Structure, thermal properties and mechanical properties of sPS-based nanocomposites with and without styrenic elastomer inclusions

Materials Chemistry and Physics 132 (2012) 145–153

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Structure, thermal properties and mechanical properties of sPS-based nanocomposites with and without styrenic elastomer inclusions Fang-Chyou Chiu ∗ , Wen-I Su Department of Chemical and Materials Engineering, Chang Gung University, Tao-Yuan 333, Taiwan, ROC

a r t i c l e

i n f o

Article history: Received 1 February 2011 Received in revised form 7 August 2011 Accepted 11 November 2011 Keywords: Composite materials Nanostructures Polymers Thermal properties

a b s t r a c t Syndiotactic polystyrene (sPS)-based nanocomposites with and without toughener inclusions were successfully prepared. One organo-montmorillonite (20A) and two styrenic elastomers (SBS and SEBS) served as the reinforcing filler and as tougheners, respectively. XRD and TEM results confirmed the achievement of intercalated and partially exfoliated sPS/20A nanocomposites. The presence of SBS or SEBS slightly depressed the dispersibility of 20A. DSC results indicated that 20A inhibited the crystallization of sPS. The presence of SBS or SEBS further retarded the crystallization of sPS; this effect was more apparent with SEBS. The presences of 20A and SBS/SEBS facilitated the formation of ␣-form sPS crystals. The thermal stability enhancement of sPS/20A nanocomposites was confirmed, and was further improved with the inclusion of SBS or SEBS. The stiffness of sPS increased with the sole addition of 20A. The addition of SBS or SEBS greatly increased the impact strength of the composites, especially with the addition of SEBS. The achievement of toughened sPS-based nanocomposites was confirmed. © 2011 Elsevier B.V. All rights reserved.

1. Introduction Syndiotactic polystyrene (sPS), a material possessing superior properties, is a promising engineering thermoplastic [1,2]. Extensive studies have been carried out since its synthesis. Four major crystalline forms (termed ␣, ␤, ␥ and ␦) are identified for sPS, and their occurrence depends on the thermal treatments [3–7]. Forms ␣ and ␤ are the predominant polymorphs under normal crystallization conditions. The less stable ␣-form exhibits a hexagonal structure [3,4], whereas the thermodynamically stable ␤-form is orthorhombic [5]. The monoclinic forms ␥ and ␦ can only be obtained through solvent-induced crystallization [6]. The crystallization kinetics and melting behavior of sPS have also been investigated [8–15]. Cimmino et al. [8] noted that, under the same undercooling, the spherulite growth rate of sPS is more than one order of magnitude faster than that of isotactic polystyrene (iPS). Hong et al. [9] observed that sPS shows three melting endotherms after isothermal crystallization treatments. The lowest-temperature and the middle-temperature melting are attributed to melting of ␤- and ␣-form crystals, respectively; the highest-temperature melting arises from the melting of recrystallized ␤-form crystals. Guerra et al. [10] and Sun and Woo et al. [11,12] systematically investigated the relationships among crystallization kinetics, polymorphic crystals, and the resulting

∗ Corresponding author. Tel.: +886 3 2118800x5297; fax: +886 3 2118668. E-mail address: [email protected] (F.-C. Chiu). 0254-0584/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.matchemphys.2011.11.011

multiple melting peaks in sPS. Basically, a slower crystallization process and a higher pre-melting temperature facilitate the formation of ␤-form crystals. Studies on sPS-based blends have been conducted to enhance the competitiveness of sPS [16–24]. The thermal properties and crystalline morphology of sPS are influenced strongly by the incorporation of miscible counterparts. Guerra et al. [16] concluded that the addition of miscible poly(phenylene oxide) (PPO) but atactic polystyrene (aPS) would be favorable for the formation of ␤-form sPS crystals. Cimmino et al. [17] reported that the crystallization rate of sPS decreases with the addition of PPO. Another study [18] has revealed that decreasing the molecular weight of aPS decreases the ability of sPS to crystallize, as well as the melting temperatures for sPS/aPS blends. Other than the blend systems, sPS-based nanocomposites have recently attracted considerable attention due to their potential in displaying more advanced properties [25–30]. Among the nano-fillers studied previously, montmorillonite (MMT) clay is recognized as an appropriate choice for manufacturing highperformance polymer-based nanocomposites [31,32]. MMT has a hydrophilic character and is frequently modified with certain organic surfactants to improve compatibility with organic polymer matrices (modified MMT is denoted as O-MMT hereafter) for nanocomposites achievement. Wang et al. [25] employed two types of alkyl-imidazoliums as MMT surfactants to fabricate sPS/O-MMT nanocomposites. The nanocomposites exhibited enhanced thermal stability in comparison with neat sPS, and the ␤-form sPS crystal became dominant. Tseng et al. [26] demonstrated a solution-mixing technique that could lead to sPS nanocomposite formation; O-MMT

