Tailored Nanostructured Thermoplastic Elastomers from

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Industrial & Engineering Chemistry Research

Tailored nanostructured thermoplastic elastomers from polypropylene and fluoroelastomer: Morphology and functional properties Shib Shankar Banerjeea and Anil K. Bhowmickb* a

Department of Materials Science and Engineering, Indian Institute of Technology Patna, Patna 800013, India. b Rubber Technology Centre, Indian Institute of Technology Kharagpur, Kharagpur 721302, India.

Abstract Conventional thermoplastic elastomer from rubber – plastic blends has rubber domains of a few micrometers dispersed in a plastic matrix. Nanotailoring of novel thermoplastic elastomeric blends from polypropylene (PP) and fluoroelastomer (FKM) by a special green technique generated unique morphology where FKM was dispersed in the nanometer range (50-70 nm) into the continuous PP matrix. The evolution of unique nanostructure morphology was explained with the help of viscoelastic drop break up mechanism. The nanostructured blends exhibited superior physical properties and outstanding oil swelling resistance and thermal stability, which increased on vulcanizing the rubber phase in-situ. For example, volume swell of nanostructured PP/FKM blends was 8-10 % in ASTM oil #3 at 100 °C for 72 hr, which was the lowest in this type of thermoplastic elastomers. Insights into how the unique morphology influenced the functional properties of thermoplastic elastomers was provided. These new thermoplastic elastomers can be used in the automotive industry, particularly as a liner material for car hoses where superior flexibility, oil resistance and heat resistance are desirable properties. Keywords: polypropylene, fluoroelastomer, thermoplastic elastomer, nanostructured materials, morphology, functional properties *Corresponding author: TEL: (91-3222) 283180; FAX: (91-3222)-220312, E-mail: [email protected] (Anil K. Bhowmick) 1 ACS Paragon Plus Environment


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1. Introduction Thermoplastic elastomers (TPEs) are a particular group of high performance materials which can be processed like thermoplastics, but exhibit flexibility and resilience of rubbers at room temperature.1-4 On vulcanization of the rubber during simultaneous mixing with plastics, thermoplastic vulcanizates (TPV) are obtained, which have improved elasticity and tension set properties.5-9 From the structural point of view, the shape and size of the dispersed rubber domain in the plastic matrix has a significant role on processability and properties of TPEs and TPVs. Most studies have reported that the rubber phase is dispersed in the plastic matrix with a diameter ranging from 0.5–3.0 µm.10,11 Attempts were made in the past to reduce the size of the dispersed rubber particles which may lead to a significant improvement of properties of TPEs and TPVs, but met with limited success.12 Also, preparation of new materials by a green approach (without the use of any solvents) is an area of current research interest. In addition, stringent demand of the automotive industry dictates development of materials having improved properties. Based on the above background, the current investigation is undertaken. Fluoroelastomers (FKM) are exceptionally stable and have outstanding heat, light, ozone, solvent, and chemical resistance compared to all rubbers. These lead to huge important applications in automotive, high performance sealing and gasket industries. 13 On the other hand, polypropylene (PP) has attractive characteristics, including light weight, low cost, high heat distortion temperature, excellent processability, mechanical properties, recyclability, etc. all of which make this an ideal candidate for commodity and industrial applications.14,15 Therefore, blending of fluorocarbon rubber with polypropylene thermoplastic to generate thermoplastic elastomer is a good challenge. Excellent properties of PP and FKM and earlier reports on TPEs and TPVs motivate us to ask the following questions:

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1. Can we make thermoplastic elastomers using PP and FKM, which has not been reported till date? It is noteworthy that simple rubber–plastic blends are not TPEs, even after vulcanization of the rubber phase. A TPEs/TPVs should have high elongation (>100 %) and low set () and weight average diameter (< dw >) 20 were calculated using the following equation.

dn =

∑n d ∑n i

i

(1)

i

dw

∑n d = ∑n d i

2

i

i

(2) i

where , ni is the number of particles in class interval i with diameter di. 2.4.2. Atomic Force Microscopy (AFM). Intermittent Contact Mode Atomic Force Microscopy, ACAFM (Agilent 5500 Scanning Probe Microscope, USA) was carried out to analyze the microstructure of the TPEs and TPVs. The resonance frequency and force constant of the tip were 130-280 kHz and 48 N/m respectively.



