Tcc

CCP

a

ARRA

KBAS

1

ictostmacaacpp

viie

0h

J. of Supercritical Fluids 87 (2014) 83–92

Contents lists available at ScienceDirect

The Journal of Supercritical Fluids

j our na l ho me page: www.elsev ier .com/ locate /supf lu

owards high-performance polypropylene and its randomopolymer: Insight into toughening mechanism of supercriticalarbon dioxide assisted annealing

hengzhen Geng, Guanghui Yang, Hongwei Bai, Yuhan Li, Qiang Fu ∗, Hua Dengollege of Polymer Science and Engineering, State Key Laboratory of Polymer Materials Engineering, Sichuan University, Chengdu 610065,eople’s Republic of China

r t i c l e i n f o

rticle history:eceived 5 October 2013eceived in revised form 2 January 2014ccepted 2 January 2014

a b s t r a c t

A supercritical carbon dioxide (scCO2) assisted annealing technique is adopted to toughen polypropy-lene (PP) and its random copolymer with 3.8 wt% co-monomer content (PPR). By using this technique,PP with 4-time improvement in toughness or PPR with superior toughness are obtained. Meanwhile, thestiffness of the materials could be maintained. To understand the toughening mechanism, scanning elec-

eywords:alanced mechanical propertynnealingtructure–property relations

tron microscope, wide angle X-ray diffraction, calorimetry, dynamic mechanical analysis and small angleX-ray scattering characterizations are carried out. It is demonstrated that the toughening mechanismis strongly related to the microstructure evolutions, including decreased chain density in the mobileamorphous phase, improved crystalline phase and improved rigid amorphous phase after annealing.These microstructure changes could result in higher cavitation ability and diminish inner stress in theamorphous phase, which consequently favors improvement in toughness.

© 2014 Elsevier B.V. All rights reserved.

. Introduction

The toughness of polymers is considered as one of the mostmportant properties to influence their potential industrial appli-ation. However, most polymers are intrinsically brittle at lowemperatures and/or high strain rates, particularly in the presencef flaws and notches. So toughening of polymers has been the mainubject of researches for more than half a century [1–9]. Generally,here are three approaches to improve the toughness of poly-

ers, namely, blending with elastomers [1,2], copolymerization [3]nd blending with inorganic rigid particles [4,5]. Elastomers andopolymerization can improve the toughness of polymers remark-bly, but at the same time, the stiffness and strength of the materialre significantly sacrificed. Blending with inorganic rigid particlesould also toughen polymers in some cases, but very complicatedrocess is usually necessary to achieve desired dispersion status ofarticles and adequate interfacial adhesion.

Recently, annealing at an elevated temperature is found to beery effective in enhancing the toughness of polymers [6–8]. The

mpact strength of polymers (usually blends) could be significantlymproved without sacrificing the stiffness of material [6–8]. How-ver, annealing is usually time-consuming and also required to be

∗ Corresponding author. Tel.: +86 28 85461795; fax: +86 28 85405402.E-mail address: qiangfu@scu.edu.cn (Q. Fu).

896-8446/$ – see front matter © 2014 Elsevier B.V. All rights reserved.ttp://dx.doi.org/10.1016/j.supflu.2014.01.003

carried out at high temperatures (near the melting point of mate-rial), which means low efficiency and risk of degradation [6–8].Besides, its toughening effect for pure polymers is marginal, andeffective toughening is only reported in blends or nanocomposites[6–8]. Therefore, a novel and simple approach is still in demandto prepare high-performance polymer materials with balancedmechanical properties.

On the other hand, supercritical carbon dioxide (scCO2) has beenincreasingly used as a solvent and swelling agent in polymer chem-istry and processing, such as grafting [10], synthesis [11], foaming[12,13] and impregnation of fillers [14]. CO2 is non-poisonous, non-ignitable, relatively cheap, and its critical conditions (Tc = 31.1 ◦C,Pc = 73.8 bar) is relatively easy to be realized [10–14]. The dissolu-tion of polymers in scCO2 is able to increase the chain segmentalmobility [15,16]. It is reported that the plasticization effect of scCO2in amorphous polymers (or in the amorphous portion of poly-mers) can lower of the energy barrier of some solid-solid phasetransitions, change the phase transition process or make thempossible at a lower temperature [15,16]. Since annealing-inducedtoughening should be strongly related to the microstructure rear-rangements during annealing [3,6,17], using scCO2 to assist theannealing process is expected to promote higher chain mobil-

ity of polymers, and thus have an even stronger effect on thefiner structure rearrangements as well as on the material tough-ness. However, such work has yet to be reported to the best ofknowledge.