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played a vital role in facilitating ␤-form crystal formation. Wu et al. [27] reported the crystallization kinetics and polymorphic behavior of sPS/O-MMT nanocomposites. The presence of O-MMT increased heterophase nucleation of the ␣-form sPS crystal. Sorrentino et al. [28] revealed the effect of incorporating O-MMTs on the final properties of injection-molded sPS nanocomposites. The addition of a small amount of O-MMT (1%) was noted to widen the processing window of sPS. As anticipated, sPS/O-MMT nanocomposites exhibit better stiffness and strength in comparison with neat sPS. Nevertheless, the impact strength (toughness) of sPS might decline after the addition of O-MMTs [30]; the versatility of sPS would thus be worsened. From an industrial viewpoint, “low toughness” must be overcome if sPS/O-MMT nanocomposites are to be employed in numerous engineering applications. From an academic viewpoint, the effect of simultaneously incorporating O-MMT and a toughener on the structure and thermal properties of sPS should be disclosed. To enhance the toughness of a nanocomposite, a low-modulus elastomeric component is usually incorporated into the matrix [33,34]. Thus far, little work has been conducted on the toughening of sPS/OMMT nanocomposites. The goal of this work is to examine the combined effects of adding O-MMT and an elastomer on sPS/OMMT nanocomposite preparation and the resulting mechanical properties. The crystallization kinetics, melting behavior, and crystal structure of sPS in the neat state and in the fabricated composites are investigated as well. The issue regarding which sPS crystal form (␣ or ␤) is favored to grow in the presence of O-MMT is illuminated. One commercially available O-MMT and two styrenic elastomers (block copolymers) were used to melt-mix with sPS for composites fabrication. The styrene portion of the two elastomers is expected to show some affinity to the sPS matrix. The influence of the styrenic elastomers’ composition on the phase morphology and physical properties of the sPS/O-MMT nanocomposites are revealed.

2. Experimental Neat sPS resin (XAREC 130ZC) used in this study was manufactured by Idemitsu Kosan Co. Ltd., Japan, having a weight average molecular weight of approximately 2 × 105 g mol−1 . OMMT clay (Cloisite® 20A, denoted as 20A) obtained from Southern Clay Products, Inc. was used as the nano-filler. The organic modifier for 20A is dimethyl dihydrogenated tallow quaternary ammonium ion, and the tallow composition is ca. 65% C18, 30% C16, and 5% C14. Two thermoplastic elastomers were used separately as tougheners for the prepared sPS/20A composites: styrene-butadiene-styrene copolymer (SBS, Kraton D-11155J) (styrene: ca. 40 wt%), purchased from Kraton Polymer Company, and styrene–ethylene–butylene–styrene copolymer (SEBS, Septon 8006) (styrene: ca. 33 wt%), purchased from Kuraray Company. A common melt-mixing procedure with an intermeshing twinscrew extruder (SHJ-20B, L/D = 40) in the co-rotating mode was used to prepare the composites. The screw speed was maintained at 180 rpm. The barrel temperatures were kept at 250–310 ◦ C from the hopper to the die. The ingredients were weighed at predetermined ratios, and then dry-mixed before being fed into the extruder. After mixing, the extruded samples were pelletized, then oven-dried before characterization. For comparison purposes, neat sPS was also melt-extruded under the same conditions. Table 1 lists the formulation and sample designations of the extruded samples. X-ray diffraction (XRD) and transmission electron microscopy (TEM) were employed to evaluate 20A dispersibility in the composites. The crystal forms developed for sPS in the samples were evaluated through XRD as well. A Siemens D5005 X-ray unit with Cu K␣ radiation ( = 0.154 nm) was used to carry out the XRD experiments on thin film specimens at room temperature. The scanning

Table 1 Samples designation and formulations. Designation

Composition

Parts (wt%)

