*, which are compatible with the polymer. Consequently, a decrease in the glass transition temperature and an increase in the mobility of polymeric chains as a function of the plasticizer concentration could be readily associated with increases in the free volume.

By far, the importance of the phase-separated structure and the role of morphology have been clearly emphasized. For one familiar with polymer blending technology, the fine control of phase morphology even in a simple nonreactive blend composed of merely two components represents a huge challenge. Moreover, the presence of a complex flow field in industrial mixing equipment with simultaneous breakup and coalescence of the dispersed phase creates a far more complicated environment for comprehending the phase morphology development. The level of complexity may even increase when additional components such as processing aids, fillers, plasticizers, and curing system are added. This is commonly the case in the TPV industry where the fine-tuning of the final properties is achieved by combining several

The use of plasticizers in both rubber and thermoplastic elastomer industries is a wellestablished technology [15]. Despite the overall processing cost reduction of the final product in the presence of a plasticizer, some technical aspects of a plasticized compound can also be improved. For instance, plasticizers have been used to improve the low temperature mechan‐ ical properties to reduce the hardness and acts as a dispersion aid for fillers and additives [16]. It further improves the resistance to oil swell, heat stability, hysteresis, permanent set, elastic recovery, as well as the melt processability and the final appearance of the compound [13]. Several theories such as gel [17,18], lubricity [19-21], and free volume [22] theories have so far been developed and further extended to explain different mechanisms involved during plasticization. In a rather general way, the lubricity theory considers that a plasticizer reduces the intermolecular friction between polymer chains, which is originally considered as its source of rigidity. It acts as lubricant and reduces the resistance to sliding between molecules. In the gel theory, the polymer molecules are considered to form a tridimensional structure held by loose attachments along their chains. According to this theory, the stiffness of the polymer is mainly due to the presence of this tridimensional structure. In such systems, plasticizers act in favor of reducing the number of attachments between polymer molecules and, therefore, enabling the molecules to change their conformation. In the free volume theory, the friction between polymer chains is attributed to the volume between molecules. The free volume in polymers is essentially considered as the required space for chain mobility. By increasing the temperature in a nonplasticized system, the chain mobility increases as a consequence of an increase in free volume. On the other hand, the shrinkage of the free volume with decreasing temperature may reach to critical level where only limited free space is available for polymer chains to have large segmental motion. This critical temperature is known as the glass transition temperature (*Tg*). Therefore, polymers below *Tg* behave as solid glassy materials, whereas at temperatures above *Tg* they possess rubber-like properties. According to this theory, the free volume in a polymer may be increased through different paths [23]: (1) by increasing the temperature (as mentioned earlier), (2) by lowering the molecular weight of the polymer resulting in an increased concentration of end groups, (3) by incorporation and/or increasing the length of side chains in the polymer, (4) by incorporating segments with low

different reactive and nonreactive additives.

52 Thermoplastic Elastomers - Synthesis and Applications

**2. Plasticization**

Once a plasticizer is incorporated and completely dissolved at the molecular level in a desired polymer, the combination of all the aforementioned theories provides an extensive insight into the plasticization mechanisms. However, dissolution of a plasticizer into polymer and the compatibility may largely affect the efficiency of plasticization in both the short and the long term. The compatibility issue is the principal factor in determining the proper plasticizer for a given polymer. Generally, a compatible polymer/plasticizer pair is by nature a homogeneous mixture.

The incorporation of a plasticizer during an industrial compounding process is usually achieved by the expenses of mechanical energy. However, afterward the stability of the compound is directly related to the thermodynamic phase equilibrium between the compo‐ nents. As a result, the widely used compatibility concepts are, therefore, based on the theories of polymer solutions. In a pair of plasticizer/amorphous polymer with an upper critical solution temperature (UCST) phase diagram, the homogeneity of the mixture increases with temperature. At lower temperatures, a mixture containing around 15 to 40 vol% of plasticizer (a typical range in TPVs) may be phase-separated into polymer-rich and plasticizer-rich domains. In certain cases, a lower critical solution temperature (LCST) may also be observed. In contrast to UCST phase diagrams, those systems displaying LCST tend to phase-separate at higher temperature. The complexity of the phase diagram is normally increased when instead of an amorphous polymer; a semicrystalline polymer is to be plasticized. The crystal‐ lization process of the polymer tends to be the factor that complicates the formation of a homogeneous polymer/plasticizer mixture. At temperatures high enough to obtain a homo‐ geneous mixture between the amorphous fraction of the polymer and the plasticizer, but low enough that crystallinity still prevails, the presence of local crystal network affects negatively the compatibility and in certain conditions it results in migration of the plasticizer. This reasoning on the effect of crystallinity brings us to a general discussion regarding the flexibility of polymer chains. Flexible chains are known to dissolve more easily in a plasticizer. This mainly is due to lower energy requirement in separating polymer chains, and an easier diffusion of chains in the plasticizer. This process usually increases the mixing entropy and facilitates the mixing. This again demonstrates that the dissolution of crystalline or cross-linked polymer chains in a plasticizer is a laborious task and, therefore, the compatibility between these pairs is much less than the compatibility between flexible amorphous polymer chains and a plasticizer. Beside chain flexibility, crystallinity, and cross-linking, some other criteria have to be as well considered. Some are based on the polarity differences between plasticizer/ polymer pairs. Highly polar polymers do not usually dissolve in nonpolar solvents and the other way around. However, some exceptions to this rule have already been observed.

Up to this point, most of the discussion surrounding the compatibility issue was qualitative. However, to choose a proper plasticizer one should estimate the compatibility based on a measurable value derived from the Flory-Huggins theory [24,25]. The miscibility criterion in this theory is determined by the Gibbs free energy of mixing, which is a combination of the enthalpy and entropy of mixing:

$$\frac{\Delta G^M}{RTV} = \left[\frac{\phi\_1}{V\_1}\ln\phi\_1 + \frac{\phi\_2}{V\_2}\ln\phi\_2\right] + \left[\frac{\phi\_1\phi\_2}{V\_1}\mathcal{X}\_1\right] \tag{1}$$


**Table 2.** Description of parameters of eq. 1

The term in the first bracket represents the entropy change of mixing, whereas the second one represents the enthalpy change. In a polymer/plasticizer mixture, an instantaneous miscibility (or complete compatibility) is achieved when the Gibbs free energy of mixing is negative. This can only be achieved if the enthalpy contribution is small enough in comparison to the entropic one. Therefore, the interaction parameter (*χ*1), which repre‐ sents the enthalpic contribution, is an important feature in estimating the compatibility. It characterizes the difference between the interaction energy when a plasticizer molecule is immersed in a neat polymer, versus when it is immersed in a neat plasticizer. The upper limit for *χ*1 to obtain a compatible pair is 0.5. For values greater than 0.5, incompatibility and phase-separation may be observed. Although values for *χ*1 in different polymer/ solvent (plasticizer) systems can be found in the literature and in several polymer hand‐ books [16,26,27], in practice the most useful parameter used to estimate the molecular interactions between two components is the solubility parameter (δ). This parameter is the square root of the cohesive energy density. This is the energy that has to be given to a system of pure liquid to extract a molecule from the liquid state. Therefore, it is proportion‐ al to the interaction energies between similar molecules. The change in the internal energy of mixing can be related to the solubility parameter based on the following equation [28]:

$$
\Delta \mathcal{U}^M = \Delta H^M \Big|\_{\text{Const. Volume}} = V \phi\_1 \phi\_2 \Big[\delta\_1 - \delta\_2\Big]^2 \tag{2}
$$

The combination of eqs. 1 and 2 for the enthalpy change of mixing results in an equation relating the Flory-Huggins interaction parameter to the solubility parameters of the polymer and the plasticizer:

Plasticization and Morphology Development in Dynamically Vulcanized Thermoplastic Elastomers http://dx.doi.org/10.5772/61414 55

$$\mathcal{Z}\_1 = \frac{V\_1}{RT}(\mathcal{S}\_1 - \mathcal{S}\_2)^2 \tag{3}$$

#### 3 <sup>1</sup> *<sup>V</sup>* Molar volume of plasticizer, m /mol é ù ë û

measurable value derived from the Flory-Huggins theory [24,25]. The miscibility criterion in this theory is determined by the Gibbs free energy of mixing, which is a combination of the

> 1 2 1 2 12 1 12 1

The term in the first bracket represents the entropy change of mixing, whereas the second one represents the enthalpy change. In a polymer/plasticizer mixture, an instantaneous miscibility (or complete compatibility) is achieved when the Gibbs free energy of mixing is negative. This can only be achieved if the enthalpy contribution is small enough in comparison to the entropic one. Therefore, the interaction parameter (*χ*1), which repre‐ sents the enthalpic contribution, is an important feature in estimating the compatibility. It characterizes the difference between the interaction energy when a plasticizer molecule is immersed in a neat polymer, versus when it is immersed in a neat plasticizer. The upper limit for *χ*1 to obtain a compatible pair is 0.5. For values greater than 0.5, incompatibility and phase-separation may be observed. Although values for *χ*1 in different polymer/ solvent (plasticizer) systems can be found in the literature and in several polymer hand‐ books [16,26,27], in practice the most useful parameter used to estimate the molecular interactions between two components is the solubility parameter (δ). This parameter is the square root of the cohesive energy density. This is the energy that has to be given to a system of pure liquid to extract a molecule from the liquid state. Therefore, it is proportion‐ al to the interaction energies between similar molecules. The change in the internal energy of mixing can be related to the solubility parameter based on the following equation [28]:

 ff

> c

> > /mol]

2

ë û (2)

12 1 2 .

The combination of eqs. 1 and 2 for the enthalpy change of mixing results in an equation relating the Flory-Huggins interaction parameter to the solubility parameters of the polymer

ff d dé ù

*Const Volume* D =D = - *UH V* (1)

 f

 <sup>D</sup> <sup>é</sup> ùé ù =+ + <sup>ê</sup> úê ú êë úê ú ûë û

ff

*<sup>M</sup> <sup>G</sup> ln ln RTV V V V*

f

]

*ϕ***1,** *ϕ***<sup>2</sup>** Volume fractions of plasticizer and polymer, [-]

*M M*

*V***1,** *V***<sup>2</sup>** Molar volumes of plasticizer and polymer, respectively, [m3

enthalpy and entropy of mixing:

54 Thermoplastic Elastomers - Synthesis and Applications

*T* Temperature, [K]

**V** Total volume, [m3

**Table 2.** Description of parameters of eq. 1

and the plasticizer:

*χ***<sup>1</sup>** Interaction parameter, [-]

*R* Universal gas constant, [J/K.mol]

A more appropriate relationship documented in the literature for a mixture of polymer/low molecular weight liquid (plasticizer) is the following:

$$\chi\_1 = 0.34 + \frac{V\_1}{RT} (\delta\_1 - \delta\_2)^2 \tag{4}$$

Earlier, the limit of compatibility between polymer/plasticizer pairs was determined by a critical value (*χ*1=0.5). According to this criterion, compatibility can be approximately estimat‐ ed by matching the solubility parameters. A typical low molecular weight liquid (plasticizer) may have a molar volume around 100 to 400 cm3 /mol. Accordingly, this translates into a critical difference between the solubility parameters around 0.48 to 0.97 (cal/cm3 ) 0.5 (based on eq. 4). As an example, for natural rubber with solubility parameter of 8.30 (cal/cm3 )0.5 [29], a proper plasticizer shall possess a solubility parameter between 8.3 ± 0.97 (cal/cm3 ) 0.5. This simple method is widely used in industry and, although it is considered as a rapid screening technique to identify a proper plasticizer, there might be cases as we will demonstrate shortly that the prediction is not quite exact.

A more appropriate approach for compatibility studies in terms of solubility parameter is the use of the three-dimensional solubility components proposed by Hansen [30]. In this approach, the individual solubility parameter for each and every phase involved in the system is composed of contributions from van der Waals dispersion forces (δ*d*), dipole–dipole interaction between molecules (δ**p**), and the contribution from hydrogen bonding (δ*h*):

$$\mathcal{S} = \left(\mathcal{S}\_d^2 + \mathcal{S}\_p^2 + \mathcal{S}\_h^2\right)^{1/2} \tag{5}$$

The advantage of using the three-dimensional components of the solubility parameter is its ability in distinguishing between different chemical interactions, which might be present. For instance, two substances may have exactly the same overall solubility parameter, but with different proportions of δ*d*/δ*p*/δ*<sup>h</sup>* components. Ethylene carbonate and methanol both possess quite similar overall solubility parameters, around 29.0 MPa1/2 or 14.2 (cal/cm3 ) 0.5 [31]. How‐ ever, comparing their corresponding Hansen solubility components, one could clearly understand their difference. Ethylene carbonate with (δ*d*=18.0; δ*p*=21.7; δ*h*=5.1 MPa1/2) and methanol with (δ*d*=14.7; δ*p*=12.3; δ*h*=22.3 MPa1/2) have a huge difference in terms of hydrogenbonding interactions [31]. As a result, if these two are to be in contact with a polymer, the difference between their solubility capabilities could be readily predicted through the differ‐ ence in their three-dimensional values in the δ*d*,δ*p*,δ*h* coordinates. The three-dimensional distance between two substances is usually calculated based on the following equation:

$$D = \left(4 \times \left(\mathcal{S}\_{d,1} - \mathcal{S}\_{d,2}\right)^2 + \left(\mathcal{S}\_{p,1} - \mathcal{S}\_{p,2}\right)^2 + \left(\mathcal{S}\_{h,1} - \mathcal{S}\_{h,2}\right)^2\right)^{1/2} \tag{6}$$

A smaller distance is an indication of similarity and thermodynamic compatibility between two molecules. In the case of the previous example, the distance between ethylene carbonate and methanol molecules is around 20.7 MPa1/2. This shows that these two molecules are not at all similar and they will not behave in a similar manner when subjected to a polymer (regard‐ less of the type of polymer). In a plasticizer/polymer mixture, as the distance increases, the compatibility and the solubility decreases in a way that after a certain distance known as the polymer radius the compatibility is negligible. The simplest and the most practical way to calculate the compatibility between a polymer and a plasticizer is to calculate the threedimensional distance based on the Hansen solubility components and, then, divide *D* by the radius of the polymer. This ratio is known as the relative energy difference (RED). A mixture with a RED value smaller than 1 is compatible or even soluble in the best-case scenario; whereas a RED value greater than 1 is a sign of incompatibility and insolubility. This approach provides a more accurate prediction for compatibility; however, care must be taken when dealing with complex mixtures with cross-linked or crystalline polymers. These complexities reduce the radius of a polymer and affect negatively the compatibility with a plasticizer.