8 critica

fieatwntcamfsman

2

2

Lmiw(wwmtnCafp

ttPpa

2

ptvcwacsta

(tswc

a

4 C. Geng et al. / J. of Super

Therefore, scCO2 is adopted to assist the annealing processor improving the toughness of polymers in current study. Its expected that scCO2-assisted annealing has better tougheningffect than thermal annealing even for pure polymers. Besides,nnealing can be carried out at relatively low temperatures withhe aid of scCO2 to diminish the possible degradation. In thisay, materials with improved toughness and unaffected rigid-ess may be obtained via a simple annealing process. To validatehis idea, polypropylene (PP) and its random copolymer (PPR) arehosen as the basal resins. ScCO2 assisted annealing and thermalnnealing are adopted to compare their toughening behavior. Theicrostructure changes after both annealing processes are care-

ully studied using various characterizations. Considering the finertructure changes and combining observations of plastic defor-ation behaviors, the structure–property relationship for both

nnealing processes is discussed, and possible toughening mecha-ism is proposed.

. Experimental

.1. Materials and sample preparation

Polypropylene (PP) was a commercial product of Dushanzi Co.td, China (mark name T30S). Its density was 0.91 g/cm3 andelt flow index (MFI) was 2.5 g/10 min (230 ◦C, 2.16 kg) accord-

ng to the producer. Polypropylene random copolymer (PPR)as a commercial available product supplied by Hyosung, Korea

mark name R200P). Its weight average molecular weight (Mw)as 72.2 × 104 g/mol, molecular weight distribution (Mw/Mn)as 4.5, density was 0.91 g/cm3 and percentage of ethylene co-onomer was 3.8 wt% according to the supplier. Pentaerythritol

etrakys 3-(3,5-ditert-butyl-4-hydroxyphenyl) propionate (markame Irganox 1010) was purchased from Ji Xin Commerce Tradeo., Ltd (China) and used as antioxidant. Its purity was ≥98.0%ccording to the producer. Carbon dioxide (CO2) was purchasedrom Air Products Taiyu Company (Chengdu, China) with certifiedurity greater than 99.5 wt%. All reagents were used as received.

The PP and PPR granules with 0.2 wt% antioxidant were respec-ively injection molded into standard dumbbell bars at an injectionemperature of 220 ◦C and injection speed of 7 cm3/s using aS40E5ASE injection machine (Japan). Mold temperature, injectionressure, hold time and hold pressure were set at 30 ◦C, 65 MPa, 20 snd 65 MPa, respectively.

.2. Annealing treatment

The PP and PPR injected samples were placed into a high-ressure vessel, which was preheated to a desired saturationemperature (40 ◦C, 60 ◦C, 80 ◦C, 100 ◦C, respectively). Then theessel was flushed with low pressure CO2 for 5 min, followed byompressing CO2 to 15 MPa with a syringe pump. The specimensere saturated in scCO2 for 3 h. Then the pressure was released at

very low rate of about 0.005 MPa/s to prevent the foaming pro-ess. (As proved by the micrographs in Fig. S1, the scCO2-annealedamples were not foamed at such a low decompression rate.) Afteraken from the vessel, the samples were cooled down in ambientir with the cooling rate of about 20 ◦C/min.

For comparison, samples treated with high pressure (HP) CO215 MPa) at room temperature (23 ◦C), at which CO2 is not withinhe supercritical zone, and samples thermally annealed at corre-ponding temperatures (40 ◦C, 60 ◦C, 80 ◦C, 100 ◦C, respectively)

ithout CO2, were also prepared using the same vessel. Other pro-

esses were kept the same as the scCO2 treated ones.Prior to testing, all the samples were placed in a vacuum oven

t 23 ◦C for 24 h to eliminate the remaining CO2, if any, in the

l Fluids 87 (2014) 83–92

samples. Then the samples were conditioned at 23 ◦C and 50%relative humidity for 48 h.