sPS sPS/C3 sPS/C5 sPS/C5/B10 sPS/C5/B20 sPS/C5/E10 sPS/C5/E20

sPS sPS/20A sPS/20A sPS/20A/SBS sPS/20A/SBS sPS/20A/SEBS sPS/20A/SEBS

100 97/3 95/5 85/5/10 75/5/20 85/5/10 75/5/20

rate was set at 0.01◦ s−1 with the X-ray generator operated at 40 kV and 35 mA. The TEM observations were performed on ultrathin cryo-microtomed sections of the composite films (ca. 80 nm) with a JEOL JEM-1230 system using an acceleration voltage of 80 kV. To evaluate phase morphology, scanning electron microscopy (SEM) experiments were performed on cryo-fractured surfaces of the samples using a Hitachi S-3000N system. The samples were immersed in liquid nitrogen for 5 min before being fractured, and then were coated with Au prior to SEM observation. A Perkin Elmer DSC 7 analyzer equipped with an inter-cooler was used for sample crystallization and melting behavior measurements. Nitrogen gas was consistently purged into the DSC during the scans to prevent specimens from thermal degradation at high temperatures. The samples were first melted at 320 ◦ C for 5 min, followed by cooling to 20 ◦ C at different rates for non-isothermal crystallization experiments, or fast-cooling to pre-determined temperatures (Tc s) for isothermal crystallization experiments. The crystallized samples were subsequently heated to 320 ◦ C for melting behavior evaluation. Thermal stability of the samples was characterized using a thermogravimetric analyzer (TGA) on a TA Q50 system under a N2 environment and an air environment at a heating rate of 20 ◦ C min−1 . Dynamic mechanical properties of compression-molded specimens (20 mm × 5 mm × 1.9 mm) were measured using a Perkin Elmer DMA 7e system. The measurements were carried out in the three-point bending mode at a 5 ◦ C min−1 heating rate at a frequency of 1 Hz under ambient atmosphere. Tensile properties (Young’s modulus and elongation at break) of the dumbbellshaped specimens (according to ASTM D638) were determined at a crosshead speed of 10 mm min−1 using a MTS Sintech 5/G system. Flexural moduli of the specimen were determined by the same MTS Sintech 5/G system at a crosshead speed of 2 mm min−1 . Notched Izod impact tests were performed using a CEAST impact tester in accordance with ASTM D256. The tensile/flexural properties and impact strength reported were average values from at least five specimens of the same sample. 3. Results and discussion 3.1. Dispersibility of 20A The dispersibility of 20A within the sPS matrix with and without SBS or SEBS was assessed through XRD experiments (in the range of 2 < 6◦ ), as shown in Fig. 1. The (0 0 1) diffraction of pristine 20A is evident and is located around 2 = 3.7◦ , indicating an interlayer spacing (d-spacing) of 2.39 nm. For the two composites not containing SBS/SEBS (cf. sPS/C3 and sPS/C5), the (0 0 1) diffraction of 20A is much weaker (hardly detectable) and is located around 2 = 2.8◦ . This result suggests that most of the layer-tactoid structure of 20A was intercalated or exfoliated within the sPS matrix, regardless of the 20A loading. When SBS or SEBS was further incorporated into the composites (0 0 1) diffraction of 20A became somewhat discernible. Diffraction remained at a lower angle (2 < 3◦ ) than that of pristine 20A. This observation implies that, although some layered-structure of 20A remains in the sPS/SBS or sPS/SEBS matrix,

F.-C. Chiu, W.-I. Su / Materials Chemistry and Physics 132 (2012) 145–153

2θ = 3.7o

Intensity (a.u.)

20A

sPS/C3 sPS/C5 sPS/C5/B10 sPS/C5/B20 sPS/C5/E10 sPS/C5/E20

2

3

4

5

6

147

To verify the detailed dispersion status of 20A within the composites, TEM experiments were carried out. Fig. 2 displays the TEM micrographs of four composites, with and without elastomer inclusions. In Fig. 2(a) and (b) (c.f., sPS/C3 and sPS/C5), some thicker and some thinner layered-structures of 20A are distributed in the sPS matrix. These observations confirm that intercalated/partially exfoliated nanocomposites had developed. Fig. 2(c)–(f) illustrates the micrographs of the composites with the presence of SBS or SEBS. The SBS and SEBS are represented by the shadowed domains in Fig. 2(c) and (e) of a lower magnification. As exhibited in the figures, the dispersion of 20A was somehow repressed with the presence of SBS or SEBS. The average thickness of 20A layered-structures is thicker than those displayed in Fig. 2(a) and (b), corresponding to the XRD results. Scarce 20A is noted to be present inside the SBS/SEBS phases; that is, 20A was mostly distributed within sPS matrix. These above results indicate that intercalated and partially exfoliated sPS/20A nanocomposites were developed through the twin-screw extruder mixing process. The further incorporation of SBS or SEBS resulted in a slightly larger size of multi-layered 20A structure within the nanocomposites.

2θ ( o) 3.2. Phase morphology Fig. 1. XRD patterns of 20A and the prepared composites.

the intercalation and partial exfoliation of 20A continued to occur upon mixing. The presence of SEBS led to a slightly higher angle 20A diffraction than that of SBS, indicating that SEBS results in a less extent of 20A intercalation and exfoliation.

The SEM micrographs of the cryo-fractured neat sPS and the composites are illustrated in Fig. 3. For clear observation, the SBS and SEBS portions in the composites were etched out with p-xylene. Neat sPS [Fig. 3(a)] and sPS/C5 [Fig. 3(b)] exhibit similar phase morphologies, and no evident 20A aggregates are detectable in sPS/C5.

Fig. 2. TEM micrographs of the composites: (a) sPS/C3; (b) sPS/C5; (c) sPS/C5/B20 (scale bar: 1 ␮m); (d) sPS/C5/B20 (scale bar: 200 nm); (e) sPS/C5/E20 (scale bar: 1 ␮m); (f) sPS/C5/E20 (scale bar: 200 nm).

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Fig. 3. SEM micrographs of (a) sPS; (b) sPS/C5; (c) sPS/C5/B10; (d) sPS/C5/B20; (e) sPS/C5/E10; (f) sPS/C5/E20.

