ery behavior of samples blend at different temperatures.

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Experimental analysis and prediction of viscoelastic creep properties of PP/EVA/LDH nanocomposites using master curves based on time–temperature superposition

Reza Mahdavi,¹Vahabodin Goodarzi ,²Seyed Hassan Jafari,³Mohammad Reza Saeb,⁴ Shahrokh Shojaei,⁵Hossein Ali Khonakdar^5,6

1Department of Chemical Engineering, Faculty of Engineering, Central Tehran Branch, Islamic Azad University, P.O. Box 19585-466, Tehran, Iran

2Applied Biotechnology Research Center, Baqiyatallah University of Medical Sciences, P.O. Box 19945-546, Tehran, Iran

3School of Chemical Engineering, College of Engineering, University of Tehran, P.O. Box 1115-4563, Tehran, Iran

4Department of Resin and Additives, Institute for Color Science and Technology, P.O. Box 16765-654, Tehran, Iran

5Department of Polymer Processing, Iran Polymer and Petrochemical Institute, P.O. Box 14965-115, Tehran, Iran

6Leibniz-Institute f€ur Polymerforschung Dresdene. V., Hohe Straße 6, D-01069 Dresden, Germany Correspondence to: V. Goodarzi (E - mail: [email protected] or [email protected]) and H. A. Khonakdar (E - mail: [email protected] [email protected])

ABSTRACT:Prediction of viscoelastic behavior of polymers over a long-term period is of vital importance for engineering applications.

An attempt was made to uncover the interplay between the morphology and viscoelastic behavior of compatibilized polypropylene/

ethylene vinyl acetate (EVA) copolymer blends in the presence of layered double hydroxide (LDH) nanoplatelets. The time–temperature superposition (TTS) principle and WLF equations were merged to obtain master curves of storage modulus at defined reference temperatures enabling prediction of storage modulus at high frequency ranges which are not experimentally measureable. Moreover, the creep compliance master curves were acquired for different reference temperatures to predict the creep compliance of nanocomposites over long period of times. It was found that the presence of LDH decreases the creep compliance at long period of times while it decreases the unrecoverable deformation of EVA domains. A simple mechanism was proposed to explain the creep and recovery behavior of samples blend at different temperatures.^V^C 2018 Wiley Periodicals, Inc. J. Appl. Polym. Sci.2018,135, 46725.

KEYWORDS:applications; morphology; structure–property relationships Received 19 March 2018; accepted 16 May 2018

DOI: 10.1002/app.46725

INTRODUCTION

It is well-known that the proportion of dimensions and mechanical properties of polymeric materials change dramati- cally under loading and temperature over a long-term period.^1–3 From this perspective, prediction and characterization of viscoelastic properties of polymers allows deeper understanding of their performance in engineering applications.

The creep and stress relaxation analyses are identified as criteria for assessing time-dependent deformations of polymers. Such characteristics can be controlled by various factors such as degree of crystallinity, flexibility of polymer chain, free volume, and the crosslink density.⁴The free volume of polymer plays a key role in determining their creep behavior.⁵ Diverse attempts have been made to increase the creep resistance of polymers, among which addition of fillers appeared as versatile and

effective approach.^6,7 Fillers play the role of nucleating agent in semicrystalline polymers, but at the same time they are barriers against free motion of polymer chains. Numerous fillers having nanosize dimension in the form of sphere, tube, and platelet have been examined to achieve higher creep resistance. Accord- ing to the literature,⁸the introduction of high aspect ratio fillers with large surface area such as nanotubes has stronger influence on the viscoelastic properties as compared to the fillers with lower aspect ratio i.e., spherical particles. This is mainly due to reinforcing role of high aspect ratio fillers acting as a strong barrier against polymer chains movements.

In the case of spherical nanoparticles, time–temperature superposition (TTS) appeared successful in demonstrating the role of free volume in nonlinear viscoelastic creep behavior of polypropylene (PP) containing two types of silica nanoparticles (high and low surface areas). It has been shown that the silica

V^C 2018 Wiley Periodicals, Inc.

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nanoparticles with high surface area have stronger effect on mechanical properties of PP, while incorporation of PP-g-maleic anhydride (MA) significantly decreases its creep compliance.⁹ The effect of introduction and increasing the weight fraction of BaTiO₃nanoparticles on mechanical and creep behavior of pol- y(vinylidene fluoride) (PVDF) matrix was reported earlier. It was shown that the beta and gamma transitions of PVDF in nanocomposites strongly depend on the applied frequency. Bur- ger’s model presented good fits on the various experimental data. The TTS principle and Williams-Landel-Ferry (WLF) equations were used to obtain master curves at some reference temperatures. According to the generated creep compliance master curves, the creep compliance of the developed nanocomposites was predicted for about 10 years.¹⁰

For nanotube-like fillers, creep and recovery behavior of PP/carbon nanotubes (CNTs) nanocomposites have been explored. It was shown that unrecovered creep strain of nanocomposites loaded with 1 and 2.8 vol % of CNTs was decreased by 53%

and 73%, respectively, compared to that of the pristine polymer.