The as-obtained samples are denominated according to the PPtype and annealing condition. For example, PP-A represents ther-mally annealed polypropylene, PPR-CO2 represents scCO2-assistedannealed polypropylene random copolymer. If not specially noted,the characterized samples were annealed at 100 ◦C.

2.3. Characterizations and measurements

2.3.1. Mechanical testsThe notched Izod impact strength of the specimens was mea-

sured with a VJ-40 Izod machine according to ASTM D256-04.Standard tensile tests were performed using an SANS Univer-sal tensile testing machine, the moving speed of crosshead was50.00 mm/min. The dumbbell-shaped samples for tensile testswere injection-molded according to DIN 53455. The thickness andwidth of the specimens are 4 and 10 mm, respectively. All the mea-surements were carried out at room temperature (23 ◦C) and theaverage value reported was derived from at least five specimens.

Dynamic mechanical analysis (DMA) was performed using aDMA Q800 analyzer (TA instruments, USA). The three-point-bendmode was used, and the heating program was set from −40 ◦C to110 ◦C at a heating rate of 3 ◦C/min. An oscillatory strain of 0.1% andfrequency of 1 Hz were adopted.

2.3.2. Scanning electron microscope (SEM)Scanning electron microscopy (SEM) experiments were per-

formed to observe the impact fractured surfaces, cross-section ofthe impact fractured surfaces and the necking zone of tensile-deformed samples. The cross-section of the impact fracturedsurfaces and the necking zone of tensile-deformed samples wereproduced by cryo-fracture [5,7]. After immersing the sample in liq-uid nitrogen for more than 60 min, the sample was cleaved by awedge in liquid nitrogen. The surfaces were coated with gold andinspected using an FEI Inspect F SEM instrument with an accelera-tion voltage of 20 kV.

2.3.3. Wide-angle X-ray diffraction (WAXD)A Philips X’Pert pro MPD apparatus with Ni-filtered Cu K� radi-

ation (� = 0.154 nm) was adopted to obtain the WAXD spectra inthe diffraction angle (2�) range of 10o–25o at 40 kV and 40 mA.The analysis method of WAXD spectra was depicted by Huo et al.[18]. The overall crystallinity, Xc was calculated according to thefollowing equation:

Xc =∑

Acryst∑Acryst +

∑Aamorp

(1)

where Acryst and Aamorp are the fitted areas of crystal and amorphousregion, respectively.

2.3.4. Differential scanning calorimetry (DSC)Conventional DSC was carried out on a Perkin-Elmer pyris-1 DSC

(USA). For each measurement, 5 mg specimen was sealed in alu-minum pan and heated from 30 ◦C to 200 ◦C with a heating rate of10 ◦C/min.

Temperature-modulated differential scanning calorimetry(TMDSC) was performed with a temperature modulated DSC Q200(TA Instruments, USA). A heating rate of 3 ◦C/min was selectedwith a modulation period of 60 s and temperature amplitude of

0.5 ◦C. About 5 mg specimen sealed in aluminum pan was tested.

Both DSC and TMDSC were calibrated before the measurementswith high-purity indium as a standard. All the measurements wereperformed under dry nitrogen atmosphere.

C. Geng et al. / J. of Supercritical Fluids 87 (2014) 83–92 85

Fig. 1. Effect of annealing on impact strength of PP and PPR. (a) Effect of annealing condition. (b) Effect of annealing temperature. The insets in (a) are pictures of impact-f on dio

2

r0wdct

cm

we

K

wp

ractured bars of the corresponding samples. HP CO2 is short for high pressure carb

.3.5. Small-angle X-ray scattering (SAXS)SAXS measurements were performed using a NanoStar X-

ay diffractometer (Bruker AXS Inc., USA). The wavelength was.154 nm. A piece of specimen sheet with a thickness of about 1 mmas cut from the center of injection molded bar along the flowirection. The two-dimensional SAXS pattern of each sample wasollected with a HI-STAR detector in transmission mode at roomemperature. The sample-to-detector distance was 2300 mm.