For composites with SBS or SEBS inclusions, a two-phased morphology is evident, with the etched-out SBS/SEBS phase dispersed randomly in the sPS matrix. The SBS mostly formed sphere-like domains, as shown in Fig. 3(c) and (d). The average size of the domains increased from ca. 0.26 to 0.38 ␮m with the increase in SBS content. The micrographs of composites with SEBS [Fig. 3(e) and (f)] depict that, in addition to sphere-like domains, SEBS developed non-sphere-like domains as well. The SEBS domains increased from an average of ca. 0.41 to 0.53 ␮m with increasing SEBS content. SBS shows a smaller domain size than SEBS; this suggests that SBS is more compatible with sPS than SEBS. The higher styrene content in SBS (ca. 40 wt%) compared with SEBS (ca. 33 wt%) might have affected in this observation. 3.3. Crystallization and melting behavior Fig. 4(a) shows the DSC cooling thermograms of the samples at 10 ◦ C min−1 . The crystallization peak temperature (Tp , temperature at the exotherm minimum) of neat sPS is 238.8 ◦ C, whereas Tp shifts to lower temperatures with the addition of 20A exclusively. Thus, 20A did not exhibit a nucleation effect for sPS crystallization; rather, it inhibited sPS crystallization. Higher 20A loading

resulted in a lower Tp . These observations can be ascribed to the formation of confined/constrained environments for sPS crystallization by intercalated/partially exfoliated 20A [35]. The increase of melt viscosity (data not shown for brevity) due to the presence of 20A, which resulted in the decline of sPS chain mobility upon crystallization, should also be taken into consideration. With further incorporation of SBS or SEBS, the Tp of sPS remains lower than that of neat sPS. The crystallization enthalpy (Hc ) of sPS also decreases slightly in the presence of 20A and SBS/SEBS elastomers. The thermograms of the samples cooled at a rate of 40 ◦ C min−1 are depicted in Fig. 4(b). Similar formulation-dependent behaviors were observed, with the exception of Tp , which depressed to a lower temperature for individual samples. The representative Tp s and Hc s of 40 ◦ C min−1 -cooled sPS in the samples are shown in Table 2. The overall isothermal crystallization kinetics of the samples was studied by monitoring the evolution of heat during crystallization at distinct temperatures (Tc s) as a function of time. Fig. 5 illustrates the representative DSC thermograms of neat sPS and sPS/C5 upon isothermal crystallization. The crystallization peak time (tp ) is defined as the time when the crystallization exotherm minimum occurred. A longer tp indicates slower crystallization.

F.-C. Chiu, W.-I. Su / Materials Chemistry and Physics 132 (2012) 145–153

(a)

149

(b)

sPS

sPS sPS/C3

Endothermic →

sPS/C3

Endothermic →

sPS/C5 sPS/C5/B10 sPS/C5/B20 sPS/C5/E10

sPS/C5 sPS/C5/B10 sPS/C5/B20 sPS/C5/E10 sPS/C5/E20

sPS/C5/E20

180

200

220

240

260

280

180

200

220

Temperature (oC)

240

260

280

o

Temperature ( C)

Fig. 4. DSC cooling thermograms of the samples at: (a) 10 ◦ C min−1 cooling rate; (b) 40 ◦ C min−1 cooling rate.

Tc: 252oC

Tc: 252oC

o 250 C o 248 C o 246 C

o 250 C

→

→

o 248 C

o 244 C o 242 C

o 242 C

(a)

tp 0

o 244 C

Endothermic

Endothermic

o 246 C

5

10

15

tp 20

0

(b) 5

Time (min)

10 Time (min)

15

20

Fig. 5. DSC isothermal crystallization curves of (a) sPS; (b) sPS/C5 at distinct Tc s.

Fig. 6 depicts tp−1 versus Tc for neat sPS and its composites. Since tp−1

is proportional to the overall crystallization rate, the results indicate that the crystallization rate decreases with Tc for each sample, as anticipated by the nucleation-controlled crystal growth

theory. Furthermore, at the same Tc , neat sPS crystallizes the fastest (the highest tp−1 value), followed by the composites incorporated with 20A exclusively. The composites with both 20A and elastomer (either SBS or SEBS) inclusions exhibit even slower overall

Table 2 Representative physical data of the samples. Samples

sPS sPS/C3 sPS/C5 sPS/C5/B10 sPS/C5/B20 sPS/C5/E10 sPS/C5/E20 a

Properties Tp (◦ C)/Hc (J g−1 )

P␣ (%)

Td,0.1 (◦ C)a

YM (GPa)

EB (%)

FM (GPa)

IS (J m−1 )

224.0/28.6 222.6/27.5 220.6/27.3 220.0/26.9 218.6/26.2 220.6/26.6 219.3/25.8

44.8 66.6 68.5 89.5 88.4 76.5 77.8

370/338 403/361 404/370 417/377 411/378 405/373 408/371

2.20 2.35 2.36 2.01 1.61 2.03 1.65

2.2 2.4 2.1 2.5 2.7 2.2 6.8

3.13 3.18 3.27 2.86 2.17 2.91 2.23

7.6 8.0 7.5 10.0 17.0 16.0 32.4

Value in N2 /value in air.