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2.4.3. Transmission electron microscopy (TEM). The nanostructure morphology of PP/FKM blends was confirmed using TEM (JEOL 1210), operated at 100 kV. Ultra-thin sections were cryogenically microtomed at -100 °C using a Reichert Ultracut S ultramicrotome from injection molded specimen. These sections were then vapour stained with OsO4 solution for at least 4 h in order to differentiate in the blend components.

2.4.4. Mechanical properties. Tensile tests were performed according to ISO 527-2-5A specification by using dumb-bell shaped test pieces which were prepared using micro-injection molding. The tests were carried out in a Zwick Universal Testing Machine (UTM Z010) at room temperature at a test speed of 100 mm/min. The result reported here is the average of five samples from the same batch. Hardness (Shore D) of the pristine polymers and their blends were measured by DIN ISO 7619 test method at room temperature. Tension set was measured using UTM Z010. Specimen was extended up to 50 % elongation in the tensile direction at a rate of 20 mm/min and kept that position for 10 min. The change in length after 48 hr at room temperature was measured.

2.4.5. Thermogravimetric analysis (TGA). Thermal stability of the pristine polymers and their blends were determined by TGA analysis using TA instruments (TGA Q600). The ramp rate during the experiment was set at 20 °C/min where ~5 mg of the sample was programmed to be heated under nitrogen atmosphere up to 600 °C. The data were analyzed by TA universal analysis software. T0 (onset of degradation) and Tmax (temperature at which maximum degradation takes place) were measured for each sample.


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2.4.6. Dynamic Mechanical Analysis (DMA). DMA study was performed using a TA Instruments (model Q800) in tension mode. The samples in the form of rectangular bars were prepared using micro-injection molding with dimensions of 16 mm x 9.8 mm x 1 mm. The experiments were conducted at a constant frequency of 1 Hz, at a heating rate of 2 °C/min and a strain amplitude of 10 μm over a temperature range of - 50 °C to + 50 °C.

2.4.7. Oil swelling test. Circular test specimens of diameter 25 mm (1.4 mm thickness) were obtained from microinjection molding. These were weighed accurately and immersed in ASTM oil #3 at 100 °C for 72 h. The specimens were removed from oil, blotted and weighed in a glass-stoppered bottle. The volume swelling for the specimens was obtained as follows 21 W ρ q − 1 =  2 − 1 C  W1  ρ S

(3)

where, q is the ratio of swollen volume to original unswollen volume. W1 and W2 are specimen weights before and after swelling respectively, ρc is the density of composition and ρs is the density of oil. From the known density of ASTM oil #3 and density of the composition (calculated from the volume fraction of plastic and rubber phases), % volume swell was calculated.

3. RESULTS AND DISCUSSION 3.1. Mixing Torque. Mixing torque-time relationship of the blends was recorded from Haake rheomixer at 190 °C and 100 rpm rotor speed. Figure 1a-b shows the mixing torque behavior of the PP/FKM blends. For the blend system, PP was first added into the internal mixing chamber leading to



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a sharp increase in the torque due to resistance exerted on the rotors by the unmelted PP followed by a decrease in torque rapidly as the PP melted. After 2 min of mixing, FKM was charged into the mixing chamber and the torque increased again due to the initial high viscosity of FKM. As the FKM blended with PP and both the materials were melted and softened, the torque started to decrease and reached a steady-state value after a certain time of mixing. In order to understand the theoretical torque, the calculation was done by using the following equation: Q = Q1φ1 + Q2φ2

(4)

where Q is the torque of the blend, Q1 and Q2 are the torque values of the plastic and the rubber respectively. φ1 is the volume fraction of the plastic and φ2 is the volume fraction of the rubber. The theoretical and experimental values of torque for all the blend compositions were compared (see Table S1, supporting information). It is interesting to note that the theoretical torque values were higher for all the compositions and the deviation increased with increase of the rubber content in the blends. For example, deviation of the torque was ~ 55 % when the rubber content was 40 % by weight in the blend. This negative deviation might be due to the local phase segregation producing a nanometric lubricated interface that generated the macroscopic velocity discontinuity between the two phases. As a result of this, slippage occurred between the blend components and the torque decreased. On vulcanization of the rubber phase in-situ during mixing (dynamic vulcanization), the torque value further increased after 3 min exponentially and reached a plateau after 5 min (Figure 1b). For example, the torque value of TPE at 5 min of mixing, when the torque curve reached a steady-state, was ~ 2 Nm which increased to ~ 3.5 Nm on vulcanization of the rubber phase in-situ. The increase in torque is due to increase in crosslink density of the rubber phase, which gives resistance to the moving rotor. 8 ACS Paragon Plus Environment

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Figure 1. Torque vs. mixing time at 190 °C and 100 rpm rotor speed of (a) PP/FKM blends and (b) TPE vs. TPV.