The Burger’s model and Weibull distribution function were employed to the experimental data to obtain creep and recovery parameters.¹¹

Moreover, the creep behavior of some rubbery thermoplastics materials containing CNTs such as polyurethane (PU)/CNTs¹² and ethylene vinyl acetate (EVA)/CNTs nanocomposites have been examined.¹³ It was found that the presence of 1 wt % of CNTs in the PU/CNTs system strongly reduces its creep strain.¹² For the EVA/CNTs nanocomposites two types of EVA with high and low melt flow index (MFI) were examined and the results indicated that the nanocomposites based on EVA with high MFI value had high resistance against creep strtain.¹³ Compared to neat polymer, the filled system exhibits higher creep resistance due to the presence of filler which acts as a barrier against flow.

Tang et al. studied the effect of various types of carbon fillers such as multi-walled carbon nanotubes (MWCNTs), carbon black (CB) and chemically reduced graphene oxide (CRGO) on creep behavior of polystyrene (PS). The results indicated that the CRGO has prominent effect on reducing creep strain and unrecovered strain of PS.¹⁴

Blending approach is another promising route for increasing the creep stability of polymers. Again a number of factors including interfacial tension between polymer pairs, and blend morphology are found to be effective on creep behavior.¹⁵ Kolarık et al. comprehensively studied creep behavior of polymer blends: polyethylene/cycloolefin copolymer (COC),¹⁶ PP/

COC,⁵ rubber-toughened PP/poly(styrene-co-acrylonitrile),¹⁷ and poly(ethylene terephthalate)/impact modifier.¹⁸ They concluded that the hard second phase can improve the creep resistance while the soft component has an opposite effect on the creep resistance.

Based on literature review on this subject, nano-reinforced blends are expected to reveal a more complex viscoelastic behavior as compared to neat blends. Perezet al. reported a significant improvement of creep resistance of starch/polycaprolac- tone blends upon addition of clay nanoplatelets. They showed

that the incorporation of 7.5 wt % of clay into the blends has significantly improved the creep resistance of the blends.¹⁹ Mata-Padilla and coworkers reported the positive impact of clay on improving the creep resistance of PP/EVA blend nanocomposites. They concluded that the systems with co- continuous morphology exhibit more resistance against creep strain as compared to the systems with droplet-in-matrix morphology.²⁰

Liuet al. investigated the effect of CNTs on tensile creep stability of immiscible poly(L-lactide)/poly(EVA) blend and highlighted creep stability improvement owing to CNTs and proposed a mechanism for the enhancement of creep stability.

They showed that interfacial location of CNTs at the blend interface and formation of percolated CNTs network structure are importance factors in increasing creep resistance of the nanocomposites.²¹

The creep and recovery of polymer blend nanocomposites were frequently addressed by researchers, but very little is known about their time-dependent viscoelastic properties.

The complexity of enhancement of creep in such systems springs from the multiplicity of independent and/or dependent morphological, chemical, and physical parameters con- tributing to viscoelastic response. There are sufficient works focused on different aspects of PP/EVA blends containing organo-modified montmorillonite (OMMT).^22–27 The use of organo-modified layered double hydroxide (LDH) on morphological, thermal, and mechanical properties of PP/EVA blends was also studied in a number of previous investigations of this group.²⁸ Overall, LDHs because of their typical metal-hydroxide-like chemistry and conventional clay-like layered crystalline structure through endothermic decomposition and char formation donate higher thermal stability and flame properties to polymeric matrix and therefore it can be a better candidate for improving creep resistance purpose at high temperatures as compared to OMMT with limited thermal stability.^29–31

In this study, we have performed a detailed viscoelastic creep studies along with multi-frequency Dynamic-Mechanical Ther- mal Analysis (DMTA) investigations on a novel nanocomposite system [75/25 (wt/wt) blends of PP/EVA containing of LDH nanoplatelets] for the first time and obtained the related highly useful and practically important master curves (based on TTS principle and WLF equation) for properties prediction at test conditions not accessible easily. Moreover, analysis of the creep and recovery data based on Burger and Weibull models are reported Therefor, such detailed creep studies on multicomponent nanocomposite systems are very rare. The information presented in this manuscript is highly useful for analyzing the creep and recovery behavior of other complex multicomponent nanocomposite systems.

THEORETICAL BACKGROUNDS Creep and Recovery Model

The Burger’s model has been applied in investigation of creep behavior of viscoelastic materials. The model results from

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combination of Maxwell and Voigt models in series form as shown in Figure

F1 1.