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values for the composites are indicative of the nucleation changes to a mixed athermal/thermal type, thereby also reducing the crystal growth dimension. These crystallization kinetics changes can be attributed to the confined/constrained effect from dispersed 20A upon the crystallization of sPS. The kinetic K constants were determined to be larger for neat sPS at individual Tc s, followed by the sole sPS/20A nanocomposites. The incorporation of SBS or SEBS further lowered the K values. Fig. 7(a) and (b) illustrates the DSC melting curves of the samples cooled from the melt at a rate of 10 and 40 ◦ C min−1 , respectively. For each sample pre-cooled at 10 ◦ C min−1 , a curve with double melting peaks (Peaks I and II) is exhibited, whereas a pre-cooling rate of 40 ◦ C min−1 causes one major melting peak (Peak II) with an ex-minor exothermic shallow. These complex melting behaviors are ascribed primarily to the melting–recrystallization–remelting process of sPS crystals (␤-form crystal) during the heating scans [9,12,14]. A faster pre-cooling rate of 40 ◦ C min−1 leads to the development of less stable original crystals, which results in a profound recrystallization process overlapping greatly with the original melting peak (Peak I) during the subsequent heating. Melting of the original crystals (Peak I) is thus hardly observable; the dominant remelting peak (Peak II) occurs. Alternatively, a slower pre-cooling rate of 10 ◦ C min−1 causes the formation of more stable original crystals, which results in a noticeable melting Peak I partially overlapping with a fractional recrystallization process, followed by an evident remelting (Peak II). Fig. 7(a) and (b) also illustrates a worthy noting phenomenon of intensity ratio variation between Peaks II and I for neat sPS and the composites. The estimated intensity ratio of melting Peak II/Peak I for each 10 ◦ C min−1 -cooled sample is included in Fig. 7(a). The ratio evidently is higher for the composites than that of neat sPS. In Fig. 7(b), Peak I is less discernible for the composites than for neat sPS. These features imply that the presence of 20A and SBS/SEBS caused a higher degree of melting–recrystallization–remelting process of sPS crystals (i.e., facilitating the formation of Peak II-associated stable crystals) during the heating scans. Accordingly, as revealed by the above Avrami plot analysis, the crystallization kinetics change due to the formation of nanocomposites should have induced a greater number of less stable crystals development in comparison with neat sPS under the same cooling conditions. The melting behavior of isothermally crystallized samples was also investigated. Fig. 8(a) and (b) respectively depict the representative melting curves of isothermally crystallized neat sPS

1.4

sPS sPS/C3 sPS/C5 sPS/C5/B10 sPS/C5/B20 sPS/C5/E10 sPS/C5/E20

1.2

tp-1 (min-1)

1.0

0.8

0.6

0.4

0.2

0.0 240

242

244

246

248

250

252

254

Tc (oC) Fig. 6. tp−1 versus Tc for neat sPS and its composites.

crystallization rates. Higher loadings of 20A and elastomers led to further depression in the crystallization rate. The Avrami equation [36] was used to analyze the crystallization kinetics of the samples: 1 − Xt = exp(−Kt n )

(1)

this is always cast into ln[−ln(1 − Xt )] = ln K + n ln t

(2)

where Xt is the fractional crystallinity at time t, n is the Avrami exponent, and K is the kinetic constant. The Avrami exponent n can provide qualitative information on the nature of nucleation and crystal growth geometry. The K constant is crystallization raterelated; a larger K value suggests a faster crystallization rate. The values of n, determined from the Avrami plot (i.e., ln[−ln(1 − Xt )] versus ln t) of the samples, range from 4.5 to 5.1 for neat sPS at different Tc s. However, the value declines to the range of from 3.5 to 4.3 for the composites. According to Wunderlich [37], the n values for neat sPS suggest a thermal nucleation process followed by a three-dimensional and branching crystal growth. The lower n

(a)

II

sPS

3.6

sPS/C5

2.9

sPS/C5/B10

3.2

sPS/C5/B20

3.2

sPS/C5/E10

250

I

260

(b) II

1.6

3.6

240

I

I

sPS

Endothermic→

Endothermic→

II/I ratio:

sPS/C5 sPS/C5/B10 sPS/C5/B20

II

II

sPS/C5/E20

sPS/C5/E20

270

280

Temperature (oC)

290

300 240

sPS/C5/E10

250

260

270

280

290

300

Temperature (oC)

Fig. 7. DSC heating thermograms of (a) 10 ◦ C min−1 -cooled samples; (b) 40 ◦ C min−1 -cooled samples.

F.-C. Chiu, W.-I. Su / Materials Chemistry and Physics 132 (2012) 145–153

(a)

(b)

I

151

I

Tc :

Tc : o 252 C

o 252 C

o 250 C

Endothermic →

Endothermic →

o 250 C o 248 C o 246 C

o 248 C o 246 C o 244 C

o 244 C

II

I

240

250

260

270

I

o 242 C

280

290

300

240

250

260

III

o 242 C

II

270

280

290

300

Temperature (oC)

Temperature (oC)

Fig. 8. DSC heating thermograms of isothermally crystallized (a) sPS; (b) sPS/C5/B20 samples.