In industry, a more complex situation is generally encountered. A plasticizer is usually incorporated into a blend of two or more polymeric constituents. Since the majority of polymer blends are known to be immiscible due to thermodynamic limitations, a question arises regarding the distribution of the plasticizer when mixed in multiphase polymer blends. The answer to this question could simply clarify whether the properties of both phases are affected equally in the presence of a plasticizer, or only one of the phases will be largely affected due to its higher affinity with the plasticizer. To quantify the characteristics of the distribution, a quantity known as distribution or partition coefficient (*K*A/B) has been widely used. This coefficient is the ratio of the weight fraction of the plasticizer in the polymer A over the weight fraction of the plasticizer in the polymer B, i.e., *K*A/B = *wp*,A/*wp*,B. A value of *K*A/B = 1 means that the plasticizer is equally distributed in both polymers; whereas, values lower than 1 mean that the plasticizer has a tendency toward polymer *B*. According to Mishra et al. [32], in those immiscible blends where the distribution of an additional low molecular weight component is entropically driven, the lowest free energy of mixing is usually achieved in the vicinity of *K*A/B ~ 1. This implies that in a system where the plasticizer has a close affinity with both polymers, a uniform distribution is thermodynamically favored. On the other hand, in a blend where only one of the components is compatible with the plasticizer, the plasticizer will obviously migrate to that specific phase. An example of such system is the blend of poly(methyl methacrylate) (PMMA) and acrylonitrile butadiene styrene (ABS) [33]. These two polymers are mutually miscible. However, in the presence of a plasticizer composed of ethylene carbonate and propylene carbonate (1:1), a phase-separated system is obtained. Since the plasticizer mixture happens to have more affinity with PMMA, a plasticized PMMA phase and an ABS-rich phase are obtained. In this special case, both ABS and PMMA are amorphous polymers. Blends with semicrystalline polymers could behave quite differently. In a conven‐ tional compounding process, the semicrystalline polymer crystallizes upon cooling from the molten state. During this stage, the plasticizer even though miscible or compatible with the semicrystalline polymer in the molten state usually migrates to the other phase. As a result, the resulting distribution of the plasticizer in the molten state could be largely different from that of the solid state. In a blend composed of polypropylene (PP) and styrene-ethylenebutadiene-styrene (SEBS), Ohlsson et al. [34] have concluded that the distribution coefficient (*K*PP/SEBS) should vary between 0.33 to 0.47 for blends with 10 to 90 wt% of polypropylene (a semicrystalline polymer). This indicates a preferential distribution of the plasticizer in the SEBS. In dynamically vulcanized blends composed of EPDM/PP in the presence of a paraffinic oil, the distribution of plasticizer has been shown to depend on the concentration and the molecular weight distribution of the polypropylene, as well as on the crystallized or molten state of the material [35]. The *K*PP/EPDM was lower or close to 1 at room temperature, whereas at an elevated temperature, i.e., 190 °C where polypropylene is molten, the distribution coeffi‐ cient was larger than 1. Based on a micro-mechanical modeling approach, it has also been shown that the distribution coefficient in a nonreactive PP/SEBS and dynamically vulcanized EPDM/PP blends is in favor of the elastomeric component, but varying with the composition [36]. By considering the rigid nature and difficulty of plasticizing the styrenic blocks in SEBS and crystalline portion of polypropylene, the average corrected *K*PP/SEBS has been reported to be around 0.51 for the molten state, and 0.76 for the solid state [37]. The same ratio (*K*PP/EPDM) of plasticized dynamically vulcanized PP/EPDM blends was reported to be around 0.89 in the solid state. The scanning electron microscopy (SEM) images of rapidly cooled extruded strands of nonreactive EPDM/PP blends have also provided a significant insight into the distribution of plasticizer [14]. The oriented structure of the EPDM phase in the plasticized extrudate versus the nonoriented structure observed in the nonplasticized blends was associated with the predominant presence of plasticizer in the elastomer phase. The possibility of obtaining an oriented elastomeric structure in the presence of plasticizer was explained in terms of large drop in the rheological properties of the elastomer in the presence of plasticizer to a level at which the polypropylene phase could deform the EPDM phase, which was otherwise a highly viscous and elastic material. Once again, these studies illustrate that even though the chosen plasticizer could have a high affinity with both polymeric components, the preferential distribution of the plasticizer is toward the elastomeric component. In a few studies, however, a separate phase mainly composed of plasticizer was also reported [38,39].

ence in their three-dimensional values in the δ*d*,δ*p*,δ*h* coordinates. The three-dimensional distance between two substances is usually calculated based on the following equation:

( ) ( ) ( ) 2 2 <sup>2</sup> 1/2

A smaller distance is an indication of similarity and thermodynamic compatibility between two molecules. In the case of the previous example, the distance between ethylene carbonate and methanol molecules is around 20.7 MPa1/2. This shows that these two molecules are not at all similar and they will not behave in a similar manner when subjected to a polymer (regard‐ less of the type of polymer). In a plasticizer/polymer mixture, as the distance increases, the compatibility and the solubility decreases in a way that after a certain distance known as the polymer radius the compatibility is negligible. The simplest and the most practical way to calculate the compatibility between a polymer and a plasticizer is to calculate the threedimensional distance based on the Hansen solubility components and, then, divide *D* by the radius of the polymer. This ratio is known as the relative energy difference (RED). A mixture with a RED value smaller than 1 is compatible or even soluble in the best-case scenario; whereas a RED value greater than 1 is a sign of incompatibility and insolubility. This approach provides a more accurate prediction for compatibility; however, care must be taken when dealing with complex mixtures with cross-linked or crystalline polymers. These complexities reduce the

 dd

(6)

,1 ,2 ,1 ,2 ,1 ,2 (4 ) *D dd pp hh* =´ - + - + -

 dd

radius of a polymer and affect negatively the compatibility with a plasticizer.

In industry, a more complex situation is generally encountered. A plasticizer is usually incorporated into a blend of two or more polymeric constituents. Since the majority of polymer blends are known to be immiscible due to thermodynamic limitations, a question arises regarding the distribution of the plasticizer when mixed in multiphase polymer blends. The answer to this question could simply clarify whether the properties of both phases are affected equally in the presence of a plasticizer, or only one of the phases will be largely affected due to its higher affinity with the plasticizer. To quantify the characteristics of the distribution, a quantity known as distribution or partition coefficient (*K*A/B) has been widely used. This coefficient is the ratio of the weight fraction of the plasticizer in the polymer A over the weight fraction of the plasticizer in the polymer B, i.e., *K*A/B = *wp*,A/*wp*,B. A value of *K*A/B = 1 means that the plasticizer is equally distributed in both polymers; whereas, values lower than 1 mean that the plasticizer has a tendency toward polymer *B*. According to Mishra et al. [32], in those immiscible blends where the distribution of an additional low molecular weight component is entropically driven, the lowest free energy of mixing is usually achieved in the vicinity of *K*A/B ~ 1. This implies that in a system where the plasticizer has a close affinity with both polymers, a uniform distribution is thermodynamically favored. On the other hand, in a blend where only one of the components is compatible with the plasticizer, the plasticizer will obviously migrate to that specific phase. An example of such system is the blend of poly(methyl methacrylate) (PMMA) and acrylonitrile butadiene styrene (ABS) [33]. These two polymers are mutually miscible. However, in the presence of a plasticizer composed of ethylene carbonate and propylene carbonate (1:1), a phase-separated system is obtained. Since the

dd

56 Thermoplastic Elastomers - Synthesis and Applications

## **3. Morphology development in simple nonreactive thermoplastic/ elastomer blends**

It was mentioned earlier that the majority of polymer blends form immiscible systems due to thermodynamic limitations. Consequently, blending usually results in a multiphase hetero‐ geneous morphology, which along with other properties of the constituent polymers dictates the ultimate properties of the resulting compound. Nowadays, compounding is performed in conventional melt-mixing equipment, such as twin-screw extruders, Banbury, or any other internal mixers. The morphology development during the compounding stage strongly depends on the rheological and the interfacial properties of the constituent polymers, their concentrations, the processing conditions and whether or not other additives such as plasti‐ cizers, curatives or fillers are present in the system. As a result, for a given immiscible blend composed of only two polymeric components, a wide range of morphologies can be tailored to specifically fulfill the requirements related to the end-use application of the blend. A droplet/ matrix type morphology, if well designed, may improve the impact properties and the toughness of materials; the lamellar type will improve the barrier properties; the fibrillar type may improve the tensile properties and the stiffness of the material; and finally the cocontinuous morphology may have combinatorial effects due to the simultaneous contributions of both phases at the same time.

Generally, compounding process in conventional melt-mixing equipment begins with the materials in their solid state, e.g., pellets, powders, or bales. These are subjected to intensive mixing conditions such as elevated temperature alongside with shear, which transforms the initial solid state to a molten liquid. As a result, materials are gradually softened, deformed, and finally become molten polymeric components with corresponding viscoelastic properties. The initial transformation of solid pellets to micron-size particles in the early stage of mixing process was shown to be achieved through sheet formation [40]. This mechanism is based on the mutual contact between the solid particles of the dispersed phase with the hot metal surface of the mixing equipment, which alongside with shear results in the formation of sheets/ribbons. The final micron-size particles are finally formed through transformation of these sheets into cylinders, which themselves are broken-up through Rayleigh instabilities. Although this explains the dispersion mechanism of the minor phase, one also needs to take into account the melting behavior of the constituent polymers in the blend. At the early stage of mixing, the melting sequence has a great importance [41,42]. In the case where the minor phase has a lower melting point than the major phase, it rapidly melts and encapsulates the major phase. However, as the temperature of the bulk increases with time during the mixing operation, the major phase begins to melt and becomes the matrix encapsulating the minor phase. As a result, in those blends with component melting characteristics far apart, the melting sequence could be crucial in determining the final morphology.

Once all polymeric components are molten, the mixed medium becomes essentially a mixture of viscoelastic fluids. The morphology refinement at this stage depends on several other parameters, such as shear and elongation rates, mixing time, rheological properties of the constituent polymers, blending composition, interfacial properties, and the presence of other solid or liquid additives (e.g., fillers or plasticizers). In dynamic mixing conditions, the flow and deformation of polymeric components are closely associated to their rheological proper‐ ties. Among all the rheological properties, the viscosity ratio of the polymeric components at the processing conditions is one of the well-known factors that have been directly related to the morphology of the blend [43,44]. For viscosity ratios greater than 1 (ηA/η<sup>B</sup> > 1), the size of the polymeric domains is known to increase monotonically. Furthermore, Favis and Chalifoux [43] have observed that in contrast to Newtonian dispersions, an immiscible blend with a high viscosity ratio (ηminor/ηmajor ~ 17) is still deformable. This essentially means that in a complex flow field, where shear and elongational flow coexist, it is still possible to deform a highly viscous dispersed phase. In the other side of viscosity ratio range, a composition dependency has been observed [43]. In the low composition range, the minimum domain size in polypro‐ pylene/polycarbonate (PP/PC) blends was reported to be reached around (ηminor/ηmajor ~ 0.15 and below this value no significant change in the domain size was observed [43]. A similar observation where the polydispersity and the average domain size increased with the viscosity ratio has been reported in EPDM/PP blends with EPDM as the minor phase [45].

the ultimate properties of the resulting compound. Nowadays, compounding is performed in conventional melt-mixing equipment, such as twin-screw extruders, Banbury, or any other internal mixers. The morphology development during the compounding stage strongly depends on the rheological and the interfacial properties of the constituent polymers, their concentrations, the processing conditions and whether or not other additives such as plasti‐ cizers, curatives or fillers are present in the system. As a result, for a given immiscible blend composed of only two polymeric components, a wide range of morphologies can be tailored to specifically fulfill the requirements related to the end-use application of the blend. A droplet/ matrix type morphology, if well designed, may improve the impact properties and the toughness of materials; the lamellar type will improve the barrier properties; the fibrillar type may improve the tensile properties and the stiffness of the material; and finally the cocontinuous morphology may have combinatorial effects due to the simultaneous contributions

Generally, compounding process in conventional melt-mixing equipment begins with the materials in their solid state, e.g., pellets, powders, or bales. These are subjected to intensive mixing conditions such as elevated temperature alongside with shear, which transforms the initial solid state to a molten liquid. As a result, materials are gradually softened, deformed, and finally become molten polymeric components with corresponding viscoelastic properties. The initial transformation of solid pellets to micron-size particles in the early stage of mixing process was shown to be achieved through sheet formation [40]. This mechanism is based on the mutual contact between the solid particles of the dispersed phase with the hot metal surface of the mixing equipment, which alongside with shear results in the formation of sheets/ribbons. The final micron-size particles are finally formed through transformation of these sheets into cylinders, which themselves are broken-up through Rayleigh instabilities. Although this explains the dispersion mechanism of the minor phase, one also needs to take into account the melting behavior of the constituent polymers in the blend. At the early stage of mixing, the melting sequence has a great importance [41,42]. In the case where the minor phase has a lower melting point than the major phase, it rapidly melts and encapsulates the major phase. However, as the temperature of the bulk increases with time during the mixing operation, the major phase begins to melt and becomes the matrix encapsulating the minor phase. As a result, in those blends with component melting characteristics far apart, the melting sequence could

Once all polymeric components are molten, the mixed medium becomes essentially a mixture of viscoelastic fluids. The morphology refinement at this stage depends on several other parameters, such as shear and elongation rates, mixing time, rheological properties of the constituent polymers, blending composition, interfacial properties, and the presence of other solid or liquid additives (e.g., fillers or plasticizers). In dynamic mixing conditions, the flow and deformation of polymeric components are closely associated to their rheological proper‐ ties. Among all the rheological properties, the viscosity ratio of the polymeric components at the processing conditions is one of the well-known factors that have been directly related to the morphology of the blend [43,44]. For viscosity ratios greater than 1 (ηA/η<sup>B</sup> > 1), the size of the polymeric domains is known to increase monotonically. Furthermore, Favis and Chalifoux

of both phases at the same time.

58 Thermoplastic Elastomers - Synthesis and Applications

be crucial in determining the final morphology.