3.2. Morphological analysis. Figure 2a-d shows the FESEM phase morphology of PP/FKM blends (TPEs and TPVs) after injection molding. FKM was observed as dispersed white particles embedded in the continuous PP matrix. Two phase morphological structure is clearly reflected from the micrographs. It was found that the viscosity of FKM was higher than PP (Figure S1, supporting information). Therefore, lower viscosity of PP phase will tend to form the matrix which is further confirmed from theoretical analysis of complex modulus (Kerner’s model) and etched FESEM morphology (discussed later). It should be mentioned that the PP/FKM thermoplastic elastomeric blends exhibited fine nanostructures of elastomeric domains in thermoplastic polypropylene matrix for the composition range of 80 PP/20 FKM to 20 PP/80 FKM (w/w) after injection molding (Figure 2a-d). The formation of nanostructure of the dispersed rubber domains is clearly reflected from the magnified FESEM images (see inset of Figure 2a-d). The number average particle diameter, dn and weight average particle diameter, 9

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dw, of the dispersed rubber phase varied between 40-60 nm and 70-90 nm respectively. It was larger with higher FKM content.

A much finer morphology with a dispersed

nanometer size rubber phase was obtained on vulcanization of the rubber phase in-situ (dynamic vulcanization) (Figure 2b′). Dispersion of the rubber was much more uniform after dynamic vulcanization as compared to the unvulcanized blend (see inset of Figure 2b and Figure 2b′). In order to have insight into structure development of the dispersed rubber phase, FESEM micrographs of the blends were captured before injection molding for each composition (Figure 2e-h). It is interesting that size of the dispersed rubber phase was found to be in the micron range (0.3 - 0.4 µm) before injection molding. By inspections of large number of samples, it was also shown that number average particle diameter, dn and weight average particle diameter, dw increased slightly with an increase of the FKM content. For example, dn of the rubber phase of the 80PP/20 FKM (w/w) blend composition was 0.3 µm which increased to 0.4 µm for the 20PP/80 FKM (w/w) blend composition. The morphology of polymer blends depends on several parameters, such as blend composition, miscibility between the components, viscosity ratio, elasticity, shear force, interfacial modification, etc.22-24 Here, shear rate plays a crucial role in the development of nano-structured rubber domain in the plastic matrix, which has been explained in detail in the forthcoming section. It is worth mentioning here that the finer nano-dispersion of elastomeric domains in thermoplastic matrix after injection molding leads to the improvement of mechanical properties in comparison with the specimen before injection molding (discussed later). The nanometer dispersion or phase structure of the rubber in thermoplastic matrix might indicate lower interfacial tension between the two phases. In fact, the most of the commercial TPEs and TPVs based on various rubber-plastic combinations usually have the dimension of rubber in the micron range and most of them are made by using compression molding. Here, micro-


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injection molding was adopted to fabricate the nanostructured thermoplastic elastomers and vulcanizates.

Figure 2. FESEM micrographs of PP/FKM blends (a) 20/80 (w/w) TPE-after injection molding, (b) 40/60 (w/w) TPE-after injection molding, (b′) 40/60 (w/w) TPV-after injection molding (c) 60/40(w/w) TPE after injection molding, (d) 80/20 (w/w) TPE after injection molding. (e) 20/80 (w/w) TPE-before injection molding, (f) 40/60 (w/w) TPE-before injection molding, (f′) 40/60 (w/w) TPV-before injection molding (g) 60/40(w/w) TPE before injection molding and (h) 80/20 (w/w) TPE before injection molding. In order to confirm the nanostructure morphology of thermoplastic elastomeric blends, additional investigations were conducted using AFM and TEM. Figure 3a-a′ reflects the 11 ACS Paragon Plus Environment