The long period (L) could be obtained the from the I(q) − q2

urve, calculated by L = 2�/qmax, where qmax is the position of theaximum intensity on the I(q) − q2 curve.The one-dimensional electron density correlation function K(z)

as obtained from the scattered intensity I(q) by the followingquation proposed by Strobl and Schneider [19,20]:

∫ ∞0

I(q)q2 cos(qz)dq∫

(z) = ∞

0I(q)q2dq

(2)

here z is correlation distance, I(q) is the one dimensional intensityrofile described above.

Fig. 2. Typical stress–strain curves for (a) PPR an

xide, and scCO2 is short for supercritical carbon dioxide.

3. Results and discussion

3.1. Mechanical properties

The impact properties of the samples treated at different condi-tions are shown in Fig. 1.

It is obvious that annealing has significant toughening effecton both PP and PPR while scCO2 assisted annealing has a muchstronger effect than thermal annealing as presented in Fig. 1(a). Forexample, the impact strength of PPR is increased moderately uponthermal annealing from 14.5 kJ/m2 to 29 kJ/m2. By contrast, PPR-CO2 exhibits superior toughness of 56.8 kJ/m2, about 4-fold thatof untreated PPR and 2-fold that of PPR-A. The impact strength ofPP increases only slightly from 5.0 kJ/m2 to 6.9 kJ/m2 after thermalannealing, while it increases about 4 times to 19 kJ/m2 after scCO2-assisted annealing. The impact strength of PP-CO2 is even higher

than that of untreated PPR.

The insets in Fig. 1(a) could give some additional informationon the improved toughness. Stress whitening, which could indicatemicro-void formation and/or plastic flow of the matrix, is far more

d (b) PP treated under different conditions.

86 C. Geng et al. / J. of Supercritical Fluids 87 (2014) 83–92

F R. (b)

w hly inm

dftv

isdtaaao

dsfis

ig. 3. SEM micrographs of the impact fractured surfaces of various samples. (a) PPhite arrows show the crack propagation direction, and the white rectangles rougagnification image.

istinct and much larger in area for impact-fractured PPR-CO2 thanor PPR-A or untreated PPR. This suggests that annealing-inducedoughening mechanism may be related to the formation of micro-oid and/or plastic flow of the matrix.

Fig. 1(b) exhibits the effect of annealing temperature on thempact strength of PPR. As expected, the toughness of the annealedamples rises nonlinearly with annealing temperature. If weefine “the impact strength of toughened PPR divided by that ofhe untreated PPR” as toughening efficiency, then scCO2-assistednnealing shows a higher toughening efficiency than thermalnnealing at the same annealing temperatures. This means scCO2-ssistance can lower the annealing temperature. As a result, the riskf degradation for annealing could be diminished.

Tensile behaviors of PPR and PP treated under different con-itions are presented in Fig. 2(a) and (b), respectively. After

cCO2-assisted annealing, yield strength of PPR is decreased onlyrom 27 MPa to 26 MPa. Considering the strong enhancementn toughness of scCO2-assisted annealing, the slight decrease intrength is acceptable. For PP-CO2, though its strength decreases

PPR-A. (c) PPR-CO2. The numbers 1, 2 and 3 represent different magnifications. Thedicate the location of the higher magnification image in the corresponding lower

slightly compared to its untreated counterpart, it is still muchstronger than untreated PPR. In addition, its impact strength isalso about 30% higher than that of untreated PPR. This means PPwith better balanced mechanical properties than its copolymer canbe obtained simply using scCO2-assisted annealing. For PPR-CO2,its impact strength is within the “super tough” zone (higher than50 kJ/m2) and comparable to that of PP/30 wt% POE blends (about55 kJ/m2 [21]), but its strength is higher than that of the PP/30wt% POE blends (about 22 MPa [21]). So scCO2-assisted annealing isvery effective in preparing polymers with well-balanced mechani-cal properties.

3.2. Fracture morphology

To investigate the toughening mechanism of scCO2-assistedannealing, SEM is adopted to inspect the impact-fractured surfacesof the samples. Typical micrographs of impact-fractured surfacesare shown in Fig. 3.