The blend composition is another major factor affecting the morphology. When two immiscible polymers are compounded, the morphology in the low composition range mainly consists of droplets of the minor phase in the matrix of the major component. The size and the polydis‐ persity of the emulsion largely depend on the compatibility and, therefore, the interfacial tension and ratio of rheological properties between the polymeric components. As the interfacial tension gets smaller and the viscosity ratio approaches unity, finer droplets of the minor phase are usually observed [46,47]. Further increase in the concentration of the minor phase results in a coarser morphology due to coalescence, where droplets are coalesced with each other and form an emulsion of larger size droplets and increased polydispersity. The increase in the concentration of the minor phase eventually transforms the morphology into a co-continuous type. Each polymeric phase in a co-continuous structure is interconnected throughout the whole bulk of the material and, therefore, both components are expected to contribute simultaneously in the overall properties of the material. Immiscible polymer blends with highly viscous and elastic components generally show a rather wide range of cocontinuous composition range. At both extremities of the co-continuity interval, the polymer that constitutes the major phase tends to encapsulate the minor phase. This results in break‐ down of the co-continuous structure by transforming it into dispersed-type morphology. The aforementioned co-continuity range and the factors affecting it have major importance in thermoplastic/elastomer blends. As mentioned earlier, most reactive dynamically vulcanized thermoplastic/elastomer blends are produced from their corresponding nonreactive precursor blend at a composition range that coincides with the co-continuity interval [13]. This is mainly due to the fact that in this composition range, there is sufficient amount of elastomeric and thermoplastic components present, which eventually after dynamic vulcanization the elasto‐ mer is able to provide the flexibility and rubber-like behavior, whereas the thermoplastic phase guarantees the complete encapsulation of the elastomer phase when it is transformed into dispersed particles.

Several parameters may affect the co-continuity range. Parameters such as the viscosity of the polymeric components, interfacial tension, presence of plasticizer, and mixing time are those which have been carefully studied [48-50]. The early works in understanding the concept of co-continuity were merely concentrated in investigating the effects of concentration and viscosity ratio. Recent studies, however, did not necessarily consider the viscosity ratio as an independent parameter when evaluating the onset of co-continuity [48,51]. Basically for a given viscosity ratio, the onset and therefore the width of the co-continuity range has largely been attributed to the viscosity of the major phase (ηmajor) regardless of the viscosity of the minor one (ηminor); after all, it is the former that is responsible for imposing stresses on the minor phase during processing. Based on some theoretical background on the stability of deformed threadlike emulsions and their packing density, it has been shown that a larger viscosity of the major phase shifted the onset of co-continuity to a lower concentration of the minor phase component [51]. This has also been experimentally observed [49] in a blend composed of EPDM/PP, where the presence of a highly viscous and elastic EPDM in comparison to a less viscous and elastic PP resulted in an asymmetric co-continuity interval, with the onset of cocontinuity shifted to a lower concentration of the less viscous PP, and higher concentration of highly viscous EPDM phase as illustrated in Fig. 2.

**Figure 2.** Continuity index of EPDM phase in both plasticized and nonplasticized blends. An asymmetric co-continuity can be clearly observed especially in the case of the nonplasticized system (from [49])

In blends with an identical viscosity ratio, the width of the co-continuity interval could be linked to the interfacial tension. Comparing two distinct immiscible blends with identical viscosity ratio, the one with a higher interfacial tension will have an onset of co-continuity shifted to higher concentrations with a narrow co-continuity interval [48]. Li et al. [52] proposed an interfacial dependent mechanism for the co-continuous morphology develop‐ ment. The authors proposed two distinctive coalescence mechanisms responsible for the cocontinuity formation, thread–thread coalescence for low interfacial tension blends, whereas droplet–droplet coalescence mechanism for high interfacial tension blends. Accordingly, the onset and the width of the co-continuity interval is not merely associated with the rheological properties of the constituent polymers, but also to the mobility of interface and the ease of coalescence and percolation of the minor phase.

The use of low molecular weight plasticizers in an immiscible thermoplastic/elastomer blend may simultaneously affect the rheological properties of the components, the interfacial tension, and the volume of the individual phases through swelling. Based on the earlier discussion on the plasticization of thermoplastic/elastomer blends in the molten state, plasticizers have a slight tendency toward the elastomeric phase. This could clearly indicate that the rheological properties of the elastomer could be largely dropped in highly plasticized systems. Conse‐ quently, the ease of deformation and coalescence between elastomeric domains with simulta‐ neous swelling in the presence of plasticizer could result in a coarser morphology, as shown in Fig. 3 [49]. Meanwhile in plasticized blends, a faster percolation and higher continuity index of the elastomeric phase at a lower concentration illustrated in Fig. 2 can be readily understood when the morphologies of nonplasticized and plasticized blends are compared (Fig. 3). tendency toward the elastomeric phase. This could clearly indicate that the rheological properties of the elastomer could be largely dropped in highly plasticized systems. Consequently, the ease of deformation and coalescence between elastomeric domains with simultaneous swelling in the presence of plasticizer could result in a coarser morphology, as shown in Fig. 3 [49]. Meanwhile in plasticized blends, a faster percolation and higher continuity index of the elastomeric phase at a lower concentration illustrated in Fig. 2 can be readily understood when the morphologies of nonplasticized and plasticized blends are compared (Fig. 3).

The use of low molecular weight plasticizers in an immiscible thermoplastic/elastomer blend may simultaneously affect the rheological properties of the components, the interfacial tension,

plasticization of thermoplastic/elastomer blends in the molten state, plasticizers have a slight

interfacial dependent mechanism for the co-continuous morphology development. The authors proposed two distinctive coalescence mechanisms responsible for the co-continuity formation, thread–thread coalescence for low interfacial tension blends, whereas droplet–droplet coalescence mechanism for high interfacial tension blends. Accordingly, the onset and the width of the co-continuity interval is not merely associated with the rheological properties of the constituent polymers, but also to the mobility of interface and the ease of coalescence and

percolation of the minor phase.

been attributed to the viscosity of the major phase (ηmajor) regardless of the viscosity of the minor one (ηminor); after all, it is the former that is responsible for imposing stresses on the minor phase during processing. Based on some theoretical background on the stability of deformed threadlike emulsions and their packing density, it has been shown that a larger viscosity of the major phase shifted the onset of co-continuity to a lower concentration of the minor phase component [51]. This has also been experimentally observed [49] in a blend composed of EPDM/PP, where the presence of a highly viscous and elastic EPDM in comparison to a less viscous and elastic PP resulted in an asymmetric co-continuity interval, with the onset of cocontinuity shifted to a lower concentration of the less viscous PP, and higher concentration of

**Figure 2.** Continuity index of EPDM phase in both plasticized and nonplasticized blends. An asymmetric co-continuity

In blends with an identical viscosity ratio, the width of the co-continuity interval could be linked to the interfacial tension. Comparing two distinct immiscible blends with identical viscosity ratio, the one with a higher interfacial tension will have an onset of co-continuity shifted to higher concentrations with a narrow co-continuity interval [48]. Li et al. [52] proposed an interfacial dependent mechanism for the co-continuous morphology develop‐ ment. The authors proposed two distinctive coalescence mechanisms responsible for the cocontinuity formation, thread–thread coalescence for low interfacial tension blends, whereas droplet–droplet coalescence mechanism for high interfacial tension blends. Accordingly, the onset and the width of the co-continuity interval is not merely associated with the rheological properties of the constituent polymers, but also to the mobility of interface and the ease of

The use of low molecular weight plasticizers in an immiscible thermoplastic/elastomer blend may simultaneously affect the rheological properties of the components, the interfacial tension, and the volume of the individual phases through swelling. Based on the earlier discussion on

can be clearly observed especially in the case of the nonplasticized system (from [49])

coalescence and percolation of the minor phase.

highly viscous EPDM phase as illustrated in Fig. 2.

60 Thermoplastic Elastomers - Synthesis and Applications

Fig. 3. Coarsening effect of a plasticizer in plasticized nonreactive blends of EPDM/PP 25/75 (wt/wt%): (a) nonplasticized (b) plasticized; dark phase: EPDM; bright phase: PP (from [49]) **Figure 3.** Coarsening effect of a plasticizer in plasticized nonreactive blends of EPDM/PP 25/75 (wt/wt%): (a) nonplasti‐ cized (b) plasticized; dark phase: EPDM; bright phase: PP (from [49])

Processing time is another factor that may or may not affect the morphology and it has been investigated by several authors [45,50,53,54]. In few studies performed on blends with high viscosity ratio ( *<sup>η</sup>minor phase <sup>η</sup>major phase* <sup>≥</sup>1) and low composition ranges of the dispersed phase, the effect of mixing time was shown to be insignificant regardless of the mixing equipment used, i.e., continuous twin-screw extrusion or internal mixer [53,54]. This is mainly due to the fact that the main deformation and disintegration process is considered to take place within the first few minutes of mixing. A more thorough study covering a wider composition range and blends with viscosity ratios lower or equal to unity demonstrated that the mixing time affected the morphology development, especially in the low viscosity ratio blends [50]. Bu and He [46] work investigated two blends with interfacial tensions of the same order of magnitude, but with different viscosity ratios. The morphology at longer processing times appeared not to be influenced by the viscosity ratio and only by the blend composition. However, at an early stage of mixing, the morphology seemed to depend on the viscosity ratio and whether the low or high viscous component formed the minor phase. For the low viscosity ratio system, when the low viscosity polymer formed the minor phase, the morphology consisted of droplets at low compositions and with further increases in the composition a fibrillar and eventually a cocontinuous morphology was formed at a relatively low composition range. On the other hand, when the high viscosity component was the minor phase, it appeared in the form of dispersed droplets and at high concentrations it was transformed into a continuous phase. For blends with viscosity ratio in the vicinity of one, the dispersed phase in either side of the composition range appeared to be in the form of droplets up to a range where a co-continuous morphology was formed. Furthermore, regardless of the blending system and viscosity ratio, the width of the co-continuity interval was shown to decrease with mixing time. Among other processing parameters, the rotational speed (or indirectly the deformation rate) in an internal mixer has been shown not to have a significant effect on the morphology, especially in the low compo‐ sition range [53,55].

In both rubber and plastic industries the compounding is usually performed in batch internal mixers or via continuous twin-screw extruders. The twin-screw extrusion process is widely known to be an excellent and versatile technique especially due to its modular design capa‐ bilities. The possibility and the advantage of designing different screw configurations gener‐ ally results in intensive mixing conditions, which provides an efficient distributive and dispersive flow characteristics. Therefore, the comparison between the final morphology and its evolution in an internal mixer and a twin-screw extruder is not as straightforward as it could be imagined. However, although the types of flow fields and their intensities are largely different in those two types of equipment, the overall morphology evolution passes through the same sequences as shown by Sundararaj et al. [54] for polystyrene/polyamide (PS/PA) and polystyrene/polypropylene (PS/PP) blends. The same sequences of sheet formation, transfor‐ mation into elongated domains, and eventually formation of droplets of the minor phase were observed. Regarding the final domain size of the morphological features, twin-screw extrusion has generally resulted in similar or even finer morphology in the blends [14,56,57]. In both nonplasticized and plasticized thermoplastic/elastomer blends within the co-continuous composition range, Shahbikian et al. [14] have shown that the use of twin-screw extrusion substantially refined the morphology and increased the interfacial area as illustrated in Figs. 4 and 5. The more refined structure especially in the co-continuous composition range provides a more desirable initial morphological state for further reactive and dynamic vulcanization of thermoplastic/elastomer blends.

**Figure 4.** Atomic force microscopy images of EPDM/PP 50/50 (wt/wt%) TPOs: (a) internal mixer, (b) twin-screw ex‐ truder. (Column 1: non-plasticized, Column 2: plasticized; Dark phase: EPDM; Bright phase: PP) (from [14])

the co-continuity interval was shown to decrease with mixing time. Among other processing parameters, the rotational speed (or indirectly the deformation rate) in an internal mixer has been shown not to have a significant effect on the morphology, especially in the low compo‐

In both rubber and plastic industries the compounding is usually performed in batch internal mixers or via continuous twin-screw extruders. The twin-screw extrusion process is widely known to be an excellent and versatile technique especially due to its modular design capa‐ bilities. The possibility and the advantage of designing different screw configurations gener‐ ally results in intensive mixing conditions, which provides an efficient distributive and dispersive flow characteristics. Therefore, the comparison between the final morphology and its evolution in an internal mixer and a twin-screw extruder is not as straightforward as it could be imagined. However, although the types of flow fields and their intensities are largely different in those two types of equipment, the overall morphology evolution passes through the same sequences as shown by Sundararaj et al. [54] for polystyrene/polyamide (PS/PA) and polystyrene/polypropylene (PS/PP) blends. The same sequences of sheet formation, transfor‐ mation into elongated domains, and eventually formation of droplets of the minor phase were observed. Regarding the final domain size of the morphological features, twin-screw extrusion has generally resulted in similar or even finer morphology in the blends [14,56,57]. In both nonplasticized and plasticized thermoplastic/elastomer blends within the co-continuous composition range, Shahbikian et al. [14] have shown that the use of twin-screw extrusion substantially refined the morphology and increased the interfacial area as illustrated in Figs. 4 and 5. The more refined structure especially in the co-continuous composition range provides a more desirable initial morphological state for further reactive and dynamic vulcanization of

**Figure 4.** Atomic force microscopy images of EPDM/PP 50/50 (wt/wt%) TPOs: (a) internal mixer, (b) twin-screw ex‐

truder. (Column 1: non-plasticized, Column 2: plasticized; Dark phase: EPDM; Bright phase: PP) (from [14])

sition range [53,55].

62 Thermoplastic Elastomers - Synthesis and Applications

thermoplastic/elastomer blends.

**Figure 5.** Specific interfacial area (Q) of EPDM/PP 50/50 (wt/wt%) prepared in internal mixer and twin-screw extruder (from [14])

## **4. Morphology development in dynamically vulcanized nonplasticized and plasticized thermoplastic/elastomer blends**

The evolution of phase morphology in reactive blends generally takes place while the rheo‐ logical, interfacial, and thermodynamic properties of the components are changing due to the chemical reaction. In the case of dynamic vulcanization of thermoplastic/elastomer blends, it is mainly the selective cross-linking reaction of the elastomeric component which influences all the aforementioned properties. The gradual formation of cross-linked elastomer network increases both the viscosity and the elasticity of the elastomer, and affects drastically the morphology development. Throughout this process, the elastomeric major phase, although within a co-continuous structure with its thermoplastic counterpart, is transformed into dispersed particles encapsulated by the thermoplastic polymer. This morphological transfor‐ mation is widely known as phase inversion. The result is a material most likely with rubberlike properties. Furthermore, dynamically vulcanized blends are melt-processable through existing thermoplastic processing equipments and their final morphology is stable and cannot be altered by any downstream operation. These special advantages of dynamically vulcanized blends have led the interest of both academia and industry toward the improvement of properties through continuously optimizing the parameters affecting their morphology development.