and sPS/C5/B20. For crystallization at low Tc s, two distinct melting peaks (Peaks I and II) for neat sPS and three melting peaks (Peaks I, II and III) for sPS/C5/B20 are evident. The multiple peaks then evolve to one dominant melting peak as Tc increases for both samples. The low-melting temperature peak (Peak I) shifts to a higher temperature with Tc ; however, the high-melting temperature peak (Peak II) stays nearly constant with Tc . The intensity ratio between Peaks II and I decreases with Tc . These features further confirm that the melting–recrystallization–remelting of sPS crystals arises not only for neat sPS, but also for the composites. Higher Tc results in more perfect sPS crystals in the neat state and in the composites; the probability for recrystallization upon heating thus declines. Consequently, one dominant melting peak (Peak I) located at a higher temperature exists. Furthermore, while comparing the Peak II/Peak I intensity ratio between neat sPS and the composite crystallized at the same Tc , the composite again exhibits a higher value. This observation corresponds to that displayed in Fig. 7 for the non-isothermally crystallized samples,

suggesting a higher amount of less stable ␤-form crystals development for the composite in comparison with that of neat sPS under the same isothermal crystallization temperature. The middle-melting temperature peak (Peak III) observed in sPS/C5/B20 crystallized at lower Tc s (i.e., 242, 244, and 246 ◦ C) is attributed to the melting of ␣-form crystals [9,12,14]. As revealed in the following XRD results, the presence of 20A would facilitate the formation of ␣-form sPS crystals. Fig. 9 illustrates the XRD spectra of the samples crystallized under different conditions. The ␣-form crystal content in the individual specimens could be quantitatively estimated via the following relation [10]: P␣ (%) =

(b)

Intensity (a.u.)

β Intensity (a.u.)

β

α

α

1 2 3

β α

1

α

2 3

4

4 5

5

α β 6

8

10

2θ (o)

12

14

(3)

where 1.8 is the ratio between the intensities of the 2 diffraction peaks located at 11.6◦ and 12.2◦ , respectively, for specimens with the same thickness and crystallinity in the pure ␣- and ␤-forms.

β

(a)

1.8A(11.6)/A(12.2) × 100% 1 + 1.8A(11.6)/A(12.2)

αβ 6

8

10

12

14

2θ (o)

Fig. 9. XRD spectra of (a) air-quenched samples; (b) 40 ◦ C min−1 -cooled samples (1: sPS; 2: sPS/C3; 3: sPS/C5; 4: sPS/C5/B20; 5: sPS/C5/E20).

F.-C. Chiu, W.-I. Su / Materials Chemistry and Physics 132 (2012) 145–153

Meanwhile, A(11.6) and A(12.2) are the areas of the 2 diffraction peaks located at 11.6◦ and 12.2◦ , respectively. Fig. 9(a) shows the results of air-quenched samples. Both characteristic diffraction peaks with a stronger 2 = 12.2◦ peak intensity appear for neat sPS, indicating that the crystals formed are mixtures of ␣- and ␤-forms. The two characteristic diffractions exist for the composites as well, and the intensity of 2 = 11.6◦ peak becomes stronger than that of 2 = 12.2◦ peak. The inclusion of SBS or SEBS in the composites enhances the relative intensity of 2 = 11.6◦ peak. Fig. 9(b) depicts the spectra of 40 ◦ C min−1 melt-crystallized samples. Of the two characteristic diffractions, only the 2 = 12.2◦ diffraction exists for neat sPS. For the composites, a weak 2 = 11.6◦ peak along with the dominant 2 = 12.2◦ peak is observed. The relative intensity of 2 = 11.6◦ peak is again noted to be stronger for the composites with SBS or SEBS inclusion compared with those without SBS/SEBS. The representative ␣-form sPS crystal content in the air-quenched samples owing to Eq. (3) is listed in Table 2. Scarce reports exist regarding the effect of adding an elastomer on the polymorphism of crystalline sPS. The additions of 20A and SBS or SEBS therefore increase the possibility of ␣-form sPS crystal formation. Comparing the results of air-quenched samples with those of 40 ◦ C min−1 melt-crystallized samples confirms that a slower cooling rate forms less ␣-form sPS crystals, regardless of the presence of 20A and/or SBS and SEBS.

(a)

Weight (%)

80

20

in N2 0 300

Fig. 11 depicts the storage modulus (G ) as a function of temperature for the representative samples, as evaluated by DMA. The increase in sPS chain motion with temperature caused the samples to exhibit similar trends of decreasing G . In the temperature range of 80–120 ◦ C, the G drop demonstrates the glass transition (Tg ) of sPS. The sPS/C5 showed higher G values than the other samples, up to 1.13 times those of neat sPS. This G reinforcement is a consequence of the fine dispersion of intercalated/partial exfoliated 20A. The composites including both 20A and either SBS or SEBS had G values lower than those of neat sPS. This behavior suggests that the elastomeric nature of SBS/SEBS played a more important role than 20A in controlling the G of the sPS-based samples. Furthermore, of

350

400

450

500

550

Temperature (oC) (b)

sPS sPS/C3 sPS/C5 sPS/C5/B10 sPS/C5/B20 sPS/C5/E10 sPS/C5/E20

100

Weight (%)

80

60

40

20

in air 0 300

350

400

450

500

550

Temperature (oC) Fig. 10. TGA decomposition curves of the samples (a) in N2 ; (b) in air.