C. Geng et al. / J. of Supercritical Fluids 87 (2014) 83–92 87

Fig. 4. SEM micrographs of the cross-section of impact fractured surfaces of various samples. (a) PPR. (b) PPR-A. (c) PPR-CO2. The numbers 1, 2 and 3 mean 50 �m, 250 �ma positp

mfcPizC

sttoamtcbCtAhoa

nd 400 �m from the fracture surface, respectively. The white arrows indicate theosition.

Before annealing (Fig. 3(a)), the sample exhibits a brittle fractureode. Two different areas could be noticed on the fractured sur-

ace, which are the rough (crack initiation) zone extending from therack root and the flat (crack propagation) zone at the sample end.lastic deformation (rough surface) could be identified in the cracknitiation zone, while the smooth surface in the crack propagationone suggests rapid fracture of the sample after crack initiation.onsequently, the sample possesses relatively low toughness.

After thermal annealing (Fig. 3(b)), the whole fractured surfacehows plastic deformation (rough surface), and the extent of plas-ic deformation (roughness) increases significantly compared tohat of untreated PPR, consistent with its higher toughness. On thether hand, the fractured surface of PPR-CO2 (Fig. 3(c)) also shows

whole-area plastic deformation, but the extent of plastic defor-ation is more severe. When magnified to the micrometer scale,

he fractured surfaces of both PPR and PPR-A are rather smooth. Inomparison, a lot of striations (sheet-like sheared fragments) coulde noticed perpendicular to the crack propagation direction in PPR-O2. The formation of striations was reported to be strongly relatedo large local deformations and high strain energy dissipation [22].

s a result, PPR-CO2 was not fractured completely and shows theighest toughness (Fig. 1(a)). So it is clear that the different extentsf plastic deformation mainly account for the toughness variationmong the samples under different annealing conditions.

ions of micro-voids. The sketch on the left side of the figure shows the sampling

The different extents of plastic deformation can also be provedthrough the cross-section of broken samples (beneath the crackpropagation zone) in the plane perpendicular to fracture surface, asshown in Fig. 4. Untreated PPR shows no obvious plastic flow under-neath the impact fractured surface (Fig. 4(a)). In contrast, intensiveplastic flow could be noticed underneath the impact fractured sur-faces of PPR-A and PPR-CO2 (Fig. 4(b) and (c)), and the plastic flow inPPR-CO2 is more severe, in good agreement with its highest tough-ness and the observations in Fig. 3. Away from the fractured surface,the severity of the plastic deformation decreases, and micro-voids(with the size of about 100 nm) could be found 350 �m beneath thesurface (Fig. 4(b3) and (c3)), indicated by the white arrows). Theseobservations correlate well with the whitening area in the inset ofFig. 1(a).

It is generally accepted that micro-voids play a crucial role ininitiating plastic flow of PP and thus beneficial to its toughening[2,5,7]. The positions and sequence of micro-voids and plastic flowbeneath the fractured surface in Fig. 4 also suggest this initiationeffect. So it is reasonable to deduce that the variation in toughnessof untreated PPR, PPR-A and PPR-CO2 mainly results from the

difference in cavitation ability and cavitation-related tougheningmechanism (plastic deformation). Among them, the cavitationability seems in the sequence of PPR-CO2, PPR-A and untreatedPPR. This argument is also evidenced by other recent works and

88 C. Geng et al. / J. of Supercritical Fluids 87 (2014) 83–92

stretc

oa

tmpavwsmii

3

o2ttiraw

Ff

Fig. 5. SEM micrographs of the necking zone after

bservations discussing about the toughening mechanism ofnnealing [8,9].

For PP, the micro-voids are not observed underneath the frac-ured surfaces because deformation in PP would not extend so

uch as in PPR. But in the necking zone of tensile deformed sam-les, the difference in cavitation ability among untreated PP, PP-And PP-CO2 could still be evidenced, as shown in Fig. 5. Few micro-oids could be observed in untreated PP, but more micro-voidsould form during deformation after annealing, especially after

cCO2-assisted annealing. Considering the above discussions onicro-voids and the toughening mechanism, it is reasonable to

nfer that the difference in cavitation ability among the sampless also strongly related to their toughness.