It is known that the elastomer phase in the initial stage of vulcanization is strongly deformed into continuous elastomeric threads and eventually breaks up and forms the final dispersed cross-linked domains [58]. As mentioned earlier, the presence of an initial co-continuous morphology prior to dynamic vulcanization is a prerequisite in obtaining fine dispersed crosslinked elastomer phase at the end of vulcanization [58,59]. From a morphological point of view, only an initial co-continuous morphology results in an effective and overall transfer of the shear and elongation stresses from one phase to the other and, hence, guarantees the afore‐ mentioned breakup of the elastomeric component [58]. The effectiveness of stress transfer and morphology transformation in an initially co-continuous morphology can be visualized by considering the complete opposite hypothetical situation, i.e., an initial dispersed/matrix morphology with the elastomeric component as the droplet phase. In this system, the viscosity and elasticity of the elastomeric component, which already forms the dispersed phase, increase during dynamic vulcanization. As a result, the stress transferred by the low viscosity thermo‐ plastic phase becomes less and less effective in deforming the elastomeric domains and coalescence of the elastomeric domains becomes more and more hindered due to increased viscosity and elasticity of this phase. Consequently, the cross-linking reaction merely stabilizes the already existing, rather coarse dispersed morphology of the elastomeric domains, without any further morphological refinement [60-62]. With this in mind, although the co-continuity is a crucial factor in this process, a stable and unchangeable co-continuous morphology in the intermediate stage of dynamic vulcanization is not a desirable situation. This has been observed in ethylene methyl acrylate and linear low density polyethylene (EMA/LLDPE) blends where phase inversion was hindered due to the presence of a stable co-continuous morphology [63].

To obtain an optimum initial morphology prior to dynamic vulcanization, all the aforemen‐ tioned parameters discussed in the previous section have to be taken into consideration. For instance, a blend with extremely low viscosity minor thermoplastic phase (low viscosity ratio, ηthermoplastic/ηelastomer) generally forms a dispersed/matrix morphology, where the thermoplastic phase encapsulates coarse elastomeric domains [61]. This is a condition which has to be avoided. On the opposite situation, a highly viscous thermoplastic phase (a high viscosity ratio blend) hinders the dispersion of the cross-linked elastomer particles [62]. The fact that high and low viscosity ratio systems represent completely different initial morphological states prior to dynamic vulcanization means that the torque requirement for mixing could be substantially different from one system to another. Indeed, blends with high initial viscosity ratio with a dispersed thermoplastic phase in an elastomeric matrix have demonstrated a shoulder in the mixing torque during dynamic vulcanization process [61]. This shoulder appears while the blend structure passes through a co-continuous morphology prior to complete vulcanization and dispersion of the elastomeric component in the thermoplastic phase. The shoulder was seen to disappear with less overall mixing torque requirement when the initial morphological state was already a co-continuous type. The appearance of a shoulder in mixing curves has been reported several times and it has been attributed to the onset of phase inversion [60,64]. From an industrial point of view mainly based on the energy con‐ sumption during the dynamic vulcanization step, an initial co-continuous structure is desired over the dispersed/matrix one. The co-continuous state guarantees a smoother phase transition with lower energy consumption (lower torque requirement).

In several cases, however, the appearance of a shoulder in the mixing torque could be over‐ shadowed and the phase inversion process could not be easily detected due to the rapid crosslinking reaction [65]. Hence, the kinetics of cross-linking can also play a major role in the phase morphology development during dynamic vulcanization. Generally, the rate of the crosslinking reaction increases with temperature, especially in a blend with a heat reactive curing system such as phenolic resin [14,64,66]. The higher the temperature, the shorter is the time required to reach a certain level of cross-linking or gel content. The consequence of a faster reaction has been associated with the hindrance in the complete disintegration and dispersion of cross-linked elastomer phase during dynamic vulcanization [60]. At comparatively slower reaction rates, a smoother phase inversion with a better dispersion of elastomeric domains could be usually expected. With similar idea in mind, it has already been observed that the dispersion of a pre-cross-linked elastomer in a thermoplastic phase is a serious challenge especially when the gel content of the elastomer is larger than 70% [67]. Some other reactive systems other than thermoplastic/elastomer blends are as well shown to be affected by the effect of excessive gel content when dispersing one phase into the other. The dispersion of an *in situ* cured epoxy into a polystyrene (PS) is an example, where coalesced and agglomerated structure has been observed for gel content larger than 70% [68].

only an initial co-continuous morphology results in an effective and overall transfer of the shear and elongation stresses from one phase to the other and, hence, guarantees the afore‐ mentioned breakup of the elastomeric component [58]. The effectiveness of stress transfer and morphology transformation in an initially co-continuous morphology can be visualized by considering the complete opposite hypothetical situation, i.e., an initial dispersed/matrix morphology with the elastomeric component as the droplet phase. In this system, the viscosity and elasticity of the elastomeric component, which already forms the dispersed phase, increase during dynamic vulcanization. As a result, the stress transferred by the low viscosity thermo‐ plastic phase becomes less and less effective in deforming the elastomeric domains and coalescence of the elastomeric domains becomes more and more hindered due to increased viscosity and elasticity of this phase. Consequently, the cross-linking reaction merely stabilizes the already existing, rather coarse dispersed morphology of the elastomeric domains, without any further morphological refinement [60-62]. With this in mind, although the co-continuity is a crucial factor in this process, a stable and unchangeable co-continuous morphology in the intermediate stage of dynamic vulcanization is not a desirable situation. This has been observed in ethylene methyl acrylate and linear low density polyethylene (EMA/LLDPE) blends where phase inversion was hindered due to the presence of a stable co-continuous

To obtain an optimum initial morphology prior to dynamic vulcanization, all the aforemen‐ tioned parameters discussed in the previous section have to be taken into consideration. For instance, a blend with extremely low viscosity minor thermoplastic phase (low viscosity ratio, ηthermoplastic/ηelastomer) generally forms a dispersed/matrix morphology, where the thermoplastic phase encapsulates coarse elastomeric domains [61]. This is a condition which has to be avoided. On the opposite situation, a highly viscous thermoplastic phase (a high viscosity ratio blend) hinders the dispersion of the cross-linked elastomer particles [62]. The fact that high and low viscosity ratio systems represent completely different initial morphological states prior to dynamic vulcanization means that the torque requirement for mixing could be substantially different from one system to another. Indeed, blends with high initial viscosity ratio with a dispersed thermoplastic phase in an elastomeric matrix have demonstrated a shoulder in the mixing torque during dynamic vulcanization process [61]. This shoulder appears while the blend structure passes through a co-continuous morphology prior to complete vulcanization and dispersion of the elastomeric component in the thermoplastic phase. The shoulder was seen to disappear with less overall mixing torque requirement when the initial morphological state was already a co-continuous type. The appearance of a shoulder in mixing curves has been reported several times and it has been attributed to the onset of phase inversion [60,64]. From an industrial point of view mainly based on the energy con‐ sumption during the dynamic vulcanization step, an initial co-continuous structure is desired over the dispersed/matrix one. The co-continuous state guarantees a smoother phase transition

In several cases, however, the appearance of a shoulder in the mixing torque could be over‐ shadowed and the phase inversion process could not be easily detected due to the rapid crosslinking reaction [65]. Hence, the kinetics of cross-linking can also play a major role in the phase morphology development during dynamic vulcanization. Generally, the rate of the crosslinking reaction increases with temperature, especially in a blend with a heat reactive curing

with lower energy consumption (lower torque requirement).

morphology [63].

64 Thermoplastic Elastomers - Synthesis and Applications

Beside the kinetics of the curing reaction, the mixing intensity is also a crucial factor which clearly affects the morphology of reactive blends. By mixing intensity, we mainly consider the apparent shear rate of the mixing system. Although this is an important mixing parameter, it cannot be easily dissociated from the effect of temperature through viscous dissipation. In a simple argument, one may observe an increase in the mixing temperature by increasing the rotational speed of the mixing equipment, i.e., by increasing the intensity of mixing. This additional increase in temperature could readily affect the morphology development through its influence on the rate of the cross-linking reaction. In a situation like that, the morphology development is said to be mainly controlled by the fragmentation of the cross-linked elastomer rather than by the transient equilibrium between coalescence and breakup of particles [61]. It has been pointed out that the significant effect of mixing intensity on the final morphology is through its effect on the rate of the cross-linking reaction [60,63].

Mixing intensity can be otherwise investigated by the use of different types of mixing equip‐ ment, e.g., internal batch mixers or twin-screw extruders. Regardless of the mixing equipment, the transition between co-continuous to dispersed-type morphologies with well-dispersed cross-linked elastomer domains has merely been attributed to the state of gel content in the blend [64]. This basically means that for a phase transition to happen the gel content of the mixture has to reach a certain level and, if an appropriate blend formulation with welldesigned mixing conditions is set in advance, the transition will eventually happen at a similar point in the reaction regardless of the mixing equipment. In most cases with rapid reactive curing systems, the phase transformation occurs quite fast within the first one minute upon the addition of the curing system [58]. Despite this rapid onset, the final morphology of dynamically vulcanized blends depends on both the rate of the curing reaction and the total amount of shear exerted on the blend [69]. Large and to some extent coarse and interconnected vulcanized elastomeric domains have been reported for those dynamically cross-linked blends obtained in a twin-screw extruder in comparison to the ones from an internal mixer [14,69]. The observed morphological features have been mainly attributed to the fast curing reaction (due to larger viscous dissipation) and a shorter residence time in an extrusion process. A morphology comparison for dynamically vulcanized EPDM/PP blends is shown in Fig. 6. The mixture blended in an internal mixer resulted in distinct EPDM domains. On the other hand, in a twin-screw extrusion process, coarse, ruptured, and to some extent large and intercon‐ nected domains of EPDM appeared at the end of process.

Fig. 6. Atomic force microscopy images of reactive nonplasticized EPDM/PP 50/50 (wt/wt%): (a) internal mixer, (b) twin-screw extruder (Dark phase: EPDM; Bright phase: PP) (from [14]) **Figure 6.** Atomic force microscopy images of reactive nonplasticized EPDM/PP 50/50 (wt/wt%): (a) internal mixer, (b) twin-screw extruder (Dark phase: EPDM; Bright phase: PP) (from [14])

The fast curing reaction in a twin-screw extrusion process in comparison to the reaction rate in an internal mixer has readily been quantified by the gel content of the final blend [14]. In blends with especially low elastomer content, the distinguishable effect of the mixing equipment on the reaction rate can be observed in Fig. 7. One may notice that at 25 and 40 wt% of EPDM, the final gel content of the elastomeric phase is significantly lower when dynamic vulcanization is performed in an internal mixer. The fast curing reaction in a twin-screw extrusion process in comparison to the reaction rate in an internal mixer has readily been quantified by the gel content of the final blend [14]. In blends with especially low elastomer content, the distinguishable effect of the mixing equip‐ ment on the reaction rate can be observed in Fig. 7. One may notice that at 25 and 40 wt% of EPDM, the final gel content of the elastomeric phase is significantly lower when dynamic vulcanization is performed in an internal mixer.

plasticizer may also be incorporated. Already, the presence of a plasticizer is known to affect the morphological state prior to the dynamic vulcanization step. As it is discussed in the previous **Figure 7.** Gel content of the EPDM phase in both nonplasticized and plasticized dynamically vulcanized EPDM/PP blends (from [14])

section, it is expected to observe phase morphology with larger and percolated elastomeric In several reactive blends, other than the main polymeric components and the curing system, a plasticizer may also be incorporated. Already, the presence of a plasticizer is known to affect the morphological state prior to the dynamic vulcanization step. As it is discussed in the previous section, it is expected to observe phase morphology with larger and percolated elastomeric domains in plasticized blends prior to the addition of a curing system (Figs. 4 and 5). Meanwhile, whether a plasticizer could affect the rate of the cross-linking reaction may actually provide further information regarding the morphology evolution during dynamic vulcanization. Several characterization techniques such as gel content measurements and thermal analyses by differential scanning calorimetry are usually employed to evaluate a curing reaction. However, the rheological characterization of an elastomer compound by means of an oscillating disc rheometer (ODR) or a conventional rheometer is the most common technique to follow the curing behavior of an elastomer. In a conventional rheometer, one may measure the dynamic rheological properties such as the storage and loss moduli, i.e., *G'* and *G"*, and obtain the crossover point between these two quantities. The aforementioned cross‐ over is the point where the elastic properties of the material overcome the viscous ones. An easy and convenient explanation is that after the crossover point during the curing reaction, the elastomer becomes more and more capable of storing the mechanical energy rather than dissipating it by means of internal friction. As a result, this rheological point has a special importance and it is attributed to the onset of network like behavior in polymeric systems. By returning to the subject of plasticization and its effect on the cross-linking reaction, the crossover point in plasticized elastomer has been clearly shown to be reached at longer times when compared to an identical nonplasticized elastomer [66]. As an example, the effects of both temperature and plasticization on the crossover point of an EPDM containing phenolic curing system are shown in Fig. 8.