The main Burger’s equation is as follows^10,32 E5rc

Em

1rc

Ek

12exp 2t s

1rc

h_mt (1) where Eis total strain of the Burger model, andEmandhmare the modulus and viscosity of the Maxwell spring and dashpot, respectively, while E_k and h_k are the modulus and viscosity of the Kelvin spring and dashpot, respectively. The retardation time(s) (s) is determined as follows

s5h_k Ek

(2) The Weibull distribution equation is used to determine the recovery strainEr(t) defined as³³

E_rðtÞ5EVI exp 2 t2t0

h_r b_r!

" #

1E_VP (3) where E_VI is the viscoelastic strain recovery that depends on hr

and b_r. Therefore, h_r and b_r are characteristic life parameter and shape parameter, respectively. Also t0 is the time at which the load on samples is removed in tension mode while E_VP is the permanent strain that relates to viscous motion of polymer chains.

TTS Equation

The TTS principle¹⁰ can be used for anticipating viscoelastic properties of polymeric materials at arbitrary reference temperatures. The TTS theorem can be applied in experiments investi- gating time- and frequency-dependency of materials’ behavior to predict their long-term viscoelastic properties. In the TTS theorem, a time shift factor (aT) is introduced as follows:

logaT5log t tT

(4) where t and t_T are actual time and shifted time in the time- dependent experiment, respectively. By using the TTS method, a master curve of transient long-term viscoelastic properties can be obtained from short-term test results of a limited range of temperatures or frequencies. The shift factora_T[defined in eq. (4)] is an important parameter in the TTS method which can be obtained through the well-known WLF equation expressed as¹⁰

logaT52C1ðT2T0Þ

C21ðT2T0Þ (5)

where C₁ and C₂ are WLF constants and T₀ is the reference temperature being within a temperature range of T_g to Tg1100.⁴Commonly,T0is taken asTg, in that case theC1and C₂constants are equal to 17.44 and 51.6, respectively. The WLF constants for other reference temperatures can be obtained from the following equations¹⁰

C₁^T⁰5 C₁^T^gC^T₂^g C₂^T^g1T02Tg

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C₂^T⁰5C₂^T^g1 T02Tg

(7) WhereC₁^T^g andC₂^T^g are the C₁and C₂constants at T_g, respectively.¹⁰ At temperatures higher than the reference temperature (T₀), the curve is shifted to the right with a negative shift factor (logaT<0) while for lower temperatures the curve is shifted to the left with a positive shift factor (loga_T>0).

EXPERIMENTAL Materials

PP (Moplen HP500N) was purchased from LyondellBasell, Rotter- dam, Netherlands; (density50.90 g/cm³, MFI at 2308C and 2.16 kg52.1 g/10 min) was obtained from Basell Company while EVA (Escorene Ultra UL00218CC3; density50.94 g/cm³, MFI at 1908C and 2.16 kg51.7 g/10 min, vinyl acetate content518 wt

%,) was supplied by Nature Works, USA Exxon Mobile Chemical Company. A commercial grade PP grafted MA PP-g-MA, Polybond 3200, consisting of 1 wt % maleic anhydride was obtained from Chemtura Inc. Magnesium nitrate and aluminum nitrate and sodium dodecyl benzene sulfonate used for synthesis of organo- modified Mg–Al LDH by one-step route were obtained from Merck. The organo-modified LDH was synthesized based on a pro- cedure reported elsewhere.²⁷Based on our thermogravimetric analysis studies (not reported here), the organic content of the organo- modified LDH is about 10 wt %. Prior to melt blending, all the materials were dried for 24 h at 508C in a vacuum oven.

Sample Preparation

The samples were prepared in a micro-compounder (DACA) under the following processing conditions: rotation speed of 150 rpm, mixing time of 5 min, and temperature of 2108C. PP and EVA components were compounded in a single-step melt mixing process with different amounts of PP-g-MA and LDH as detailed in Table I. All the samples were prepared based on the T1 weight percent in which the PP/EVA ratio was kept constant at 3/

1 in all the blend and nanocomposite systems. For nanocomposite samples, a known fraction of the total blend, with same ratio, was Figure 1.Burger model for representation of viscoelastic behavior. [Color

figure can be viewed at wileyonlinelibrary.com]

COLORONLINEANDBWINPRINT

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replaced with the required amount of LDH. The samples for DMTA measurements were prepared by a compression molding process. The molding temperature of 2008C and pressure of 10 bars were applied. The samples were allowed to cool down to the ambient temperature at a cooling rate of 10 K/min.

Characterizations

Multi frequencies dynamic mechanical analysis (DMA, 0.1, 1, 5, 10, and 20 Hz) was done by means of dynamic mechanical

analyzer (DMA2980, TA Instrument, USA). This analysis was performed in a tension mode on the molded samples with dimensions of 20 mm 3 10 mm 3 0.5 mm. Heating rate of 3 K/min and temperature range of280 to 1508C were selected.