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AFM image of 40/60 PP/FKM TPE and TPV. The light-yellow regions in the topography images represent rubber particles, which are dispersed in the PP matrix. The dimension of rubber phase was found to be 50-80 nm, as noted from the inspection of a large number of samples. Figure 3b-b′ shows the high magnification TEM photomicrographs of 40/60 PP/FKM TPE and TPV, which revealed the nanodomain structure of the rubber in the plastic matrix (the rubber phase appeared as black domains within the white PP matrix). It is important to mention here that in-situ vulcanization of the rubber phase during melt mixing with the plastic leads to more stable and uniform dispersion of FKM in PP matrix. This observation is in line with the earlier FESEM findings. In order to further substantiate the dimension of the dispersed phase, FKM phase was selectively etched with methyl ethyl ketone at room temperature for 72 h. Figure 3c-c′ reflects the dark holes (corresponding to the rubber phase) distributed in the PP matrix. The average diameter of the hole was in the range of 70-90 nm, which supports the previous microscopic findings. It is important to mention here that the cross-section of the blends was also observed as a sphere-like domain of the rubber in the PP matrix (due to etching), as shown in Figure 3c.


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Figure 3. (a-a′) AFM images of 40/60 PP/FKM (w/w) TPE and TPV, (b-b′) TEM images of 40/60 PP/FKM (w/w) TPE and TPV and (c-c′) Etched FESEM morphology of 40/60 PP/FKM (w/w) TPE at two different magnifications.



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3.3. Theoretical analysis of nanostructure formation. The deformation and breakup of the dispersed rubber phase in the plastic matrix occur upon shear in a mixer. At a critical shear in the mixer, when the interfacial force is equivalent to the deformation rate, small particles are generated. However, as shown in the earlier section, these particles are forced to breakdown into still smaller particles under higher shear in injection molding. Experimental observation reflected that the particle size of the dispersed rubber phase in the mixer was in the range of 0.3 - 0.4 µm which is reduced to 40-60 nm after injection molding. In order to have more insight into structure development in the mixer and the injection molding machine, the size of the dispersed rubber particles was predicted theoretically from critical break up law proposed by Wu. 25 It has been reported that when the droplet is roughly spherical, the critical condition for breakup of droplet in a continuous matrix is given by the following equation.5,26,27 D = 4.(η D η m )

0.84

.

.γ 12 γ e .η m

η D ηm > 1

when

(5)

.

where, D is the number average particle diameter, γ e is the effective shear rate, η m is viscosity of the matrix phase, η D is the viscosity of the dispersed phase and γ 12 is the interfacial tension between droplet and matrix. The above equation shows that the particle diameter is an inverse function of shear rate for a particular matrix, when η m ,η D and γ 12 are assumed to be constant. Shear rate in the mixer and the injection molding machine is not constant because of complex shear deformation, flow and temperature profiles along the various sections. Hence, the viscosity of rubber and plastic were determined in the mixer and the injection molding machine at various shear rates- 40 – 100 s-1 (mixer) and 500 – 800 s-1 (injection molding) at their respective mixing temperatures 190 °C and 220 °C respectively (see supporting 14 ACS Paragon Plus Environment

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information). The shear rate range in the mixer and the injection molding has been supplied by the manufacturer (Thermo Fisher Scientific). Interfacial tension (α) between the PP matrix and FKM droplet was found to be 0.5 × 10-3 N/m, calculated from Palierne model 24, 28 which is shown as follows

G * = Gm*

1 + 3ΦH (ω ) 1 − 2ΦH (ω )

(6)

H (ω ) is given by the following expression

H (ω ) =

( (

) ( ) (

)( )(

) )

4α R 2Gm* + 5Gd* + Gd* − Gm* 16Gm* + 19Gd* 40α R Gm* + Gd* + 2Gd* + 3Gm* 16Gm* + 19Gd*

(7)

where G * , Gm* and Gd* are the complex modulus of the blend, matrix phase, and dispersed phase respectively. Φ and R are the volume fraction and the average radius of the dispersed phase respectively. R was measured accurately from FESEM, AFM and TEM images. The experimental

values were substituted in equation (7). Interfacial tension, α value was

obtained by using equations (6) and (7). The values of ηd and ηm of 40PP/60FKM (w/w) TPE at different shear rates and temperatures are reported in Table 1. It is demonstrated that the diameter of the particles is larger when the shear rate is lower. In addition, when the shear rate is high, as in the case of injection molding, the size of the particles is reduced. Theoretical calculation predicts that the average value of the particle is 285 ± 50 nm (after mixing in the mixer) which is reduced to 88 ± 5 nm after injection molding. Therefore, the experimental values of the dispersed domain size are in line with the theoretical prediction.