.3. Microstructure evolutions

Fig. 6 illustrates the WAXD spectra of the samples and thebtained crystallinity values. In the profiles, the (1 1 0) plane at� = 14.1◦, (0 4 0) plane at 16.9◦, and (1 3 0) plane at 18.6◦ are reflec-ions for �-crystals in PPR, while the (3 0 0) plane at about 16.1◦ ishe reflection for �-crystals. Obviously, the �-crystal crystallinity

s rather low in PPR. The small amount of �-crystals could beetained after thermal annealing, but they completely disappearedfter scCO2-assisted annealing at 100 ◦C. This consists with otherorks focusing on solid–solid phase transitions in the presence of

ig. 6. WAXD profiles of various (PPR) samples. The crystallinity values calculatedrom the curves are given in the profile.

hing for (a) untreated PP, (b) PP-A and (c) PP-CO2.

scCO2 [15]. The total crystallinity increases only slightly after bothannealing treatments, in agreement with other works [8,23]. Theslight change in crystallinity and polymorphic composition couldnot be the reason for annealing-induced cavitation and toughen-ing, because �-crystals of PP has higher cavitation ability and thushigher toughness compared to �-PP [8,24].

Fig. 7 presents the DSC heating curves of various samples.The endotherm signals at higher melting range (higher than theendotherm peak temperature) do not show any difference beforeand after annealing, indicating that the most stable populationsof crystalline lamellas are scarcely affected by annealing [25]. Themost significant change on DSC curves is the appearance of a shoul-der endotherm called annealing peak (indicated by the black arrowsin Fig. 7), which has been well documented in publications con-cerning thermal annealing [8,25,26]. However, the position andmagnitude of annealing peak are rather different between thermaland scCO2-assisted annealing. For thermal annealing, the positionof annealing peak is just a few degrees higher than the correspond-ing annealing temperature, consistent with other works [8,25,26].By contrast, for scCO2-assisted annealing, the onset of annealingpeak is more than 20 degrees above the corresponding annealingtemperature. Besides, the magnitude of annealing peak is higherfor scCO2-assisted annealing than for thermal annealing at the sameannealing temperature. Considering that the annealing peak shouldbe strongly related to the microstructure evolutions during anneal-ing, this phenomenon is reasonable because the plasticizationeffect of scCO2 during annealing should promote the microstruc-ture evolutions to a higher degree.

There are three possible origins for the appearance of annealingpeak [26]. First, the double peak may be attributed to a melt-recrystallization process [27,28]. Second, new crystals may formduring the annealing process, and the melting of new crystals mayaccount for the annealing peak [29,30]. Third, the annealing peakmay be associated with an enthalpy involved structural relaxationof rigid amorphous fraction (RAF). RAF consists of the constrainedamorphous chain segments attached to the crystals across thephase boundaries [8,26,31]. Due to the physical confinement bysurrounding crystals, the chain segment mobility of RAF is signifi-cantly suppressed compared to mobile amorphous fraction (MAF).As a result, its glass transition temperature increases [23,31,32].Since the existence of RAF is evidenced for most semi-crystallizedpolymers (including PP) [23,31,32], the microstructure changes inRAF after annealing should be considered in this case.

TMDSC has been adopted to analyze the origin of anneal-

ing peak [8,29]. In this measurement, with a small sinusoidalmodulation superimposed on the conventional linear heatingprogram, the obtained total heat flow (THF) can be divided into a(capacity-related) reversible heat flow (RHF) and a (kinetic-related)

C. Geng et al. / J. of Supercritical Fluids 87 (2014) 83–92 89

F at diffT

noa

iasitRaaeasfc

ce

Ftf

ig. 7. DSC melting curves of various (PPR) samples. (a) samples thermally annealedhe black arrows indicate the positions of the annealing peaks in various samples.

onreversible heat flow (NHF). Generally, melting appears in notnly RHF but also NHF, while crystallization and enthalpy relax-tion and/or recovery can be found only in NHF [8,29].