3000 0.30

0.25

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2000

G' (MPa)

3.5. Mechanical properties

60

40

3.4. Thermal stability The TGA-scanned results of the samples under a N2 environment are shown in Fig. 10(a). The inclusion of 20A is noted to increase the thermal stability of sPS evidently. The temperature at 10 wt% (Td,0.1 ) loss increased from 370 ◦ C for neat sPS to 403 ◦ C for the sPS/C3 composite. The sPS/C5 exhibits an even higher Td,0.1 value than that of sPS/C3. The thermal stability of sPS was further enhanced with the simultaneous existence of 20A and SBS or SEBS. The presence of SBS led to higher degradation temperatures compared with SEBS-included composites. These observations can be ascribed to the fine dispersion of 20A and the higher thermal stability of SBS. Fig. 10(b) illustrates the TGA results of the samples under an air environment. Similar formulation-dependent behaviors to those shown in Fig. 10(a) are observed, with the exception of the presence of oxygen in the air, which led to an inferior thermal stability for individual samples. The composites demonstrated higher thermal stability than neat sPS. Of the composites, the SBS-included sPS/C5/B20 exhibited the highest thermal stability. The enhancement in the overall thermal stability of sPS, from the formation of nanocomposites, was more evident under air environment than under N2 environment. The apparent thermal stability enhancement for the nanocomposites was mainly a consequence of the homogeneous dispersion of layered-20A, which resulted in heat and O2 permeability reductions in the sPS matrix during the heating scans. Some representative TGA data are included in Table 2.

sPS sPS/C3 sPS/C5 sPS/C5/B10 sPS/C5/B20 sPS/C5/E10 sPS/C5/E20

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tan δ

152

0.15

0.10

0.05

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0.00 50

100

150

200

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0 0

50

100

150

200

250

Temperature (oC) Fig. 11. DMA storage modulus (G ) as a function of temperature for the samples (inset: tan ı as a function of temperature for the samples).

F.-C. Chiu, W.-I. Su / Materials Chemistry and Physics 132 (2012) 145–153

the two elastomers, due to its higher soft component(s) concentration SEBS caused a higher extent of G decrease than SBS. The loss tangent (tan ı) results of the samples are illustrated in the inset of Fig. 11. The dynamic relaxation peak, corresponding to the glass transition of sPS, is observed for each sample. The peak temperature (i.e., Tg ) was marginally higher (∼3 ◦ C) for the composites than for neat sPS, indicating an effect of dispersed 20A on the chain mobility of sPS around the glass transition region. Some of the sPS chains were confined/constrained within the interlayers of 20A, resulted from the efficient mixing process. A certain interaction occurred between the modifier of 20A and the contacting sPS chains. Further noted is that the sPS/C5 exhibited a lower relaxation peak height (lower hysteresis loss) compared with neat sPS. The SBS- and SEBS-included composites, in contrast, showed higher tan ı peak heights (higher hysteresis loss) than that of neat sPS. The interactions (compatibility) between the styrene blocks of SBS/SEBS and the amorphous portions of sPS should play an important role for this result. The measured Young’s modulus (YM), elongation at break (EB), flexural modulus (FM), and Izod impact strength (IS) are listed in Table 2. The sole addition of 20A increased Young’s modulus and flexural modulus of sPS; increasing 20A loading marginally increased both moduli. The elongation at break of sPS was hardly influenced with the presence of 20A. With a further addition of SBS or SEBS into sPS/C3 and sPS/C5, Young’s modulus/flexural modulus declined to a value lower than neat sPS; however, the elongation at break increased. The impact test results reveal that the impact strength of sPS varied only slightly, within the range 7–8 J m−1 , after adding 20A. Nevertheless, further incorporation of SBS or SEBS resulted in an evident rise in composite impact strength, demonstrating an apparent toughening effect. The improvement was more than 330% for the SEBS-included composite, reaching a value of 32.4 J m−1 . This remarkable improvement is believed to be associated with the interactions between the styrene portion of SBS/SEBS elastomers and the amorphous portion of sPS [24]. A higher soft components’ concentration (ca. 67 wt%) in SEBS played a significant role for its high toughening efficiency. From our viewpoint, if the toughness of the sPS-based nanocomposites required further enhancement, the employment of a reactive compatibilization technology is one of the alternatives [38]. The maleic anhydride grafted SBS/SEBS and maleic anhydride grafted polyolefin elastomers may serve as proper tougheners. 4. Conclusions Different formulated sPS-based nanocomposites were successfully prepared using a twin-screw extruder. The dispersibility of an O-MMT (20A) within sPS matrix with and without SBS or SEBS elastomers was evaluated. The crystalline structure and thermal/mechanical properties of the samples were also investigated. The XRD and TEM results indicate that the layered 20A was mainly intercalated and/or partially exfoliated within the sPS matrix. However, further incorporation of SBS or SEBS slightly depreciated the dispersibility of 20A. The presence of 20A and SBS/SEBS favored the ␣-form sPS crystal development in the composites. The crystallization rate of sPS was inhibited with the addition of 20A, and was further decreased in the presence of SBS/SEBS. Results show