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Table 1. Viscosity and number average particle diameter

Hakke rheomix

Injection molding

ηd

ηm

(Pa.s)

(Pa.s)

40

2250

500

353

60

1900

420

80

1630

100

1450

.

γ (s-1) at 190 °C

ηd

ηm

(Pa.s)

(Pa.s)

500

470

130

91

282

600

400

110

89

350

260

700

350

95

88

300

250

800

300

85

84

D (nm)

.

γ (s-1) at 220 °C

D (nm)

3.4. Theoretical interpretation of Complex Modulus. In order to predict the morphology of the blends, Kerner’s two phase theoretical model was applied 20 as follows: (1 − φi ) Em* + β (α + φi ) Ei* E* = γ Em* (1 + αφi ) Em* + αβ (1 − φi ) Ei*

(8)

where

α=

(1 + υ m ) 2(4 − υ m ) (1 + υ ) , β= ,γ= (1 + υ i ) (7 − 5υ m ) (1 + υ m )

where E * is the complex modulus of the polymer blend, Em* is the complex modulus of the matrix phase, Ei* is the complex modulus of the dispersed phase. υ , υi and υ m are the Poisson ratio of the blend, dispersed phase and the matrix phases respectively. The Poisson ratio which changes with temperature, was calculated from the following equation 20

{0.17 log E ( glass) − log E (T )} + 0.32 υ (T ) = {log E ( glass) − log E (rubber )} *

*

*

*

(9)

Figure 4 shows the plots of experimental complex moduli (obtained from DMA) of the thermoplastic elastomeric blends at 25 °C and the theoretical values of the complex modulus of the blends (based on Kerner’s model) as a function of volume fraction of the PP phase. It 16 ACS Paragon Plus Environment

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is evident that the experimental complex moduli for the PP/FKM thermoplastic elastomeric blends, especially at higher PP content, are very close to those from Kerner’s rigid matrix– soft filler model, clearly confirming the formation of PP as the continuous matrix and FKM as the dispersed phase. 2.0

1.5

Log E∗ (MPa)

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1.0

0.5

Kerner's rigid matrix-soft filler model Kerner's soft matrix-rigid filler model PP/FKM blends (experimental values)

0.0 0.2

0.4

0.6

0.8

1.0

Volume fraction of polypropylene

Figure 4. Experimental and theoretical complex modulii of PP/FKM blends as a function of volume fraction of PP at 25 °C.

3.5.Tensile Properties. Figure 5 exhibits tensile stress − strain curves of injection molded PP/FKM blends. The values of various important parameters like tensile strength, Young’s modulus, elongation at break (EAB), tension set and hardness were measured (see Table S2, supporting information). It was observed that 40 PP/60 FKM (w/w) blend reflected the best thermoplastic elastomeric properties, such as absence of yield point before failure, higher elongation (145 %) and lower tension set (18 %) as compared to other compositions. Therefore, this particular composition was subjected to dynamic vulcanization. Crosslinking



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of FKM phase during simultaneous mixing with plastic (dynamic vulcanization) leads to a drastic improvement in the tensile strength and modulus. On the other hand, tension set was reduced after dynamic vulcanization in comparison to the unvulcanized counterpart. Variation of rubber-plastic composition affects the properties of the blends. Tensile strength, modulus and hardness increased with increase of PP content in the blend (Table S2, supporting information). It should be noted that when PP content was higher (60-80 wt. %), necking and yielding regions were detected in the stress-strain curve. On the other hand, at higher rubber content (60-80 wt. %), the nature of the stress-strain curves significantly changed from plastic to elastomeric type with the disappearance of necking and yielding regions (Figure 5). In order to correlate the structure of the dispersed rubber domain at different processing stages with the blend properties, tensile test was performed before the injection molding. Important mechanical properties of the representative 40 PP/60 FKM (w/w) blend before injection molding were obtained. Tensile strength, modulus and EAB increased after injection molding when compared to the specimen before injection molding. For example, tensile strength, Young’s modulus and EAB were 15.5 MPa, 100 MPa and 90 % which increased to 19.5 MPa, 140 MPa and 145 % after injection molding. It is important to mention here that when EAB was less than 100 %, as in the case of before injection molding (EAB = 90 %), the composition could not be considered as TPE. The same composition generated TPE only after injection molding when EAB was more than 100 %. Adaptation of injection molding gave finer and smaller dispersion of elastomeric domains in thermoplastic matrix when compared to the phase morphology before injection molding (discussed earlier). This fine structure might help to decrease the total interfacial tension and increase the adhesion/wetting between the rubber and the plastic phases leading to an increase in mechanical properties of the TPEs. In addition, intense flow, plastication and