The tensile creep test was performed by means of DMA in a tension mode. Also, the creep and recovery cycles were con- ducted at isotherms of 0, 15, 30, 50, 70, and 908C. An equilib- rium time of 5 min was allowed for each interval before applying load. For each isotherm, a constant stress of 2.5 MPa Table I.Sample Code and Composition of the Prepared PP/EVA/LDH Nanocomposites

Sample code PP (wt %) EVA (wt %) PP-g-MA (wt %) LDH (wt %)

P100 100 — — —

P75E25 75 25 — —

P75E25C5 75 25 5 —

P75E25L5 75 25 — 5

P75E25C5L1 75 25 5 1

P75E25C5L5 75 25 5 5

Figure 2.Effect of frequency sweeps on storage modulus, loss modulus, and tan d of studied samples: (A–C) P100, (D–F) P75E25, and (G,H,K) P75E25C5L5. [Color figure can be viewed at wileyonlinelibrary.com]

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was applied for 10 min, followed by a 20 min recovery period.

Considering the conditions used for preparation of samples for the DMTA measurements (applied slow cooling rate of 10 K/

min) and also the relatively fast crystallization rate of PP and EVA, no annealing process were performed on samples before the creep measurements.

RESULTS AND DISCUSSION

Detailed investigations have been performed on morphology of PP/EVA/LDH system using scanning electron microscopy and transmission electron microscopy and the results are reported elsewhere.²⁸ A brief summary of the main morphological find- ings are as follows: the unfilled PP/EVA blend system exhibits a biphasic nature in which EVA domains are dispersed in the PP matrix. Inclusion of LDH nanoplatelets, which localize mainly in the EVA phase of the blend, refines the EVA domain size.²⁸ Further refinement of morphology is obtained when the compatibilizer is added to the filled system.

Frequency Sweep and Master Curves

In our previous work,²⁸we have examined dynamic mechanical properties of PP/EVA/LDH system at single frequency mode. It was found that with the addition of EVA to PP matrix, the storage modulus of PP decreases considerably. Addition of PP-g- MA to this blend due to compatibilization effect increased the storage modulus of the blend to some extent. Moreover, addition of LDH up to 5 wt % to the compatibilized blend further increased the storage modulus of the blend so that the storage modulus of the nanocomposite approaches the storage modulus of the neat PP indicating that PP/EVA/LDH nanocomposite

system exhibit a good balance between stiffness and toughness.

Investigations on damping behavior of PP/EVA/LDP nanocomposite system and its comparison with damping properties of the neat PP and neat PP/EVA blend showed that incorporation of EVA as rubbery phase to PP increased its damping factor while, addition of LDH as hard solid phase to this blend decreased its damping factor. Moreover, the observed three relaxation peaks of g, b, and a transitions were assigned to glass transition temperatures of EVA, PP phases and to release of PP macromolecules from crystal structures, respectively. It is to be noted that based on the DSC data reported previously,²⁸ EVA undergoes melting process in a range of temperature close toatransition (50–1008C) of PP.

Variation of viscoelastic properties of P100, P7525, and P75E25C5L5 samples against different frequencies and temperatures are shown in Figure2. Generally it is seen that as expected F2 an increase in frequency, which reduces the time scale of experiment, increases the elastic characteristics of the materials and Figure 3.The storage modulus versus frequency master curves at different temperatures: (A) P100, (B) P75E25, and (C) P75E25C5L5 along with storage modulus master curves of the same materials at different reference temperatures: (D) 258C and (E) 458C. [Color figure can be viewed at wileyonlinelibrary.com]

Table II.Predicted Storage Modulus of Selected Samples at High Frequen- cies of 100 and 10,000 Hz at Two Different Reference Temperatures

Storage modulus at Tr5258C (MPa)

Storage modulus at Tr5458C (MPa) Samples 100 Hz 10,000 Hz 100 Hz 10,000 Hz P100 1428.7 1709.5 1113.1 1368.9

P75E25 672.3 827.9 502.4 642.4

P75E25C5L5 1063.3 1268.7 808.1 968.0

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shifts all the transitions toward slightly higher temperatures.

The other transitions in the tan d curves above the observed local peak between 0 and 258C can be attributed to the relaxa- tions in the crystalline region of PP and also approaching the melting temperature of the EVA component of the systems.²⁸ In order to predict the materials behavior, such as storage modulus, over long period of times (low frequency) and temperatures, it is necessary to obtain a master curve of storage modulus using the WLF equation at different reference temperatures. For this purpose, the reference temperatures of 25 and 458C were selected. Applying eqs. (6) and (7), the constants of WLF equation were determined. The values of shift factors (a_T) for different temperatures at different reference temperatures for the studied samples were also calculated.