 Fig. 6. Atomic force microscopy images of reactive nonplasticized EPDM/PP 50/50 (wt/wt%): (a) internal mixer, (b) twin-screw extruder (Dark phase: EPDM; Bright phase: PP) (from [14])

**Figure 6.** Atomic force microscopy images of reactive nonplasticized EPDM/PP 50/50 (wt/wt%): (a) internal mixer, (b)

The fast curing reaction in a twin-screw extrusion process in comparison to the reaction rate in an internal mixer has readily been quantified by the gel content of the final blend [14]. In blends with especially low elastomer content, the distinguishable effect of the mixing equipment on the reaction rate can be observed in Fig. 7. One may notice that at 25 and 40 wt% of EPDM, the final gel content of the elastomeric phase is significantly lower when dynamic vulcanization is

The fast curing reaction in a twin-screw extrusion process in comparison to the reaction rate in an internal mixer has readily been quantified by the gel content of the final blend [14]. In blends with especially low elastomer content, the distinguishable effect of the mixing equip‐ ment on the reaction rate can be observed in Fig. 7. One may notice that at 25 and 40 wt% of EPDM, the final gel content of the elastomeric phase is significantly lower when dynamic

Fig. 7. Gel content of the EPDM phase in both nonplasticized and plasticized dynamically vulcanized EPDM/PP blends (from [14])

In several reactive blends, other than the main polymeric components and the curing system, a plasticizer may also be incorporated. Already, the presence of a plasticizer is known to affect the morphological state prior to the dynamic vulcanization step. As it is discussed in the previous section, it is expected to observe phase morphology with larger and percolated elastomeric

**Figure 7.** Gel content of the EPDM phase in both nonplasticized and plasticized dynamically vulcanized EPDM/PP

In several reactive blends, other than the main polymeric components and the curing system, a plasticizer may also be incorporated. Already, the presence of a plasticizer is known to affect the morphological state prior to the dynamic vulcanization step. As it is discussed in the previous section, it is expected to observe phase morphology with larger and percolated elastomeric domains in plasticized blends prior to the addition of a curing system (Figs. 4 and 5). Meanwhile, whether a plasticizer could affect the rate of the cross-linking reaction may actually provide further information regarding the morphology evolution during dynamic

performed in an internal mixer.

66 Thermoplastic Elastomers - Synthesis and Applications

blends (from [14])

vulcanization is performed in an internal mixer.

twin-screw extruder (Dark phase: EPDM; Bright phase: PP) (from [14])

**Figure 8.** Effects of plasticization and temperature on the curing behavior of EPDM (filled symbols: nonplasticized; open symbols: plasticized; the crossovers are indicated by arrows) (from [66])

Performing such measurement is sometimes impossible for a reactive blend with a complete set of ingredients, i.e., both elastomer and thermoplastic phases in the presence of a curing system and plasticizer. The difficulty lies in the fact that in most cases it is impossible to inhibit the curing reaction while mixing with high-melting-point thermoplastic phase. If one needs to understand the effect of certain ingredients, such as plasticizer, on the curing behavior of complete formulation, two logical paths could be envisaged. The first one is to design a special formulation with no interference between the mixing conditions, e.g., temperature, and those which may initiate the curing process [66]. In a specially designed reactive EPDM/PP blend, the presence of a plasticizer was shown to induce an initial delayed reaction with a subsequent rapid curing behavior [66]. The second approach toward characterizing the effect of a plasti‐ cizer on the curing behavior is based on the gel content analysis of dynamically vulcanized blends [14]. A low level of gel content in the presence of a plasticizer at the end of dynamic vulcanization, especially for the low elastomer content reactive blends is already reported in Fig. 7. These observations, both on the neat elastomer and dynamically vulcanized formula‐ tions, indicate that the presence of a plasticizer induces a delayed reaction. How could this affect the morphology development in plasticized dynamically vulcanized blends? According to the previous discussion in this section, a slightly slower but gradually cross-linking reaction could actually result in smooth and good dispersion of the elastomer without its excessive rupture. However, a delayed reaction, especially in a flow field of low extensional and/or shear deformations, may actually result in large coalesced elastomer domains, which is an undesired morphological state in the case of dynamically vulcanized blends [66]. The formation of such coalesced structure in the presence of plasticizer can be observed in Fig. 9, where a plasticized reactive EPDM/PP was subjected to an *in situ* curing and shearing at 0.1 s-1 in a rheometer. One may see that the initial co-continuous morphology was gradually transformed into elongated elastomeric domains, where afterward at later stage and longer times (Fig. 9d), large domains of EPDM were formed mainly due to coalescence of elastomeric domains. smooth and good dispersion of the elastomer without its excessive rupture. However, a delayed reaction, especially in a flow field of low extensional and/or shear deformations, may actually result in large coalesced elastomer domains, which is an undesired morphological state in the case of dynamically vulcanized blends [66]. The formation of such coalesced structure in the presence of plasticizer can be observed in Fig. 9, where a plasticized reactive EPDM/PP was subjected to an *in situ* curing and shearing at 0.1 s-1 in a rheometer. One may see that the initial co-continuous morphology was gradually transformed into elongated elastomeric domains, where afterward at later stage and longer times (Fig. 9d), large domains of EPDM were formed mainly due to coalescence of elastomeric domains.

Fig. 9. Atomic force microscopy images of a plasticized EPDM/PP 50/50 (wt/wt%) reactive blend at different shearing stages in a rheometer at 165 °C and 0.1 s-1: (a) initial morphology prior shearing and curing, (b) after 450 s, (c) after 2700 s, (d) after 7200 s (Dark phase: EPDM; **Figure 9.** Atomic force microscopy images of a plasticized EPDM/PP 50/50 (wt/wt%) reactive blend at different shear‐ ing stages in a rheometer at 165 °C and 0.1 s-1: (a) initial morphology prior shearing and curing, (b) after 450 s, (c) after 2700 s, (d) after 7200 s (Dark phase: EPDM; Bright phase: PP) (from [66])

Bright phase: PP) (from [66]) In a similar nonplasticized reactive blend that was subjected to an identical flowing condition, a complete different morphology evolution could be observed (Fig. 10). First of all, the initial morphology is comparably finer. Furthermore, the gradual shearing with the no delayed curing reaction transforms the co-continuous morphology into a structure where less coalescence could be observed. More interestingly, at longer times (Fig. 10d) the thermoplastic phase shows the In a similar nonplasticized reactive blend that was subjected to an identical flowing condition, a complete different morphology evolution could be observed (Fig. 10). First of all, the initial morphology is comparably finer. Furthermore, the gradual shearing with the no delayed curing reaction transforms the co-continuous morphology into a structure where less coales‐ cence could be observed. More interestingly, at longer times (Fig. 10d) the thermoplastic phase shows the tendency of encapsulating the elastomeric domains. This morphological transfor‐ mation is basically what is wished for when dynamic vulcanization process is in mind.

tendency of encapsulating the elastomeric domains. This morphological transformation is basically what is wished for when dynamic vulcanization process is in mind. To conclude, the ideal morphological state at the end of the dynamic vulcanization process with finely dispersed elastomeric particles could only be achieved with the finest initial morphology and with a well-designed formulation and processing parameters. All these together, in a perfectly fine-tuned mixing procedure based on a correct blending sequence with tendency of encapsulating the elastomeric domains. This morphological transformation is basically what is wished for when dynamic vulcanization process is in mind. Plasticization and Morphology Development in Dynamically Vulcanized Thermoplastic Elastomers http://dx.doi.org/10.5772/61414 69

smooth and good dispersion of the elastomer without its excessive rupture. However, a delayed reaction, especially in a flow field of low extensional and/or shear deformations, may actually result in large coalesced elastomer domains, which is an undesired morphological state in the case of dynamically vulcanized blends [66]. The formation of such coalesced structure in the presence of plasticizer can be observed in Fig. 9, where a plasticized reactive EPDM/PP was subjected to an *in situ* curing and shearing at 0.1 s-1 in a rheometer. One may see that the initial co-continuous morphology was gradually transformed into elongated elastomeric domains, where afterward at later stage and longer times (Fig. 9d), large domains of EPDM were formed

Fig. 9. Atomic force microscopy images of a plasticized EPDM/PP 50/50 (wt/wt%) reactive blend at different shearing stages in a rheometer at 165 °C and 0.1 s-1: (a) initial morphology prior shearing and curing, (b) after 450 s, (c) after 2700 s, (d) after 7200 s (Dark phase: EPDM; Bright phase: PP) (from [66])

In a similar nonplasticized reactive blend that was subjected to an identical flowing condition, a complete different morphology evolution could be observed (Fig. 10). First of all, the initial morphology is comparably finer. Furthermore, the gradual shearing with the no delayed curing reaction transforms the co-continuous morphology into a structure where less coalescence could be observed. More interestingly, at longer times (Fig. 10d) the thermoplastic phase shows the

mainly due to coalescence of elastomeric domains.

**Figure 10.** Atomic force microscopy images of a nonplasticized EPDM/PP 50/50 (wt/wt%) reactive blend at different shearing stages in a rheometer at 165 °C and 0.1 s-1: (a) initial morphology prior shearing and curing, (b) after 450 s, (c) after 2700 s, (d) after 7200 s (Dark phase: EPDM; Bright phase: PP) (from [66])

respect to all ingredients and their corresponding effects on the phase morphology evolution, will guarantee the success of the dynamic vulcanization process.

## **5. Concluding remarks**

formulation with no interference between the mixing conditions, e.g., temperature, and those which may initiate the curing process [66]. In a specially designed reactive EPDM/PP blend, the presence of a plasticizer was shown to induce an initial delayed reaction with a subsequent rapid curing behavior [66]. The second approach toward characterizing the effect of a plasti‐ cizer on the curing behavior is based on the gel content analysis of dynamically vulcanized blends [14]. A low level of gel content in the presence of a plasticizer at the end of dynamic vulcanization, especially for the low elastomer content reactive blends is already reported in Fig. 7. These observations, both on the neat elastomer and dynamically vulcanized formula‐ tions, indicate that the presence of a plasticizer induces a delayed reaction. How could this affect the morphology development in plasticized dynamically vulcanized blends? According to the previous discussion in this section, a slightly slower but gradually cross-linking reaction could actually result in smooth and good dispersion of the elastomer without its excessive rupture. However, a delayed reaction, especially in a flow field of low extensional and/or shear deformations, may actually result in large coalesced elastomer domains, which is an undesired morphological state in the case of dynamically vulcanized blends [66]. The formation of such coalesced structure in the presence of plasticizer can be observed in Fig. 9, where a plasticized reactive EPDM/PP was subjected to an *in situ* curing and shearing at 0.1 s-1 in a rheometer. One may see that the initial co-continuous morphology was gradually transformed into elongated elastomeric domains, where afterward at later stage and longer times (Fig. 9d), large domains

smooth and good dispersion of the elastomer without its excessive rupture. However, a delayed reaction, especially in a flow field of low extensional and/or shear deformations, may actually result in large coalesced elastomer domains, which is an undesired morphological state in the case of dynamically vulcanized blends [66]. The formation of such coalesced structure in the presence of plasticizer can be observed in Fig. 9, where a plasticized reactive EPDM/PP was subjected to an *in situ* curing and shearing at 0.1 s-1 in a rheometer. One may see that the initial co-continuous morphology was gradually transformed into elongated elastomeric domains, where afterward at later stage and longer times (Fig. 9d), large domains of EPDM were formed

Fig. 9. Atomic force microscopy images of a plasticized EPDM/PP 50/50 (wt/wt%) reactive blend at different shearing stages in a rheometer at 165 °C and 0.1 s-1: (a) initial morphology prior shearing and curing, (b) after 450 s, (c) after 2700 s, (d) after 7200 s (Dark phase: EPDM; Bright phase: PP) (from [66])

**Figure 9.** Atomic force microscopy images of a plasticized EPDM/PP 50/50 (wt/wt%) reactive blend at different shear‐ ing stages in a rheometer at 165 °C and 0.1 s-1: (a) initial morphology prior shearing and curing, (b) after 450 s, (c) after

In a similar nonplasticized reactive blend that was subjected to an identical flowing condition, a complete different morphology evolution could be observed (Fig. 10). First of all, the initial morphology is comparably finer. Furthermore, the gradual shearing with the no delayed curing reaction transforms the co-continuous morphology into a structure where less coalescence could be observed. More interestingly, at longer times (Fig. 10d) the thermoplastic phase shows the tendency of encapsulating the elastomeric domains. This morphological transformation is

To conclude, the ideal morphological state at the end of the dynamic vulcanization process with finely dispersed elastomeric particles could only be achieved with the finest initial morphology and with a well-designed formulation and processing parameters. All these together, in a perfectly fine-tuned mixing procedure based on a correct blending sequence with

In a similar nonplasticized reactive blend that was subjected to an identical flowing condition, a complete different morphology evolution could be observed (Fig. 10). First of all, the initial morphology is comparably finer. Furthermore, the gradual shearing with the no delayed curing reaction transforms the co-continuous morphology into a structure where less coales‐ cence could be observed. More interestingly, at longer times (Fig. 10d) the thermoplastic phase shows the tendency of encapsulating the elastomeric domains. This morphological transfor‐ mation is basically what is wished for when dynamic vulcanization process is in mind.

basically what is wished for when dynamic vulcanization process is in mind.

of EPDM were formed mainly due to coalescence of elastomeric domains.

mainly due to coalescence of elastomeric domains.

68 Thermoplastic Elastomers - Synthesis and Applications

2700 s, (d) after 7200 s (Dark phase: EPDM; Bright phase: PP) (from [66])

The morphology evolution and the state of the final phase structure have crucial effects on the physical and mechanical properties of a reactive thermoplastic/elastomer system. A successful approach toward product development of such materials usually begins by identifying the essential requirements of the final product. Accordingly, the ingredients should be carefully chosen as each fulfills a special role in the overall formulation. Thermoplastic elastomers may contain several different chemicals such as polymeric components, curing system, fillers, additives, and plasticizer. Product development could only be successful if the role and the effect of each ingredient on the phase morphology are thoroughly investigated in advance. Throughout this procedure, the rheological, interfacial, and thermal properties of complete formulations with the curing behavior have to be all considered. Only based on this informa‐ tion, an effective mixing and blending strategy can be established, which at the end will guarantee the desired target properties.

## **Author details**

Shant Shahbikian1 and Pierre J. Carreau2\*

\*Address all correspondence to: pcarreau@polymtl.ca

1 CREPEC, Chemical and Biotechnological Engineering Department, University of Sherbrooke, QC, Canada

2 CREPEC, Chemical Engineering Department, Polytechnique Montreal, Montreal, QC, Canada

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## **Environmental Degradability of Polyurethanes**

Katarzyna Krasowska, Aleksandra Heimowska and Maria Rutkowska

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/60925

## **Abstract**

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[67] G. Martin, C. Barres, P. Sonntag, N. Garois and P. Cassagnau. Morphology develop‐ ment in thermoplastic vulcanizates (TPV): Dispersion mechanisms of a pre-cross‐

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[69] P. Sengupta and J.W.M. Noordermeer. Effects of composition and processing condi‐ tions on morphology and properties of thermoplastic elastomer blends of SEBS-PPoil and dynamically vulcanized EPDM-PP-oil. J Elastom Plast. 2004; 36(4): 307-331.