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freezing properties of injection molding were also responsible for superior mechanical properties of injection molded specimen. 40

Tensile Stress (MPa)

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20 PP/80 FKM (w/w) 40 PP/60 FKM (w/w) 60 PP/40 FKM (w/w) 80 PP/20 FKM (w/w)

30

20

10

0 0

15

30

45

60

75

90 105 120 135 150

Tensile Strain (%)

Figure 5. Tensile stress vs. tensile strain plots of PP/FKM blends. 3.6. Thermogravimetry analysis. Thermal degradation temperatures of PP, FKM and their blends (TPEs and TPVs) were measured using thermogravimetry. Figure 6a-b represents the thermogravimetric curves (TG) and the corresponding derivative curves (DTG) respectively of PP, FKM, 40 PP/60 FKM (w/w) TPE and TPV. Various thermal degradation parameters such as initial degradation temperature (T0), the temperature corresponding to a maximum rate of degradation (Tmax), and weight loss at 450 °C of the pristine polymers and their blends were measured (Table S3, supporting information). Pure FKM and PP showed T0 of 420 °C and 436 °C and the corresponding Tmax of 466 °C and 482 °C respectively. Therefore, FKM has very high thermal stability compared to PP in nitrogen atmosphere which is also evident from the weight loss vs. temperature plot (Figure 6a). Interestingly, Tmax values of PP and FKM in 40 PP/60 FKM w/w TPE shifted towards higher temperatures when compared with the values of pristine PP and FKM. In addition, overall thermal stability of TPE increased over neat PP. The effect of composition ratio on thermal stability of the blends is also given (Table S3 and



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Figure S2a-b, supporting information). Table S3 shows the percent weight loss of the sample remained at 450 °C. A gradual decrease in weight loss of the blends was observed with increasing the FKM content in the blends. On the other hand, PP showed maximum weight loss at 450 °C. This indicates that the thermal stability of the blends increases with increase in FKM content in the blends. Thermal stability of polymer blends mainly depends on the morphology and miscibility and thermal stability of the components. Morphology of the blends showed FKM of higher thermal stability is dispersed in nanometer scale in PP matrix. The thermal degradation results suggested that even though PP/FKM blends are immiscible (two glass transition temperatures, as shown later), the thermal stability of PP can be improved by the incorporation fluorocarbon rubber. From the structural point of view, lower the particle size of the dispersed phase, higher is the contact surface area and wetting between the two polymers. Therefore, due to the nanostructured morphology, more rubber particles are adhered and wetted to the PP matrix, which increase the thermal resistance of the blends. Effect of dynamic vulcanization on thermal stability was also investigated. Dynamically vulcanized blend (TPV) showed two step degradation. Tmax of PP remained unchanged, whereas, Tmax of FKM shifted towards the higher temperature (Figure 6a-b) along with a decrease in weight loss of TPV at 450 °C (Table S3, supporting information). During dynamic vulcanization, a three dimensional rubber network was formed in the PP matrix which delay the thermal degradation of the blend. Decreased volume swell of the thermoplastic vulcanizate compared to the thermoplastic elastomer (discussed later) gave evidence for the formation of network rubber in the PP matrix. Excellent thermal stability of PP/FKM blends motivates us to ask whether we can use FKM as a thermal stabilizer of PP. In order to answer this, one composition of PP (95 PP/5 FKM w/w) was prepared using 5 wt. % FKM as given in the experimental section. Figure 6c-d shows the TG and DTG plots of pristine PP and PP-blend having 5 wt. % FKM. It should be


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noted here that the thermal stability of PP increased in the entire temperature range after incorporation of small amounts of FKM. In addition, T0 and Tmax of PP are shifted towards higher temperatures (shifted by about 5 degrees). Outstanding thermal stability of FKM is responsible to improve the thermal stability of PP. It is well known that the thermal degradation mechanism of PP is governed by the formation of the free radicals followed by chain fragmentation processes which might slow down in the presence of thermally stable nanodimensional fluoroelastomer. From the above study it could be inferred that small amount of FKM could be used as a thermal stabilizer of PP.