A typical result of TMDSC is shown in Fig. 8. The observed anneal-ng peak on the THF seems to reflect only on the NHF, hardlyffecting the RHF. The TMDSC of other samples presented in Fig. S2how similar results. These results suggest that the annealing peaks originated from the relaxation of RAF (the third mechanism). Ifhe first mechanism is valid, there should be an endothermic peak inHF and an exothermic peak in the NHF. As for the second mech-nism, the annealing peak should be found not only in NHF butlso in RHF. But expectations of neither mechanism agree with ourxperimental result. One may argue that the second model couldlso be valid if the newly formed crystals cannot melt in the timecale of the 60-s modulation period. But if a lot of new crystals areormed, long period of the material should decrease, which is not

onsistent with our SAXS result discussed below.

So it is reasonable to infer that the annealing peak is mainlyontributed by the nonreversible enthalpy recovery of RAF, which isxpected to show only in NHF on the TMDSC curves. The appearance

ig. 8. TMDSC scan of PPR-A (thermally annealed at 100 ◦C). The arrows indicatehe positions of the annealing peaks on corresponding curves. THF, NHF, RHF standor total heat flow, nonreversible heat flow and reversible heat flow, respectively.

erent temperatures. (b) samples scCO2-assisted annealed at different temperatures.

of annealing peak means that the volume of RAF which vitrifiesat the corresponding temperature (with corresponding degree oforder) increases due to the microstructure rearrangements duringannealing.

Fig. 9 exhibits plots of the mechanical loss factor (tan ı) as afunction of temperature for PPR treated under different conditions.Two obvious relaxation peaks/maximums can be noticed in theseplots. The peak at lower temperature is related to the �-relaxationresulting from the glass transition in MAF, while the other maxi-mum at higher temperature is related to the �-relaxation, whichmay be contributed by an intra-lamellar block-slip process in thecrystalline phase [33] or segment relaxation in RHF [34]. As can beclearly seen in Fig. 9, the peak position of �-relaxation shifts a littletoward a higher temperature after thermal annealing, and toward amuch higher temperature after scCO2-assisted annealing. There aretwo possibilities for the enhancement in �-relaxation temperature,the first is crystal perfection resulting from annealing, and the sec-ond is the enhancement of RAF in ordering or quantity as evidenced

in TMDSC. In our opinion, both possibilities should account for theenhancement in �-relaxation. On the other hand, the glass tran-sition temperature (�-relaxation temperature) decreases by about2 ◦C after thermal annealing and about 3.5 ◦C after scCO2-assisted

Fig. 9. Mechanical loss factor (tan ı) of various (PPR) samples as a function of tem-perature.

90 C. Geng et al. / J. of Supercritical Fluids 87 (2014) 83–92

F whicp vario

aomda

cs

lbtpRiPrPaTtfcssofs3

aS

TMd

ig. 10. (a) The curve of one-dimensional correlation function for unannealed PPR, inhase (Ltr) are indicated. (b) The curves of one-dimensional correlation function for

nnealing, suggesting that annealing promotes the chain mobilityf MAF to different degrees. Fig. S3 presents the storage and lossodulus of the tested samples. Loss modulus results also show a

ecreased Tg after annealing. The origin of this will be discussedbout later.

The above characterizations (WAXD, DSC and DMA) are alsoonducted on PP and its annealed samples, and the results are veryimilar to those of PPR, as shown in Figs. S4–S6, respectively.

Fig. 10(a) presents the one-dimensional electron density corre-ation function of the SAXS results for untreated PPR. As indicatedy the red arrows, the thickness of various structures at the crys-alline lamella level could be obtained from the curves, namely, longeriod (L), crystalline lamellar thickness (Lc), and the thickness ofAF (Ltr). The thickness of MAF (La) can be acquired by subtract-

ng 2 Ltr and La from L. Fig. 10(b) shows the correlation function ofPR under different treated conditions, and quantitative analysisesults of L, Lc, La and Ltr are listed in Table 1. The long period ofPR is increased from 12.7 nm to 14.3 nm after thermal annealingt 100 ◦C, and to 15.2 nm after scCO2-assisted annealing at 100 ◦C.he long period (L) values obtained from one-dimensional elec-ron density correlation function are very similar to those obtainedrom the I(q) − q2 curve (Fig. S7), ensuring the reliability of theorrelation function results. In particular, Lc and Ltr experience alight increase after thermal annealing and a further increase aftercCO2-assisted annealing, but the rise in L after annealing is mostlyriginated from the increase in La. For instance, La for PPR increasesrom 5.7 nm to 6.9 nm after thermal annealing and to 7.4 nm aftercCO2-assisted annealing. The increment in La is as high as 21% and0%, respectively.