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74 Thermoplastic Elastomers - Synthesis and Applications

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The growing interest in environmental issues and increasing demands to develop materials that do not burden the natural environment significantly are currently observed. In this connection many studies on polymer degradation in different environments are carried out. It is important to consider the influence of synergistic action of various factors in order to understand the environmental degradation of synthetic polymers. This requires understanding of interactions between polymer and living organisms.

This paper reviews current authors research on environmental degradation of polyurethanes.

The comparison of environmental degradability of polyurethanes in the Baltic Sea water and compost under natural weather depending conditions is presented. The environmental degradation of poly(ester-urethane) based on poly(ethylene-butyleneadipate) and poly(ester-urethane) based on poly (*ε* -caprolactone) was evaluated.

The characteristic parameters of sea water (temperature, pH, salt, and oxygen contents) and of compost (temperature, pH, moisture content, and activity of dehydrogenases) were monitored and their influence on degradation of polyur‐ ethanes was discussed.

The environmental degradability of polyurethanes was investigated by changes of weight, tensile strength, morphology, and crystallinity of polyurethanes after incubation in environment. The investigated polyurethanes were degradable in both natural environments and their environmental degradability depends on the chemical structure and the kind and conditions of environment.

**Keywords:** Poly(ester-urethane), environmental degradation, sea water, compost

© 2015 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

## **1. Introduction**

Polyurethanes are an important and versatile class of synthetic polymers that can be synthe‐ sized on a large scale and can be processed in a variety of different ways. For these reasons, polyurethanes are widely used in many aspects of modern life. They have found a widespread use in the industrial fields, for example as furniture coatings, adhesives, construction materials, flame retardants, fibers, paddings, paints, elastomeric parts, and synthetic skins [1-5]. More‐ over, over the past 40 years polyurethanes have also been used in various biomedical appli‐ cations such as vascular prostheses, artificial skin, pericardial patches, soft-tissue adhesive, drug delivery devices, and scaffolds for tissue engineering [6-15].

Polyurethane is a general term used for a class of polymers that are synthesized from three basic components: a diisocyanate, a polyol, and a chain extender. The terminal hydroxyl group allows for alternating blocks, called "segments", to be inserted into polyurethane chain. Soft segments are derived from polyols such as polyester, while the hard segments are formed from the combination of diisocyanate and a chain extender component. The chain extender is usually a small molecule with either hydroxyl or amine groups.

On the one hand, hard segments contribute to hardness, tensile strength, impact resistance, stiffness, and modulus. On the other hand, soft segments contribute to water absorption, elongation, elasticity, softness, and degradability. Hence, from the point of view of applica‐ tions, it is possible to produce various polyurethanes, which properties can be easily modified by varying structures of soft and hard segments [16].

The nature of polyurethanes chemistry is the central point for understanding why some polyurethanes are non-degradable and other undergo fast degradation. Both non-degradable and degradable polyurethanes can be designed through a proper selection of building segments. Non-degradable polyurethanes are characterized by their excellent chemical stability, abrasion resistance, and mechanical properties. Now despite the xenobiotic origins of polyurethanes, they have been found to be susceptible to degradation by naturally occurring microorganisms.

Degradable polyurethanes are generally achieved by incorporating labile and hydrolysable moieties into polymer chain. It is well known that the biodegradation of polyurethanes depends on their structure, which is conditioned by several factors such as the nature of the used polyol. It has been reported that poly(ester-urethanes) are prone to microbial degradation due to the presence of ester bonds that are known to be enzymatic hydrolysable.

Incorporation of biodegradable fillers or biodegradable aliphatic isocyanates can also enhance biodegradability of polyurethanes and then reduce negative influence on the environment. For example, using a lignin-based polyols in the polyurethanes was observed higher fungal degradation level [17].

Considering the variety of possible applications very often the degradability of polyurethanes has become an important or even a deciding factor. For example, polyurethanes used for insulation of under-water cables should have high resistance to environmental degradation. In contrast to this, biodegradable polyurethane wastes could be place for utilisation or accumulation in natural environment.

**1. Introduction**

76 Thermoplastic Elastomers - Synthesis and Applications

microorganisms.

degradation level [17].

Polyurethanes are an important and versatile class of synthetic polymers that can be synthe‐ sized on a large scale and can be processed in a variety of different ways. For these reasons, polyurethanes are widely used in many aspects of modern life. They have found a widespread use in the industrial fields, for example as furniture coatings, adhesives, construction materials, flame retardants, fibers, paddings, paints, elastomeric parts, and synthetic skins [1-5]. More‐ over, over the past 40 years polyurethanes have also been used in various biomedical appli‐ cations such as vascular prostheses, artificial skin, pericardial patches, soft-tissue adhesive,

Polyurethane is a general term used for a class of polymers that are synthesized from three basic components: a diisocyanate, a polyol, and a chain extender. The terminal hydroxyl group allows for alternating blocks, called "segments", to be inserted into polyurethane chain. Soft segments are derived from polyols such as polyester, while the hard segments are formed from the combination of diisocyanate and a chain extender component. The chain extender is usually

On the one hand, hard segments contribute to hardness, tensile strength, impact resistance, stiffness, and modulus. On the other hand, soft segments contribute to water absorption, elongation, elasticity, softness, and degradability. Hence, from the point of view of applica‐ tions, it is possible to produce various polyurethanes, which properties can be easily modified

The nature of polyurethanes chemistry is the central point for understanding why some polyurethanes are non-degradable and other undergo fast degradation. Both non-degradable and degradable polyurethanes can be designed through a proper selection of building segments. Non-degradable polyurethanes are characterized by their excellent chemical stability, abrasion resistance, and mechanical properties. Now despite the xenobiotic origins of polyurethanes, they have been found to be susceptible to degradation by naturally occurring

Degradable polyurethanes are generally achieved by incorporating labile and hydrolysable moieties into polymer chain. It is well known that the biodegradation of polyurethanes depends on their structure, which is conditioned by several factors such as the nature of the used polyol. It has been reported that poly(ester-urethanes) are prone to microbial degradation

Incorporation of biodegradable fillers or biodegradable aliphatic isocyanates can also enhance biodegradability of polyurethanes and then reduce negative influence on the environment. For example, using a lignin-based polyols in the polyurethanes was observed higher fungal

Considering the variety of possible applications very often the degradability of polyurethanes has become an important or even a deciding factor. For example, polyurethanes used for

due to the presence of ester bonds that are known to be enzymatic hydrolysable.

drug delivery devices, and scaffolds for tissue engineering [6-15].

a small molecule with either hydroxyl or amine groups.

by varying structures of soft and hard segments [16].

Degradation of polymers is determined by different factors; apart from polymer type the nature of the environment is also important.

The initial breakdown of a polymer, which is the first step of the biological degradation process, can result from physical and biological forces. Physical forces such as heating/cooling, freezing/ thawing, or wetting/drying can cause mechanical damage such as cracking of polymeric materials. The growth of many microorganisms can also cause small-scale swelling and bursting of polymeric materials. Most polymers are too large to pass through cellular mem‐ branes, so they must be depolymerized to smaller monomers before they can be adsorbed and degraded within microbial cells. The monomers, dimers, and oligomers of a polymer's repeating units are much easily degraded and mineralized because they can be assimilated through the cellular membrane and then further degraded by cellular enzymes. Two categories of enzymes are involved in the biological degradation of polymers: extracellular and intracel‐ lular depolymerases. During degradation, exoenzymes from microorganisms break down complex polymers, yielding smaller molecules of short chains that are small enough to pass semi-permeable outer bacterial membranes and then to be utilized as carbon energy sources. Under oxygen conditions, aerobic microorganisms are mostly responsible for the degradation of polymer. Biomass, carbon dioxide, and water are the final products of deterioration. As opposite to this, under anoxic conditions, anaerobic microorganisms play the main role in polymer destruction. The primary products are methane, water, and biomass [18].

Figure 1 represents the general mechanism of polymer biological degradation under aerobic conditions in a natural environment such as sea water and compost, which are studied by authors.

According to the literature microorganisms such as fungi and bacteria are involved in the degradation process of polyurethanes [18].

Generally, three types of polyurethane degradations have been identified: fungal degradation, bacterial degradation, and degradation by polyurethanase enzymes. However, polyurethanes are especially susceptible to fungal attack. Soil fungal communities are involved in polyur‐ ethanes degradation. For example, four species of fungi, namely, *Curvularia senegalensis, Fusrium solani, Aureobasidium pullulans and Cladosporium*, were obtained from soil and found to degrade ester-based polyurethanes [2, 19]. Bacteria known to degrade poly(ester-urethanes) also produce polyurethanes degrading enzymes, such as polyurethane-esterases. Two kinds of polyurethane-esterase enzymes, such as an intracellular polyurethane-esterase and an extracellular polyurethane-esterase, play predominant and various roles in polyurethane biodegradation process. The intracellular enzyme provides access to the hydrophobic poly‐ urethane surface. Then the extracellular enzyme sticks on the surface of the polyurethane. During these enzymatic actions, the ability of bacteria to adhere to the polyurethanes surface and to hydrolyse polyurethane substrates to metabolites is observed [18].

**Figure 1.** General mechanism of polymer biodegradation under aerobic conditions.

According to the literature, the degradation of polyurethanes is mostly studied in laboratories, in many cases under stable and favourable temperature conditions, providing additional nutrients to microorganisms and using highly concentrated enzymes to promote degradation [2, 4, 6, 8, 17, 20-24].

In the case of degradation in natural environment, very often the synergistic action of various factors leads to polymers degradation. Each natural environment contains different macroor‐ ganisms, microorganisms, and enzymes (in terms of species diversity and population). Different physical and chemical parameters, which have influence on rates of microbial activity, affect the rate of the degradation process. Thus, it is very interesting to estimate degradability of polyurethanes under natural weather - depending conditions.

The sea is a very complicated natural environment for degradation because microorganisms, animals, salt, sunlight, fluctuation of water, rain, etc., all play a part in degradation in nature. But it is known that a wide population of living organisms can also exist in the compost. Therefore, sea water and compost under natural weather depending conditions could be used for accumulation and utilization of polyurethane wastes.

The aim of this paper is to summarize and compare the results of environmental degradation of poly(ester-urethanes) in the Baltic Sea water [25] and in compost under natural weather depending conditions [26].

## **2. Experimental**

## **2.1. Materials**

According to the literature, the degradation of polyurethanes is mostly studied in laboratories, in many cases under stable and favourable temperature conditions, providing additional nutrients to microorganisms and using highly concentrated enzymes to promote degradation

In the case of degradation in natural environment, very often the synergistic action of various factors leads to polymers degradation. Each natural environment contains different macroor‐ ganisms, microorganisms, and enzymes (in terms of species diversity and population). Different physical and chemical parameters, which have influence on rates of microbial activity, affect the rate of the degradation process. Thus, it is very interesting to estimate

The sea is a very complicated natural environment for degradation because microorganisms, animals, salt, sunlight, fluctuation of water, rain, etc., all play a part in degradation in nature. But it is known that a wide population of living organisms can also exist in the compost. Therefore, sea water and compost under natural weather depending conditions could be used

The aim of this paper is to summarize and compare the results of environmental degradation of poly(ester-urethanes) in the Baltic Sea water [25] and in compost under natural weather

degradability of polyurethanes under natural weather - depending conditions.

**Figure 1.** General mechanism of polymer biodegradation under aerobic conditions.

for accumulation and utilization of polyurethane wastes.

[2, 4, 6, 8, 17, 20-24].

78 Thermoplastic Elastomers - Synthesis and Applications

depending conditions [26].

Two kinds of poly(ester-urethanes) were used in this work: poly(ester-urethane)A and poly(ester-urethane)B designated as PU-A and PU-B respectively. The polyurethanes were obtained by a two-step condensation reaction [27, 28].

In the first step, the prepolymers were prepared from 4,4'diphenylmethane diisocyanate /MDI/ and different polydiols. The poly(ethylene-buthylene-adipate) /PEBA/, MW=2000 was used for poly(ester-urethane)A. Poly(caprolactone)diol /PCLD/, MW=2000 was used for poly(esterurethane)B. The molar ratio of NCO:OH was 4:1 in all cases in the first step reaction. The synthesis was carried out at 80°C. Both prepolymers were further extended by 1,4-butanediol / BD/. A chain extender was added to the prepolymer in an appropriate quantity to maintain a steady NCO:OH ratio of 1.1:1. The content of hard segments (the reaction product of MDI and low molecular weight chain extender) in obtained PU-A and PU-B was 38% [25, 26].

After synthesis, the polymers obtained in a form of a sheet (thickness of 2 mm), were cut into dumbbell-shaped samples and were used for the incubation in the environments.

Before incubation in the environments, the swelling test of polyurethane samples [25] was carried out in acetone and tetrahydrofuran (THF) to estimate the partial solubility and swelling degree, which can inform us about the extent of crosslinking.

Looking at the results of swelling degree and mass loss presented in Table 1, we can notice that PU-A is soluble in THF, PU-B is partially soluble. Thus, both the studied poly(esterurethanes) differ not only in their chemical structure but in their network as well. PU-B can be partially crosslinked by allophanate bonds while PU-A is uncrosslinked.


(\*) Was calculated from the mass-ratio before and after swelling.

**Table 1.** Results of swelling test for the blind samples of polyurethanes in acetone and tetrahydrofuran.

## **2.2. Degradation environments**

The incubation of poly(ester-urethanes) took place in two natural environments:


The environmental degradation of PU-A and PU-B was carried out for up to 24 months.

During the environmental degradation in the Baltic Sea water, the samples were located in a special basket at 2 m depth under the surface of the sea, near the ship of the Polish Ship Salvage Company. The place of environmental degradation of polyurethane samples in the harbour area of the Baltic Sea is shown on Figure 2.

**Figure 2.** Place of environmental degradation of polyurethane samples in the Baltic Sea.

Degradation of polyurethane samples was also performed in the laboratory in a liquid medium containing sea water with NaN3 (0.195 g NaN3/1000cm3 ). The polyurethane samples were located in the glass aquarium, with dimensions of 40 × 40 × 20 cm, equipped with an aeration pump. The sodium azide was added to the sea water to exclude the activity of microorganisms and to evaluate the resistance of the polymers to hydrolysis.

The characteristic parameters of sea water according to the Gdynia Water Management and Meteorology Institute and of liquid medium containing sea water with NaN3 are shown in Table 2.



**Table 2.** The characteristic parameters of the Baltic Sea water and sea water with NaN3.

During the environmental degradation in the Baltic Sea water, the samples were located in a special basket at 2 m depth under the surface of the sea, near the ship of the Polish Ship Salvage Company. The place of environmental degradation of polyurethane samples in the harbour

area of the Baltic Sea is shown on Figure 2.