Figure 6. (a) Weight (%) vs. temperature of PP, FKM and their blends (TPE and TPV), (b) derivative weight vs. temperature of PP, FKM and their blends (TPE and TPV), (c) weight



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(%) vs. temperature of PP and PP having 5 wt. % FKM, (d)

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derivative weight vs.

temperature of PP and PP having 5 wt. % FKM. 3.7. Dynamic mechanical analysis. Figure 7a-b shows the loss tangent vs. temperature of rubber, plastic and their blends (TPEs and TPVs) over the temperature range of - 50 °C to + 50 °C. In the above temperature range, pristine PP showed a single loss tangent peak at about 19.5 °C corresponding to its Tg, whereas FKM showed Tg at around – 4.0 °C. In the case of all PP/FKM blends (TPEs and TPVs) two separate peaks were observed which correspond to the Tg of the rubber and the plastic phases. The appearance of two separate loss tangent peaks strongly indicates the existence of two separate phases in the PP/FKM blends (i.e. two phase morphological structure) which corroborate the findings from FESEM and AFM studies, as discussed earlier. The loss tangent results of the pristine polymers and their blends (TPEs and TPVs) are reported (Table S4, supporting information) in terms of (i) Tg of the plastic phase, Tgp, (ii) Tg of the rubber phase, Tgr, and (iii) (tan δ)max of the rubber phase. It was found that Tg of the rubber phase and (tan δ)max corresponding to the rubber phase decreased with increasing the plastic content in the blends. On the other hand, Tg of the plastic phase decreased with increasing the rubber content in the blends. The decrease of Tg of the rubber phase in the blends could be explained on the basis of thermal stresses built-up in the FKM particles during melt mixing and molding operation. As a result, a triaxial tension was developed due to lager thermal shrinkage of fluorocarbon rubber compared to PP matrix. As a consequence of this, free volume and chain mobility of the rubber phase increased, resulting in drop of primary Tg. Such type of contradicting features of fluorocarbon rubber during melt mixing with plastic was recently reported in detail by us.20 As PP was mixed with FKM which are


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soft in nature, chain flexibility and mobility of PP increased, leading to decrease in Tg of the plastic phase in the blends. Dynamic vulcanization has a strong influence on the loss tangent and the glass transition values. Due to the restriction of rubber chain mobility imparted by the curing reaction, the Tg of rubber phase increased and height of tan δ decreased after dynamic vulcanization when compared to its unvulcanized counterparts (Figure 7b). For example, the Tg of the rubber phase in TPE was –12.5 °C which increased to –8.5 °C after dynamic vulcanization.

Figure 7. Loss tangent vs. temperature of (a) PP/FKM TPEs, (inset of a; Loss tangent vs. temperature of FKM) and (b) TPE vs.TPV.

3.8. Swelling behaviour in ASTM oil. Oil swelling resistance is one of the important properties of thermoplastic elastomeric materials. Figure 8 shows the swelling behavior of the PP/FKM blends in ASTM oil #3 at 100 °C. It was found that volume swell decreased with increase of fluoroelastomer content in the blends. Volume swell of TPE was further decreased after dynamic vulcanization. For example, volume swell of 40 PP/60 FKM TPE was 10 %, which decreased to 8 % after



dynamic vulcanization. It is clear from the data that PP/FKM blends is excellent in oil resistance which is obviously due to the presence of highly oil resistance fluorocarbon rubber domains in PP matrix. Coran and Patel

8,29

reported volume swell of 109 % and 22 % at 100

°C in ASTM oil #3 for PP/EDPM (40/60 w/w) and PP/NBR (50/50 w/w) TPVs respectively, whereas the PP/FKM (40/60 w/w) TPV swelled only 8 % by volume under similar conditions. The above result shows that PP/FKM TPE is superior in oil swelling resistance and can be used as an oil-resistant thermoplastic elastomers.

18

Volume swell (%)

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15

12

9

6 20

40

60

80

Weight of FKM (%)

Figure 8. Equilibrium volume swelling of PP/FKM blends in ASTM oil #3 at 100 °C versus weight fraction of the FKM in the blends.