Using the shift factors (a_T) the storage modulus master curves for PP/EVA/LDH system and its corresponding components were concluded and the results are presented in Figure

F3 3. The variation

of storage modulus of P100, P75E25, and P75E25C5L5 samples

against frequency at different temperatures are shown in Figure 3(A–C), respectively. It is seen that increase in frequency and temperature has opposing effects on storage modulus. The storage modulus of the studied samples decreases with rising the temperature, due to enhanced chain motions and partly because of transformation of crystalline zones into amorphous phases.

To evaluate the storage modulus of these materials at higher range of frequencies, the TTS principle was utilized and the master curves at two typical reference temperatures of 25 and 458C were drawn and the results are presented in Figure 3(D,E), respectively. Based on these master curves, the storage modulus of these materials can be predicted at the frequencies not being examined. As a typical example, the storage moduli of P100, P75E25, and P75E25C5L5 at frequency of 1 Hz and reference temperature of 258C are 1193.4, 537.2, and 882.7 MPa, respectively while being 882.7, 347.3, and 587.7 MPa at reference temperature of 458C, respectively. These results indicate that P75E25C5L5 sample shows acceptable properties as compared to the rest of samples. As expected the results Figure 4.Creep compliance–time curves at different applied stresses and temperatures: (A) P100 at 308C and (B) P75E25C5L5 at 908C. [Color figure can be viewed at wileyonlinelibrary.com]

Figure 5.Experimental and theoretical creep behavior of studied samples at different temperatures: (A) 08C, (B) 158C, (C) 308C, (D) 508C, (E) 708C, and (F) 908C. [Color figure can be viewed at wileyonlinelibrary.com]

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indicate that increase in temperature decreases the storage modulus of the materials.

Based on Figure 3, the storage modulus of some selected samples at frequencies not measureable by the DMTA device such as 100 and 10,000 Hz at two reference temperatures of 25 and 458C are reported in Table

T2 II.

Tensile Creep and Recovery Behavior

Creep Compliance Behavior.Determination of linear and nonlinear range of viscoelastic properties in the tensile creep analysis is very important. The range is usually determined by a tensile test. Since in this study the DMTA technique has been used for studying creep behavior of samples in the solid state,

therefore linear and nonlinear range of viscoelastic properties should be determined by the instrument.

Figure 4 shows the creep compliance–time curves for P100 and F4 P75E25C5L5samples at 30 and 908C under various applied stresses from 1.5 to 7 MPa. It is clearly seen that all the applied stresses to the samples safely fall within the linear viscoelastic region.

As a result, atypical tensile stress of 2.5 MPa was chosen to further analyze the creep behavior of samples at different temperatures within a certain linear viscoelastic range. The reason for selecting 2.5 MPa as optimal applied stress is because the applied stress higher than this value falls outside of the linear viscoelastic region.

Table III.Burger’s Model Parameters at Different Conditions

Samples Temp. (8C) hk(GPa s) Ek(GPa) hm (GPa s) Em(MPa) s(s)

P100 0 0.26 1.02 31.5 1.21 0.25

P75E25 0.19 0.91 31.4 0.78 0.02

P75E25C5 0.17 0.94 30.0 1.11 0.19

P75E25L5 0.26 1.12 40.0 33.0 0.24

P75E25C5L1 0.15 0.35 33.0 5.23 0.42

P75E25C5L5 0.15 0.58 18.1 4.00 0.26

P100 15 0.14 1.01 20.8 1.37 0.14

P75E25 2.21 0.87 23.6 1.01 2.53

P75E25C5 3.32 1.15 11.0 0.64 3.09

P75E25L5 0.12 0.72 17.1 7.31 0.17

P75E25C5L1 0.04 0.33 21.0 1.00 0.11

P75E25C5L5 0.07 0.44 7.12 3.00 0.15

P100 30 2.84 0.85 10.9 1.83 3.35

P75E25 0.09 2.78 5.81 3.71 0.32

P75E25C5 0.15 0.32 5.94 1.13 0.45

P75E25L5 0.21 0.39 5.22 0.93 0.51

P75E25C5L1 0.27 0.43 4.81 0.67 0.62

P75E25C5L5 0.33 0.45 4.84 0.81 0.66

P100 50 0.77 0.26 5.08 0.28 2.97

P75E25 0.06 0.17 2.91 0.25 0.36

P75E25C5 0.17 0.22 2.42 0.60 0.77

P75E25L5 0.22 0.22 3.31 0.32 0.89

P75E25C5L1 0.26 0.43 4.52 0.62 0.62

P75E25C5L5 0.29 0.45 3.83 0.75 0.65

P100 70 0.35 0.13 4.33 0.16 2.68

P75E25 0.04 0.08 1.00 0.18 0.82

P75E25C5 0.03 0.06 1.00 0.67 0.52

P75E25L5 0.14 0.09 1.00 0.58 1.05

P75E25C5L1 0.07 0.12 1.62 0.18 0.61

P75E25C5L5 0.06 0.09 1.23 0.48 0.71

P100 90 0.34 0.12 4.22 0.14 2.63

P75E25 0.03 0.06 0.91 0.48 0.48

P75E25C5 0.02 0.06 0.97 0.39 0.43

P75E25L5 0.08 0.09 0.91 0.32 0.84

P75E25C5L1 0.06 0.12 1.42 0.19 0.57

P75E25C5L5 0.32 0.45 5.00 0.81 0.67

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Figure

F5 5 demonstrates the creep compliance of the studied samples at various temperatures of 0, 15, 30, 50, 70, and 908C.