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that 20A and SBS/SEBS induced a higher amount of the less stable ␤-form sPS crystals in comparison with neat sPS under the same crystallization conditions. TGA data confirm the thermal stability enhancement of sPS after adding 20A; the inclusion of SBS or SEBS further increased thermal stability, especially with SBS inclusion. The storage modulus, Young’s modulus, and flexural modulus of sPS increased with the inclusion of 20A. The addition of SBS or SEBS, however, produced a decrease in these properties. Nevertheless, the addition of SBS or SEBS led to an apparent toughening effect on sPS/20A nanocomposites, especially with the addition of SEBS. The successful preparation of toughened sPS-based nanocomposites was concluded. Acknowledgement The authors would like to thank the National Science Council of the Republic of China (Taiwan) for financially supporting this research under contract number NSC-99-2221-E-182-006. References [1] N. Ishihara, T. Seimiya, M. Kuramoro, M. Uoi, Macromolecules 19 (1986) 2464. [2] J. Schellenberg, H.J. Leder, Adv. Polym. Technol. 25 (2006) 141. [3] O. Greis, Y. Xu, T. Asano, J. Petermann, Polymer 30 (1989) 590. [4] Z. Sun, R.L. Miller, Polymer 34 (1993) 1963. [5] C.De. Rosa, G. Rapacciuolo, G. Guerra, V. Petraccone, P. Corradini, Polymer 33 (1992) 1423. [6] Y. Chanati, Y. Shimane, T. Inagaki, T. Ijitsu, T. Yukinari, H. Shikuma, Polymer 34 (1993) 1620. [7] E.B. Gowd, K. Tashiro, Macromolecules 40 (2007) 5366. [8] S. Cimmino, E. Di Pace, E. Martuscelli, C. Silvestre, Polymer 32 (1991) 1080. [9] B.K. Hong, W.H. Jo, S.C. Lee, J. Kim, Polymer 39 (1998) 1793. [10] G. Guerra, V.M. Vitagliano, C.De. Rosa, V. Petraccone, P. Corradini, Macromolecules 23 (1990) 1539. [11] Y.S. Sun, E.M. Woo, Macromolecules 32 (1999) 7836. [12] E.M. Woo, Y.S. Sun, C.P. Yang, Prog. Polym. Sci. 26 (2001) 945. [13] C. Wang, Y.C. Hsu, C.F. Lo, Polymer 42 (2001) 8447. [14] F.C. Chiu, K.Y. Shen, H.Y. Tsai, C.M. Chen, Polym. Eng. Sci. 41 (2001) 881. [15] C.H. Su, U. Jeng, S.H. Chen, C.Y. Cheng, J.J. Lee, Y.H. Lai, W.C. Su, J.C. Tsai, A.C. Su, Macromolecules 42 (2009) 4200. [16] G. Guerra, C.De. Rosa, V.M. Vitagliano, V. Petraccone, P. Corradini, J. Polym. Sci. Polym. Phys. 29 (1991) 265. [17] S. Cimmino, E. Di Pace, E. Martuscelli, C. Silvestre, Polymer 34 (1993) 2799. [18] F.C. Chiu, C.G. Peng, Polymer 43 (2002) 4879. [19] F.C. Chiu, M.T. Li, Polymer 44 (2003) 8013. [20] C. Wang, M.L. Wang, Y.D. Fan, Macromol. Chem. Phys. 206 (2005) 1791. [21] J. Kolarik, L. Fambri, M. Slouf, D. Konecny, J. Appl. Polym. Sci. 96 (2005) 673. [22] B.B. Johnsen, A.J. Kinloch, A.C. Taylor, Polymer 46 (2005) 7352. [23] W. Zhou, M. Lu, K. Mai, Polymer 48 (2007) 3858. [24] F. Picchioni, E. Passaglia, G. Ruggeri, F. Ciardelli, Macromol. Chem. Phys. 202 (2001) 2142. [25] Z.M. Wang, T.C. Chung, J.W. Gilman, E. Manias, J. Polym. Sci. Polym. Phys. 41 (2003) 3173. [26] C.R. Tseng, J.Y. Wu, H.Y. Lee, F.C. Chang, Polymer 42 (2001) 10063. [27] T.M. Wu, S.F. Hsu, J.Y. Wu, J. Polym. Sci. Polym. Phys. 40 (2002) 736. [28] A. Sorrentino, R. Pantani, V. Brucato, Polym. Eng. Sci. 46 (2006) 1768. [29] A.K. Ghosh, E.M. Woo, Polymer 45 (2004) 4749. [30] L. Torre, G. Lelli, J.M. Kenny, J. Appl. Polym. Sci. 100 (2006) 4957. [31] S. Pavlidou, C.D. Papaspyrides, Prog. Polym. Sci. 33 (2008) 1119. [32] H. Deka, N. Karak, Mater. Chem. Phys. 124 (2010) 120. [33] F.C. Chiu, S.W. Fu, W.T. Chuang, H.S. Sheu, Polymer 49 (2008) 1015. [34] Y. Yoo, R.R. Tiwari, Y.T. Yoo, D.R. Paul, Polymer 51 (2010) 4907. [35] F.C. Chiu, T.L. Deng, Mater. Chem. Phys. 125 (2011) 769. [36] M. Avrami, J. Chem. Phys. 7 (1939) 1103. [37] B. Wunderlich, Macromolecular Physics, vol. 2, Academic Press, New York, 1976. [38] W.M. Choi, C.I. Park, O.O. Park, J.G. Lim, J. Appl. Polym. Sci. 85 (2002) 2084.