80 Thermoplastic Elastomers - Synthesis and Applications

**Figure 2.** Place of environmental degradation of polyurethane samples in the Baltic Sea.

containing sea water with NaN3 (0.195 g NaN3/1000cm3

and to evaluate the resistance of the polymers to hydrolysis.

[°C] pH

Temperature

Table 2.

**Months**

Degradation of polyurethane samples was also performed in the laboratory in a liquid medium

located in the glass aquarium, with dimensions of 40 × 40 × 20 cm, equipped with an aeration pump. The sodium azide was added to the sea water to exclude the activity of microorganisms

The characteristic parameters of sea water according to the Gdynia Water Management and Meteorology Institute and of liquid medium containing sea water with NaN3 are shown in

> Oxygen contents [cm3 /dm3 ]

January 2.5 8.2 9.7 6.9 18.0 8.1 February 1.6 8.2 10.3 6.4 18.5 8.0 March 3.5 8.2 10.3 6.5 19.0 8.0 April 5.1 8.4 10.0 6.8 20.0 8.0 May 13.1 8.5 8.5 6.2 21.5 8.0 June 16.5 8.3 8,0 6.3 23.0 8.0

**PARAMETERS**

**Baltic Sea water sea water with NaN3**

Salt contents [ppt]

). The polyurethane samples were

Temperature

[°C] pH

**Figure 3.** The cross-section of the compost pile prepared in natural environment [26].

The compost used in this work was formed with the dehydrated sewage sludge taken from a municipal waste treatment plant in Gdynia. The compost pile was prepared under natural conditions of sewage farm. It consisted of the sewage sludge, burnt lime, and straw. Burnt lime (0.45 kgCaO/1kg dry mass of compost) was added to ravage phatogenic bacterium and eggs parasites, to deacidificate sewage sludge, and to convert sludge to compost. The straw was added to maintain the higher temperature of the compost pile and to loosen the structure of the compost pile. The compost pile prepared under natural conditions was not adequately aerated, so it is expected that a combination of conditions from aerobic at the upper part of pile, microaerophilic in the middle part, and facultative anaerobic at the bottom of the pile could occur for microorganism growth [26]. Figure 3 represents cross-section of the compost pile in natural environment.

The characteristic parameters of the compost, such as, temperature, pH, moisture content, and activity of dehydrogenases, were measured during the degradation process and are shown in Table 3.


**Table 3.** The characteristic parameters of compost.

## **2.3. Measurements**

## *2.3.1. Characterization of the compost pile under natural conditions*

The characteristic parameters of the compost such as temperature, pH, moisture content, and activity of dehydrogenases were measured during degradation process.

## *2.3.1.1. The humidity of compost*

The moisture content of the compost was determined by drying at 105°C until constant weight was obtained.

## *2.3.1.2. The pH of the compost*

The pH of the compost was determined with a Teleko N 5172 pH-meter [26].

## *2.3.1.3. The biochemical activity of compost*

To estimate the biochemical activity of microorganisms in sludge, the activity of the dehydro‐ genases was measured by a spectrophotometric method using triphenyltetrazolium chloride (TTC). The method is based on the dehydrogenation of glucose added to the compost with a subsequent transfer of hydrogen to the colourless, biologically active compound of TTC, which undergoes a reduction to triphenyloformazan (TF). The intensity of red colour compound TF was measured using a Specol colorimeter at 490 nm [26].

## *2.3.2. Investigations of polyurethanes samples*

pile, microaerophilic in the middle part, and facultative anaerobic at the bottom of the pile could occur for microorganism growth [26]. Figure 3 represents cross-section of the compost

The characteristic parameters of the compost, such as, temperature, pH, moisture content, and activity of dehydrogenases, were measured during the degradation process and are shown in

January 6.0 5.3 49 0.0281 February 4.0 5.6 50 0.0286 May 15.0 5.7 55 0.0297 July 19.0 5.5 53 0.0328 August 22.0 5.4 56 0.0431 November 8.0 5.8 60 0.0318 December 7.0 5.9 55 0.0192

The characteristic parameters of the compost such as temperature, pH, moisture content, and

The moisture content of the compost was determined by drying at 105°C until constant weight

To estimate the biochemical activity of microorganisms in sludge, the activity of the dehydro‐ genases was measured by a spectrophotometric method using triphenyltetrazolium chloride (TTC). The method is based on the dehydrogenation of glucose added to the compost with a subsequent transfer of hydrogen to the colourless, biologically active compound of TTC, which

**PARAMETERS**

Moisture content [%]

Activity of dehydrogenases [mol mg-1 d.m.]

pile in natural environment.

82 Thermoplastic Elastomers - Synthesis and Applications

Temperature [ o

**Table 3.** The characteristic parameters of compost.

*2.3.1. Characterization of the compost pile under natural conditions*

activity of dehydrogenases were measured during degradation process.

The pH of the compost was determined with a Teleko N 5172 pH-meter [26].

C] pH

Table 3.

**Months**

**2.3. Measurements**

was obtained.

*2.3.1.1. The humidity of compost*

*2.3.1.2. The pH of the compost*

*2.3.1.3. The biochemical activity of compost*

After incubation time, the samples were taken out from the environment, washed with distilled water, and dried at room temperature until constant weight.

The environmental degradability of polyurethanes was investigated by changes of weight, tensile strength, morphology, and crystallinity of polyurethane samples after incubation in the environment.

## *2.3.2.1. The changes in the polymer surface*

The changes of polyurethanes surface were evaluated at macro- and micro scale. The views of polyurethane samples surface before and after degradation were compared. The pictures were taken before and after incubation in the environment. Microscopic observations of samples surface were done at magnification of 1:300 using the optical microscope ALPHAPHOT-2YS2- H Nikon linked to the photo camera Casio QV-2900UX. The samples were observed with and without polarizer.

## *2.3.2.2. The changes of weight*

Weight changes were determined using an electronic balance Gibertini E 42s. The weight of clean and dried samples of polyurethanes after incubation in the compost was compared with those before incubation. The percentage weight changes [%] were calculated from the weight data.

## *2.3.2.3. The changes of tensile strength*

Tensile strength [MPa] was measured at room temperature using a Tensile Testing Machine ZMGi-250 according to PN-EN ISO Standard [29].

## *2.3.2.4. The changes of thermal properties*

Thermal analysis was carried out using Perkin-Elmer Pyris 1 Differential Scanning Calorimeter equipped with Intracooler 2P. The heating scans at the rate of 5 K/min in the temperature range -65-230o C in nitrogen atmosphere were recorded [26].

## **3. Results and discussion**

## **3.1. Characteristics of the degradation environments of polyurethanes samples**

The characteristic parameters of the Baltic Sea water and compost pile under natural conditions were monitored during the environmental degradation process of polyurethanes and their influence on degradation of polyurethanes was discussed. As both biotic and abiotic param‐ eters of sea water (temperature, pH, salinity, and oxygen content) and compost (temperature, pH, moisture content, and activity of dehydrogenases) have a significant influence on the development of living microorganisms in natural environment.

Looking at the parameters presented in Tables 2 and 3, we can state, that the average temper‐ ature in the Baltic Sea was about 10°C and about 12°C in the compost pile. The temperature of both environments, depending on the weather conditions (season), had been fluctuating a lot during the experiment (from 1 to 20°C in sea water and 1°C-22°C in compost). Only the average temperature of sea water and compost during summer months (July, August) was on the similar level to the preferred for enzymatic degradation (20°C-60°C) [30].

There were significant differences in the pH values of both natural environments. The average pH in the Baltic Sea was 8.2 and 5.5 in compost pile.

The very low temperature and the alkaline pH (~8) of the Baltic Sea caused that the psychro‐ trophic bacteria could play the main role in the degradation in this environment [25].

During the winter months, we could observe the very low temperature and the highest oxygen content (February 10, 3 cm3 /dm3 ). These conditions could have an influence on the activity of oxidizing enzymes, which are responsible for oxidation. It could be expected that these conditions had an influence on the development of aerobic epilithic bacteria. The metabolism of these microorganisms probably caused the decrease of oxygen content in the summer months (August 6.5 cm3 /dm3 ) and changed the concentration of carbon dioxide in sea water. However, it was not able to change significantly the pH of sea water.

It is known that the activity of dehydrogenases depends on the degree growth of microorgan‐ ism populations, which are producing enzymes involved in degradation process. During the degradation time, the activity of dehydrogenases had been changing and depending on both biotic and abiotic conditions in this environment.

The weather, as well as respiration of microorganisms, had an influence on fluctuation of moisture content in the compost. With decreasing moisture content, lower absolute value of the activity of dehydrogenases was observed.

The rather low temperatures (below 20°C) and slightly acid pH (~6) of compost under natural weather - depending conditions caused that psychrotrophic acidophilic microorganisms (fungi) could play the main role in the degradation process [26].

Considering the characteristic abiotic parameters of sea water and compost presented in Tables 1 and 2 and the different microbial communities (fungi in compost and bacteria in sea water), we could expect the different rate of degradation of polyurethanes in these two natural environments.

Incubation of polyurethane samples in the laboratory in a liquid medium containing sea water with NaN3 (Table 2) was performed in a stable temperature (about 20°C) and under alkaline conditions (pH about 8).

## **3.2. Evaluation of the changes of polyurethane samples during environmental degradation**

influence on degradation of polyurethanes was discussed. As both biotic and abiotic param‐ eters of sea water (temperature, pH, salinity, and oxygen content) and compost (temperature, pH, moisture content, and activity of dehydrogenases) have a significant influence on the

Looking at the parameters presented in Tables 2 and 3, we can state, that the average temper‐ ature in the Baltic Sea was about 10°C and about 12°C in the compost pile. The temperature of both environments, depending on the weather conditions (season), had been fluctuating a lot during the experiment (from 1 to 20°C in sea water and 1°C-22°C in compost). Only the average temperature of sea water and compost during summer months (July, August) was on the

There were significant differences in the pH values of both natural environments. The average

The very low temperature and the alkaline pH (~8) of the Baltic Sea caused that the psychro‐

During the winter months, we could observe the very low temperature and the highest oxygen

oxidizing enzymes, which are responsible for oxidation. It could be expected that these conditions had an influence on the development of aerobic epilithic bacteria. The metabolism of these microorganisms probably caused the decrease of oxygen content in the summer

It is known that the activity of dehydrogenases depends on the degree growth of microorgan‐ ism populations, which are producing enzymes involved in degradation process. During the degradation time, the activity of dehydrogenases had been changing and depending on both

The weather, as well as respiration of microorganisms, had an influence on fluctuation of moisture content in the compost. With decreasing moisture content, lower absolute value of

The rather low temperatures (below 20°C) and slightly acid pH (~6) of compost under natural weather - depending conditions caused that psychrotrophic acidophilic microorganisms

Considering the characteristic abiotic parameters of sea water and compost presented in Tables 1 and 2 and the different microbial communities (fungi in compost and bacteria in sea water), we could expect the different rate of degradation of polyurethanes in these two natural

Incubation of polyurethane samples in the laboratory in a liquid medium containing sea water with NaN3 (Table 2) was performed in a stable temperature (about 20°C) and under alkaline

). These conditions could have an influence on the activity of

) and changed the concentration of carbon dioxide in sea water.

trophic bacteria could play the main role in the degradation in this environment [25].

development of living microorganisms in natural environment.

pH in the Baltic Sea was 8.2 and 5.5 in compost pile.

84 Thermoplastic Elastomers - Synthesis and Applications

/dm3

However, it was not able to change significantly the pH of sea water.

(fungi) could play the main role in the degradation process [26].

/dm3

biotic and abiotic conditions in this environment.

the activity of dehydrogenases was observed.

content (February 10, 3 cm3

months (August 6.5 cm3

environments.

conditions (pH about 8).

similar level to the preferred for enzymatic degradation (20°C-60°C) [30].

The environmental degradation of polyurethanes in sea water and compost was evaluated visually at first. Figure 4 represents the surface view at macro scale of investigated poly(esterurethanes) before and after degradation in the Baltic Sea water and compost under natural weather depending conditions.

**Figure 4.** Macroscopic images of poly(ester-urethanes) after environmental degradation.

The blind samples of PU-A and PU-B are beige or white and opaque. The both poly(esterurethane) samples incubated in natural environments are characterized by the flaws at the surface that gradually grew into microcracks, eventually breaking the samples. At the end of the experiment in natural environments, the surface of poly(ester-urethanes) samples is rough and cracked with visible black areas after incubation in sea water. While after incubation in sea water with NaN3 the surface is only matt.

The environmental degradation of poly(ester-urethane) samples in the Baltic Sea water and the compost under natural weather-depending conditions was also evaluated on the basis of changes of surface morphology. After incubation in both natural environments, the poly(esterurethane) samples were not homogeneously destroyed over the whole polymer surface and there were different images depending on the place. The photomicrographs repeated images observed under the reflected microscope equipped with polarizer were done.

Microscopic observations after incubation in the natural environments have shown vulnera‐ bility of PU-A and PU-B to the microbiological attack. The changes of surfaces of both poly(ester-urethanes) are comparable. After incubation of poly(ester-urethanes) in natural environments the deterioration of the surfaces has been observed. The photomicrographs of PU-B are presented in Figure 5.

**Figure 5.** Microscopic images of PU-B observed under optical microscope with polarizer before and after environmen‐ tal degradation.

The changes in birefringence of PU-B observed under a polarized optical microscope are noticeable (Figure 5). The blind sample of PU-B is slightly crystalline (the bright elements on the surface before degradation). It is known that the degradation process of polyesters proceeds in two stages [31]. During the first stage of the environmental degradation of PU-B, the gradually increase of bright elements on the surface was observed. This increase of bright elements might be an evidence of increase in crystallinity as a result of the degradation amorphous phase (random hydrolytic scission of ester bonds) [25, 26]. The second stage started after degradation of the major part of the amorphous phase. At the end of incubation in the compost, we could see distinct loss of bright elements due to the degradation of the crystalline phase (Figure 5).

Comparing microscopic observations of PU-B after environmental degradation in the Baltic Sea water and in compost we can state that PU-B based on poly(ε-caprolactone) is more vulnerable to degradation in compost than in sea water. Thus, the higher deterioration of the PU-B surface have been observed. It is because of the fragment of poly(ε-caprolactone) in the main chain of PU-B, which is susceptible to microbial degradation. Moreover, in this natural environment there were conditions favourable for growth of fungi. It is well known that fungal communities are involved in poly(ester-urethanes) biodegradation [2, 6, 17, 23, 24, 32].