According to application of PP/EVA based materials in the automotive industry which are normally exposed to wide range of temperatures, it seems that properties investigations of these materials at this particular range of temperature are highly important. Specifically as mentioned before, the PP under goes a transition and EVA melts in this range of temperature, and therefore, these phenomena exert major impact on the viscoelastic properties of the materials. It is clear that the creep compliance of the samples increases with increasing the temperature. From Figure 5(A) it is seen that the creep compliance of P75E25L5 is lower than that of the rest of samples. This can be attributed to the presence of LDH layers and their impeding effect on the chain mobility.^24,25 On the other hand, addition of EVA to PP matrix decreases amount of crystallinity (increases the chain mobility), and hence, enhances the creep compliance. Figure 5(B) shows the creep compliance of samples at 158C. At this temperature, P75E25L5 sample exhibits the lowest creep compliance among all samples while, P75E25C5L5 and P75E25C5L1 samples show higher creep compliance in comparison to the neat PP. This has its root in compatibilizer and its influential role in compatibilizing the system and decreasing crystalline zones of PP phase and/or decreasing the size of soft EVA dispersed domains within PP matrix since addition of 5 wt % compatibilizer to PP/EVA blend (P75E25C5) increases the slope of the creep compliance versus time curves

more rapidly. On the other hand, the presence of LDH in the PP/EVA blend decreases its creep compliance.

The creep behavior of samples at 308C is shown in Figure 5(C).

Significant changes are observed in the creep behavior such that the neat PP (P100) exhibits the lowest creep compliance among all samples. Moreover, there is a large difference between the creep compliance of neat PP with those of other samples which contain EVA as the second phase. It seems that at 308C the mobility of EVA chains being at rubber state is high enough so that even the presence of LDH cannot reduce its segmental mobility indicating that the temperature has significant impact on increasing the entropy of the system and increasing the creep compliance.

The creep behavior of samples at 508C is shown in Figure 5(D).

Unlike the previous condition there is more similarity between the creep behaviors of all investigated samples. Again P75E25 has the highest creep compliance among all samples. The creep compliance of P75E25C5L5 approaches toward compliance of the neat PP implying that the creep compliance is under the control of LDH at 508C which is close to rubbery flow region of EVA and close to theatransition of PP.

The creep behavior of samples at 708C, which is approaching to the melting temperature of EVA, is shown in Figure 5(E). Sig- nificant rise in creep compliance and slope of the curves in blends and nanocomposite samples is observed due to large movement of the chains as EVA approaches its melting Figure 6.Creep compliances and master curves at 25, 30, 45, and 708C as reference temperatures: (A–C) P100 and (D–F) P75E25, and (G–I) P75E25C5L5. [Color figure can be viewed at wileyonlinelibrary.com]

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temperature and a transition of PP phase. The neat PP again exhibits the lowest creep compliance while the neat blend samples show the highest compliance and the LDH-filled blends placed in between these two limits. At 908C, EVA undergoes viscous flow and again similar behavior to that of 708C is obtained with slight increase in compliance [Figure 5(F)].

The creep compliance of PP was studied by other researchers such a Razavi-Nouri³⁴and Kolarık and Pegoretti.³⁵The value of creep compliance of the PP used in the present work is different from the ones used by the other researchers. The PP used by Razavi-Nouri is an Ethylene propylene rubber (EPR) copolymer containing 5 wt % of ethylene which exhibits low compliance Table IV.Anticipated Creep Compliance of the Samples at Different Reference Temperatures Obtained from Master Curves

Time (min)

Creep compliances at 258C (MPa²¹)

Creep Compliances at 708C (MPa²¹)