3.9. Applications of PP/FKM thermoplastic elastomeric materials. Owing to the unique properties of PP/FKM blends, such as superior flexibility, recyclability, outstanding oil swelling resistance and thermal stability, the materials would find their application in the automotive industry, particularly for car hoses. For the conventional vacuum hose (used for cars) which is two layer structured materials and made of rubbers, manufacturing process is complicated as rubbers are heavy, and a rubber


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kneading process and vulcanizing process are necessary. In recent years, it was reported that an olefin thermoplastic elastomer comprising PP and acrylonitrile-butadiene rubber (NBR) is used in the inner layer, and PP and ethylene-propylene-diene rubber (EPDM) is used as outer layer materials for car hoses.30 Since they have used stabilizer to the olefin thermoplastic elastomers to enhance the heat resistance, it is difficult to reduce the cost and also the stabilizer would disappear with time. On the other hand, PP/NBR and PP/EPDM exhibited higher volume swell in oil. The new material has been developed here in consideration of the above described problems and our objective thereof was to develop a material which would be superior in flexibility and outstanding in heat resistance and oil resistance without the incorporation of any additive. The results of mechanical properties showed that the olefin thermoplastic elastomers here (PP/FKM TPEs) have good elongation (~145 %) indicating the superior flexibility of this material. In addition, this thermoplastic elastomer can be recyclable. Therefore, it can directly be inserted, and attached and also can be recycled. TGA results confirmed that PP/FKM blend has excellent thermal stability without incorporation of any thermal stabilizer. Hence, there is no need of additional thermal stabilizers. We have also shown that PP/ FKM TPE has higher overall thermal stability in comparison to commercially important PP/NBR and PP/EPDM TPEs, having the same weight percent of the rubber and the plastic (Figure S3, supporting information) . For example, 40PP/60 FKM (w/w) TPE showed a 17 % weight loss at 450 °C, whereas 40PP/60 EDPM (w/w) and 40PP/60 NBR (w/w) TPEs exhibited 20 % and 30 % weight loss respectively at the same temperature. Swelling test revealed that PP/FKM TPEs exhibited outstanding oil swelling resistance. The commercial olefin thermoplastic elastomers consisting of PP and EPDM, are inferior in oil resistance and they can only be used as cover material where oil does not directly contact . This cannot be used as inner layer (liner) material . However, olefin thermoplastic elastomers consisting of PP and FKM, which



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have outstanding oil swelling and heat resistance and superior flexibility, would be suitable as the liner material. 4. CONCLUSIONS In this work, novel nanostructured thermoplastic elastomeric materials consisting of PP and FKM were prepared using micro-injection molding. Morphological analysis using FESEM, AFM and TEM confirmed that the FKM phase was dispersed in the nanometer range (50-70 nm) into the continuous PP matrix. The role of injection molding in adapting the unique nanostructure morphology of thermoplastic elastomeric blends was predicted with the help of visco-elastic drop break up mechanism. It was revealed that high shear rate in injection molding was responsible for the fine nanostructured morphology of the blends. Tensile strength, modulus and EAB increased after injection molding when compared to the values of the specimen before injection molding. For example, tensile strength, Young’s modulus and EAB were 15.5 MPa, 100 MPa and 90 %, which increased to 19.5 MPa, 140 MPa and 145 % after injection molding. TGA analysis showed that the fabricated blends exhibited excellent thermal stability, which increased with FKM content in the blends and was further enhanced after dynamic vulcanization. It was also revealed that 5 wt. % of FKM significantly improved the thermal stability of PP. Therefore, very small amount of FKM could be used as a thermal stabilizer for PP. Swelling results showed that PP/FKM blends exhibited superior oil swelling resistance (swell only 8-10 % by volume in ASTM oil #3 at 100 °C) compared to commercial PP/EDPM of PP/NBR blends. Superior flexibility, outstanding oil resistance and heat resistance characteristic revealed that these thermoplastic elastomers would be suitable as the liner

material for car hoses. This work offers a promising pathway to design high

performance thermoplastic elastomeric materials for various applications.


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“Graphical abstract”

Tailored nanostructured thermoplastic elastomers from polypropylene and fluoroelastomer: Morphology and functional properties

Shib Shankar Banerjeea and Anil K. Bhowmickb*



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“Graphical abstract” Tailored nanostructured thermoplastic elastomers from polypropylene and fluoroelastomer: Morphology and functional properties

Shib Shankar Banerjeea and Anil K. Bhowmickb*


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Tailored Nanostructured Thermoplastic Elastomers from

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