Microscopic observations after incubation in the natural environments have shown vulnera‐ bility of PU-A and PU-B to the microbiological attack. The changes of surfaces of both poly(ester-urethanes) are comparable. After incubation of poly(ester-urethanes) in natural environments the deterioration of the surfaces has been observed. The photomicrographs of

**Figure 5.** Microscopic images of PU-B observed under optical microscope with polarizer before and after environmen‐

The changes in birefringence of PU-B observed under a polarized optical microscope are noticeable (Figure 5). The blind sample of PU-B is slightly crystalline (the bright elements on the surface before degradation). It is known that the degradation process of polyesters proceeds in two stages [31]. During the first stage of the environmental degradation of PU-B, the gradually increase of bright elements on the surface was observed. This increase of bright elements might be an evidence of increase in crystallinity as a result of the degradation amorphous phase (random hydrolytic scission of ester bonds) [25, 26]. The second stage started after degradation of the major part of the amorphous phase. At the end of incubation in the compost, we could see distinct loss of bright elements due to the degradation of the crystalline

Comparing microscopic observations of PU-B after environmental degradation in the Baltic Sea water and in compost we can state that PU-B based on poly(ε-caprolactone) is more

PU-B are presented in Figure 5.

86 Thermoplastic Elastomers - Synthesis and Applications

tal degradation.

phase (Figure 5).

Susceptibility of poly(ester-urethanes) to environmental degradation was evaluated based on weight changes [%] of polymer samples after incubation. The results of the weight changes of the PU-A and PU-B after incubation in natural environments are presented in Figure 6.

**Figure 6.** The weight losses [%] of PU-A and PU-B after environmental degradation.

The obtained results indicate that both poly(ester-urethanes) degrade in natural environments such as the Baltic Sea water and compost. It confirms susceptibility of poly(esters-urethane) to biological degradation.

Generally, the results presented in Figure 6 indicate that the degradability of poly(esterurethanes) depends on their chemical structure and the kind of environment.

In the Baltic Sea water, where the low temperature and the alkaline pH were favourable for the development of aerobic bacteria, uncrosslinked PU-A based on poly(ethylene-butyleneadipate) was more prone to degradation than PU-B based on poly(ε-caprolactone). At the end of incubation in the Baltic Sea water (12 months) the changes of weight of PU-A were much higher (19%) than of PU-B (4%). It could be explained by the different networks of those polyurethanes. Unexpectedly, the introduction of a fragment of poly (*ε* -caprolactone) to the main chain of PU-B did not lead to the increase of its environmental degradability [25].

In contrary, to this there are results of the degradation of poly(ester-urethanes) in compost under natural weather-depending condition. The strongest effect of environmental degrada‐ tion in compost was observed for slightly crosslinked PU-B based on poly (*ε* -caprolactone) than not crosslinked PU-A based on poly(ethylene-butylene-adipate). It could be mainly explained by degradation of poly (*ε* -caprolactone) in this biotic environment as a result of enzymatic hydrolysis of ester bonds susceptible to fungal degradation [26]. Considering the parameters of compost pile under natural weather-depending conditions (temperature and pH) it could be stated that the psychrotrophic acidofilic microorganisms (fungi) were respon‐ sible for the level of degradability of poly (ester-urethanes).

According to the literature polyester degradation occurs primarily by chain scission in main chain of polymer and can be induced by enzymatic hydrolysis. Enzymes can attack on the surface poly(ε-caprolactone) segments of polyurethane, degrading them to smaller molecular units via hydrolytic attack. Then surface erosion takes place with farther hydrolysis process erosion. At the end, hydrolysis rate decreases after the consumption of the amorphous materials by microorganisms [6, 33, 34].

This is why the higher weight losses and deterioration of the surface have been observed during environmental degradation in compost [26].

During incubation in the sea water with NaN3 in the laboratory, the weight changes of poly(ester-urethane)s are insignificant even though the temperature was higher than in the natural environment. This was due to the absence of microorganisms in sea water with sodium azide [25].

Changes of mechanical properties of both poly(ester-urethanes) were checked by the meas‐ urement of the tensile strength before and after environmental degradation. The results are presented in Figure 7.

It is interesting to note that for blind samples, the higher tensile strength is observed for PU-B sample, which is due to its partial crosslinking. The rates of the changes in the mechanical properties (Figure 7) resemble the rate of the mass loss (Figure 6) and the changes of surface poly(ester-urethanes) (Figures 4 and 5). The data in Figure 7 show that the tensile strength of poly(ester-urethanes) had been decreased during the incubation time in sea water and compost [25, 26]. After 6 months of environmental degradation, only the samples of PU-A incubated in sea water were torn up into pieces, whereas the tensile strength of the other samples degraded in both environments could still be estimated. Probably, microorganisms existed in sea water, such as psychrotrophic bacteria, caused the breaking of polymer samples resulting in the fragmentation. After 12 months of environmental degradation all poly(ester-urethanes) lost the tensile strength.

The loss of tensile strength, discoloration, and cracking observed for environmental degraded poly(ester-urethanes) are typical for the effects of degradation of poly(ester-urethanes) as a result of microorganisms activity.

The results of thermal analysis of poly(ester-urethane) samples are shown in Table 4 and Figure 8a and b and they are in the agreement with microscopic observations (Figure 5).

**Figure 7.** The tensile strength [MPa] of PU-A and PU-B before and after environmental degradation.

In contrary, to this there are results of the degradation of poly(ester-urethanes) in compost under natural weather-depending condition. The strongest effect of environmental degrada‐ tion in compost was observed for slightly crosslinked PU-B based on poly (*ε* -caprolactone) than not crosslinked PU-A based on poly(ethylene-butylene-adipate). It could be mainly explained by degradation of poly (*ε* -caprolactone) in this biotic environment as a result of enzymatic hydrolysis of ester bonds susceptible to fungal degradation [26]. Considering the parameters of compost pile under natural weather-depending conditions (temperature and pH) it could be stated that the psychrotrophic acidofilic microorganisms (fungi) were respon‐

According to the literature polyester degradation occurs primarily by chain scission in main chain of polymer and can be induced by enzymatic hydrolysis. Enzymes can attack on the surface poly(ε-caprolactone) segments of polyurethane, degrading them to smaller molecular units via hydrolytic attack. Then surface erosion takes place with farther hydrolysis process erosion. At the end, hydrolysis rate decreases after the consumption of the amorphous

This is why the higher weight losses and deterioration of the surface have been observed

During incubation in the sea water with NaN3 in the laboratory, the weight changes of poly(ester-urethane)s are insignificant even though the temperature was higher than in the natural environment. This was due to the absence of microorganisms in sea water with sodium

Changes of mechanical properties of both poly(ester-urethanes) were checked by the meas‐ urement of the tensile strength before and after environmental degradation. The results are

It is interesting to note that for blind samples, the higher tensile strength is observed for PU-B sample, which is due to its partial crosslinking. The rates of the changes in the mechanical properties (Figure 7) resemble the rate of the mass loss (Figure 6) and the changes of surface poly(ester-urethanes) (Figures 4 and 5). The data in Figure 7 show that the tensile strength of poly(ester-urethanes) had been decreased during the incubation time in sea water and compost [25, 26]. After 6 months of environmental degradation, only the samples of PU-A incubated in sea water were torn up into pieces, whereas the tensile strength of the other samples degraded in both environments could still be estimated. Probably, microorganisms existed in sea water, such as psychrotrophic bacteria, caused the breaking of polymer samples resulting in the fragmentation. After 12 months of environmental degradation all poly(ester-urethanes) lost

The loss of tensile strength, discoloration, and cracking observed for environmental degraded poly(ester-urethanes) are typical for the effects of degradation of poly(ester-urethanes) as a

The results of thermal analysis of poly(ester-urethane) samples are shown in Table 4 and Figure

8a and b and they are in the agreement with microscopic observations (Figure 5).

sible for the level of degradability of poly (ester-urethanes).

materials by microorganisms [6, 33, 34].

88 Thermoplastic Elastomers - Synthesis and Applications

azide [25].

presented in Figure 7.

the tensile strength.

result of microorganisms activity.

during environmental degradation in compost [26].


**Table 4.** Melting temperature (Tm) and melting enthalpy (∆H) of crystallites made of soft (SS) and hard segment (HS) determined from DSC scans for PU-A and -B samples before and after incubation in compost.

The DSC analysis of PU-A and PU-B revealed the differences in their phase composition (Figure 8). Both poly(ester-urethanes) before degradation showed the little presence of crystal phases. The blind sample of PU-A contains mainly crystals made of hard segments, which is indicated by the small melting peaks at high temperatures (140°C and 184°C in Figure 8a). In the case of the blind sample of PU-B, only small melting peak at low temperatures (77°C in Figure 8b) was noticed. It is corresponding to the melting of soft segments crystals.

**Figure 8.** The DSC curves for poly(ester-urethanes): **a)** PU-A and **b)** PU -B before and after incubation in the compost for 6, 12, and 24 months [26]

Due to incubation in the compost, the evident increase in crystallinity in both poly(esterurethanes) was observed. However, in the case of PU-A, crystal phases made of both hard and soft segments appeared, but in the case of PU-B, mainly crystal phase made of soft segments (compare Figure 8a and b and Table 4) appeared. The observation that higher crystallinity develops in PU-A may be explained by its uncrosslinked structure. In the case of PU-B, the partial crosslinking constraints at some point of crystallization are especially of hard segments. Differences in crystallinity in both poly (ester-urethanes) seem to correspond with mechanical properties. The smaller decrease of mechanical properties of PU-A is due to the reinforcing effect of higher crystallinity [26].

## **4. Conclusions**

Currently, the information concerning microbial degradation of polyurethanes in the natural environment is still limited. In this study, the ability of the Baltic Sea water and compost to degrade poly(ester-urethanes) were accessed. The achieved results pointed out that the poly(ester-urethane) based on poly(ethylene-butylene-adipate) and poly(ester-urethane) based on poly (*ε* -caprolactone) are susceptible to environmental degradation under natural weather-depending conditions.

Generally the stronger effect of environments degradation—the higher weight losses and deterioration of the surface of poly(ester-urethanes)—has been observed during environmen‐ tal degradation in compost, than in sea water where the conditions were favourable for the development of aerobic bacteria.

In the Baltic Sea water uncrosslinked poly(ester-urethane) based on poly(ethylene-butyleneadipate) was very prone to degradation. Whereas slightly crosslinked poly(ester-urethane) was moderately resistant to degradation in sea water, even though it has a fragment of poly (*ε* -caprolactone) in the main chain. In contrary to this, there are results of the degradation of poly(ester-urethanes) in compost. In this environment, slightly crosslinked poly(esterurethane)B, based on poly(ε-caprolactone) was more degradable, than not being crosslinked poly(ester-urethane)A, based on poly(ethylene-buthylene adipate). Finally, the higher weight losses and deterioration of surface have been observed. It could be mainly explained by degradation of poly (*ε* -caprolactone) in this biotic environment as a result of enzymatic hydrolysis of ester bonds susceptible to fungal degradation. The psychrotrophic acidofilic microorganisms (fungi) were responsible for the level of degradability of poly(ester-urethanes) in compost.

Due to incubation in the compost, there is the evident increase in crystallinity in both poly(es‐ ter-urethanes). Differences in crystallinity are corresponding with mechanical properties as reinforcing effect of crystal phase in poly(ester-urethanes). The smaller decrease in mechanical properties of poly(ester-urethane)A, than poly(ester-urethane)B may be due to its higher crystallinity. The loss of tensile strength, discoloration, and cracking observed for environ‐ mental degraded poly(ester-urethanes) are typical for the effects of degradation of poly(esterurethanes) as a result of microorganisms activity.

The environmental degradation process of poly(ester-urethanes) indicates that the degrada‐ tion in natural environment such as the Baltic Sea water and compost is the result primarily of enzymatic hydrolysis of ester bonds, then crystallinity and network structure of poly (esterurethanes). The rate of environmental degradation process of poly(ester-urethanes) also is depended on the kind and conditions of natural environment.

## **Author details**

Due to incubation in the compost, the evident increase in crystallinity in both poly(esterurethanes) was observed. However, in the case of PU-A, crystal phases made of both hard and soft segments appeared, but in the case of PU-B, mainly crystal phase made of soft segments (compare Figure 8a and b and Table 4) appeared. The observation that higher crystallinity develops in PU-A may be explained by its uncrosslinked structure. In the case of PU-B, the partial crosslinking constraints at some point of crystallization are especially of hard segments. Differences in crystallinity in both poly (ester-urethanes) seem to correspond with mechanical properties. The smaller decrease of mechanical properties of PU-A is due to the reinforcing

(a) (b)

**Figure 8.** The DSC curves for poly(ester-urethanes): **a)** PU-A and **b)** PU -B before and after incubation in the compost

Currently, the information concerning microbial degradation of polyurethanes in the natural environment is still limited. In this study, the ability of the Baltic Sea water and compost to degrade poly(ester-urethanes) were accessed. The achieved results pointed out that the poly(ester-urethane) based on poly(ethylene-butylene-adipate) and poly(ester-urethane) based on poly (*ε* -caprolactone) are susceptible to environmental degradation under natural

Generally the stronger effect of environments degradation—the higher weight losses and deterioration of the surface of poly(ester-urethanes)—has been observed during environmen‐ tal degradation in compost, than in sea water where the conditions were favourable for the

In the Baltic Sea water uncrosslinked poly(ester-urethane) based on poly(ethylene-butyleneadipate) was very prone to degradation. Whereas slightly crosslinked poly(ester-urethane) was moderately resistant to degradation in sea water, even though it has a fragment of poly (*ε* -caprolactone) in the main chain. In contrary to this, there are results of the degradation of

effect of higher crystallinity [26].

weather-depending conditions.

development of aerobic bacteria.

**4. Conclusions**

for 6, 12, and 24 months [26]

90 Thermoplastic Elastomers - Synthesis and Applications

Katarzyna Krasowska, Aleksandra Heimowska and Maria Rutkowska\*

\*Address all correspondence to: m.rutkowska@wpit.am.gdynia.pl

Gdynia Maritime University, Faculty of Entrepreneurship and Quality Science, Department of Chemistry and Industrial Commodity Science, Gdynia, Poland

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