P100 10 2.34e²⁷ 2.61e²⁷ 3.79e²⁷ 7.07e²⁷

10,000 4.02e²⁷ 5.02e²⁷ 8.93e²⁷ 1.68e²⁶

10,000,000 1.10e²⁶ 1.47e²⁶ — —

P75E25 10 2.97e²⁷ 3.29e²⁷ 6.79e²⁷ 1.05e²⁶

10,000 8.07e²⁷ 9.98e²⁷ 1.68e²⁶ —

10,000,000 2.08e²⁶ 2.61e²⁶ — —

P75E25C5L5 10 2.68e²⁷ 2.83e²⁷ 4.19e²⁷ 7.98e²⁷

10,000 4.59e²⁷ 5.02e²⁷ 1.35e²⁶ —

10,000,000 1.88e²⁶ 1.87e²⁶ — —

Figure 7.Experimental creep recovery and theoretical recovery based on Weibull distribution function of studied samples at different temperatures: (A) 08C, (B) 308C, and (C) 708C. [Color figure can be viewed at wileyonlinelibrary.com]

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due to its slightly rubbery nature.³⁴ The various grades of PP used by Kolarık and Pegoretti differ with the one used in the present work.³⁵Moreover, the test conditions used for the creep measurements are different from the ones used here. On the other hand, the creep data reported in this work were obtained by DMTA instrument, while in those works^34,35 the tensile test had been used for the creep measurements. Therefore, it is obvious that the results which obtained on different grades of PP at different test conditions under different testing modes differ with each other.

According to Figure 5, it is seen that the Burger’s model can be well fitted to the experimental creep data. In this regard, the model parameters for each sample at different temperatures were evaluated and the results are reported in Table

T3 III. Using

these data it is possible to well predict the creep behavior of the system under various conditions not measured experimentally.

The influence of temperature on creep compliances of samples along with corresponding creep compliance master curves at 25, 30, 45, and 708C, as reference temperatures, are shown in Fig- ure

F6 6. The effect of temperature in the range of 15–908C on creep compliance of neat PP is shown in Figure 6(A). It is observed that the creep compliance increases with increasing temperature. The analysis was performed within time span of less than 10 min. In Figure 6(B,C), the creep compliance behavior of P100 is predicted at some selected reference temperatures in a broader time span from less than 1 min to 100 billion min.

Figure 6(D) shows the temperature effect on creep compliance behavior of P75E25 blend. A comparison between Figure 6(A,B) reveals that rising of temperature has significant effect on increasing the creep compliances of P75E25 blend. Also creep compliance master curves of P75E25 are shown in Figure 6(E,F). The creep compliance of P75E25C5L5 sample at different temperatures is presented in Figure 6(G). From this figure it can be seen that the addition of 5 wt % of LDH to the blend decreases the creep compliance. The master curves for sample at different references temperatures are presented in Figure 6(H,L).

In order to anticipate the creep compliance of samples at long times in a better manner, from the master curves the creep compliance data of samples at specified times of 10, 1000, and 10,000,000 min were determined and the results are presented in Table

T4 IV.

From data of Table IV, it can be concluded that increasing time and temperature has significant impact on creep compliance.

Also addition of EVA to PP increases the rate of creep compliance, while presence of LDH in the blend has a positive influence on the anticipated creep compliance at long period of time.

Recovery Behavior

The creep recovery behavior of the samples at different temperatures of 0, 30, and 708C is shown in Figure

F7 7.The recovery

behaviors of samples at other temperatures were also analyzed but the results are not reported to shorten the discussion. It is generally observed that the addition of EVA to PP matrix increases its viscous nature and therefore decreases its recovery behavior. On the other hand, the presence of LDH layers increases the elastic nature of the blend and hence increases its recoverability. The added LDH layers can fill up the free space

in between the polymer chains, and hence, decreases the viscous nature of the polymers. P75E25C5L5 sample exhibits an enhanced creep recovery behavior clearly seen in Figure 7(A) where the analysis temperature was maintained at 08C. Intro- duction of compatibilizer into the nanocomposite provides better interaction between all the blend components and therefore facilitates its creep recovery behavior.

The recovery behavior of samples at 308C is shown in Figure 7(B). At such temperature both the polymeric components of the blend are above their Tgs. It is seen that P100 has a very good recovery at this temperature.

Table V.Obtained Weibull Model Parameters for the Studied Samples

Samples

Temp.

(8C)

EVI

(%) hr

(s) br

EVP

(%)

P100 0 0.28 2.18 0.45 0.09

P75E25 0.54 0.27 0.37 0.28

P75E25C5 0.49 0.24 0.30 0.12

P75E25L5 0.15 0.23 0.30 0.40

P75E25C5L1 0.59 0.07 0.22 0.24

P75E25C5L5 0.28 0.45 0.54 0.04

P100 15 0.32 0.02 0.21 0.16

P75E25 0.37 0.42 0.38 0.20

P75E25C5 0.61 0.09 0.39 0.31

P75E25L5 0.24 0.23 0.43 0.10

P75E25C5L1 0.30 0.18 0.37 0.02

P75E25C5L5 0.33 0.35 0.50 0.09

P100 30 0.41 0.01 0.26 0.19

P75E25 1.26 0.06 0.31 0.50

P75E25C5 1.20 0.07 0.37 0.60

P75E25L5 1.20 0.02 0.30 0.91

P75E25C5L1 1.05 0.08 0.34 0.63

P75E25C5L5 1.12 0.14 0.38 0.74

P100 50 1.85 0.94 0.73 1.61

P75E25 2.93 0.05 0.27 1.71

P75E25C5 2.16 0.06 0.33 1.33

P75E25L5 2.42 0.02 0.24 1.82

P75E25C5L1 1.05 0.08 0.34 0.63

P75E25C5L5 1.80 0.03 0.26 1.18

P100 70 3.62 0.13 0.27 1.76

P75E25 5.88 0.23 0.29 2.05

P75E25C5 6.09 0.13 0.28 2.40

P75E25L5 4.20 0.11 0.29 2.76

P75E25C5L1 2.41 1.31 0.41 0.89

P75E25C5L5 4.43 0.23 0.30 1.60

P100 90 3.82 0.11 0.26 2.06

P75E25 7.02 0.14 0.28 2.42

P75E25C5 7.10 0.09 0.26 3.25

P75E25L5 4.93 0.05 0.26 3.36

P75E25C5L1 2.49 1.16 0.38 1.01

P75E25C5L5 5.25 0.14 0.28 2.25

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Also P75E25 exhibits a better recovery as compared to P75E25C5L5. This can be attributed to sensitivity of EVA chains to temperature which leads to increasing chain mobility at high temperature and therefore decreases the stress transfer between EVA and PP macromolecules and as a result enhances the viscous behaviors of the system, and hence, decreases its recovery percent. Figure 7(C) indicates recovery time dependency of studied samples at 708C. It is seen that P75E25L5 sample shows the lowest recovery percent. Since EVA begins to melt around this temperature, the presence of LDH can impose large impact on prevention of chains movement.

On the other hand, the recovery behavior of P75E25C5L1 is better than other samples. This is closely related to dispersion of LDH nanoplatelets within polymer matrix and controlling viscous behavior of EVA phases.

The Weibull model was used to fit on the recovery data and the results are also presented in Figure 7. The values of Weibull model parameters at different temperatures for the studied samples were obtained and the results are summarized in Table

T5 V.

From the obtained data it can be found that temperature and composition of samples are two important parameters which have significant effect on Weibull model parameters.

A mechanism proposed for better understanding and describing the creep and recovery behavior of polymer blend nanocomposites at different temperatures is presented in Figure

F8 8. The

schematic presents creep and recovery behavior of a polymer blend nanocomposites in the presence of EVA, LDH layers, and PP spherulites. Stretching the sample at 08C during creep test results in slight changes in the EVA domains, which recover to their original shape during the recovery stage. Also, PP acts as barrier against chains mobility. At 308C both polymers being above theirTgs have low mechanical properties and initially PP as a matrix is stretched and then the stress is transferred to the EVA phase. Therefore, the EVA spherical domains are changed

to ellipsoidal shape. In the recovery stage, the recovery of EVA deformed domains to their original spherical shape is hard and need high energy therefore the domains prefer to maintain their deformed shape. At temperature of 708C, since EVA starts to melt, elasticity is rapidly decreased. On the other hand, the PP matrix at constant stress undergoes extraordinary large strain.

Such large strain is transferred to EVA phase and changes the shape into cylindrical form. For some of the EVA domains the droplet break up can occur. Presence of LDH layers in the EVA domains can facilitate the droplet break up. This phenomenon can create a serious barrier against recovery process and therefore most of the EVA domains remain in the deformed shape.

CONCLUSIONS

Analysis of frequency and temperature sweeps on the developed nanocomposite systems was performed by DMTA technique and master curves were generated according to TTS principle and WLF equation. From the master curves of storage modulus of samples at 25 and 458C, it was found that P75E25C5L5 sample shows the best mechanical properties at high frequency in comparison with the rest of the samples. The tensile creep analysis of the studied samples at different temperatures was carried out with the help of Burger model and the creep parameters were determined. The creep compliance master curves at different reference temperatures were also obtained and the creep compliance of samples was predicted at a wide range of times i.e., 10, 10,000, and 10,000,000 min. The obtained data showed that presence of LDH layers decreases the creep compliance at long period of times. Recovery analysis of the studied samples showed that the presence of EVA and LDH layers, respectively, increases and decreases the unrecoverable deformation. The viscoelastic parameters of samples in the recovery mode were also determined by Weibull model. A simple mechanism was proposed to explain the creep and recovery behavior of PP/EVA blend reinforced by LDH layers.

Figure 8.Proposed model for creep-recovery behavior of PP/EVA/LDH nanocomposites. [Color figure can be viewed at wileyonlinelibrary.com]

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ACKNOWLEDGMENTS

The authors thank Uta Reuter, DMTA laboratory of IPF Dresden for her technical help.

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