SAXS was also conducted on PP. The changes in L, Lc and La afternnealing in PP are very similar to those in PPR, as presented in Fig.8 and Table 1.

able 1icrostructure characteristics of various samples calculated from the one-

imensional correlation function.

Sample L (nm) Lc (nm) Ltr (nm) La (nm)

PPR-un 12.7 3.6 1.7 5.7PPR-A 14.3 3.8 1.8 6.9PPR-CO2 15.2 4.0 1.9 7.4PP-un 15.3 4.0 1.6 8.1PP-A 16.0 4.1 1.7 8.5PP-CO2 16.4 4.2 1.8 8.6

h long period (L), crystalline lamella thickness (Lc) and thickness of rigid amorphousus (PPR) samples.

Considering the slightly increased crystallinity and Lc, and theimproved quality in RAF with higher order or density (indicated byTMDSC and DMA), a portion of MAF should be rearranged into RAFand the crystalline phase. With further improvement in La, chaindensity in MAF would be significantly decreased. This should bethe cause of the enhanced chain mobility in MAF proved by DMA.It is also noteworthy that the situations in thermal annealing andscCO2-assisted annealing are rather similar, but only to differentextents. The higher extent of microstructure rearrangements forscCO2-assisted annealing may be the result of the enhanced chainmobility by scCO2 during annealing.

3.4. Toughening mechanism

Summarizing the above results and discussions, the annealing-induced improvement in toughness is strongly related to themicrostructure rearrangements for the annealed samples. And themicrostructure rearrangements in this system are mainly lamellathickening/perfection, improved RAF (in ordering or quantity) anddecreased mobility in MAF. There is no difference in the nature oftoughening mechanism or structure evolutions for thermal anneal-ing and scCO2-assisted annealing, the difference lies only in theextent in the structure changes. Then the question is how themicrostructure rearrangements affect the cavitation ability andthus the material toughness.

In our opinion, increased chain mobility in MAF is the resultof significantly decreased chain density. Probably, the much lowerchain density in MAF favors the cavitation process and diminishedinner stress, both beneficial to the toughening mechanism. Basedon this assumption, a sketch of the possible toughening mecha-nism for PP is shown in Fig. 11. After microstructure adjustmentby annealing, MAF of PP becomes looser and RAF is improved inordering and quality. As a result, the larger difference betweenthem makes the mechanical coupling weaker, which is favorablefor initiation of cavitation at RAF/MAF interface before destruc-tion of the crystalline phase [8,35]. More importantly, looser MAFcould also significantly contribute to easier cavitation [8]. As aresult, a large amount of micro-voids form in the local deforma-tion process, which will facilitate shear yielding of the crystalline

phase to a large extent and in a large area range (which showshigh extent of plastic deformation on the fractured-surface), dis-sipating much energy. Besides, the diminished inner stress in themolecular chains is also believed to be an important reason for the

C. Geng et al. / J. of Supercritical Fluids 87 (2014) 83–92 91

F lamelp , the

ihttemoe

piipbltaP

4

aomoma

emasiem

t

ig. 11. A sketch showing the formation of micro-voids after low deformation at the

hase and rigid amorphous phase, respectively. Note that the entanglement density

mproved toughness. After annealing, MAF becomes looser withigher mobility. This means the molecular chains relax from theense state and lie in a low energy state, which will also account forhe higher toughness of the material. This is the proposed tough-ning mechanism of annealing on PP, irrespective of the annealingethod. ScCO2-assisted annealing would only improve the extent

f microstructure rearrangements and thus exhibits higher tough-ning efficiency.

For PPR, the situation seems more complex because rubberyhase will form in the samples, and cavitation should take place

n rubbery phase in this case. But as shown in Fig. S9, the anneal-ng process does not affect the size and morphology of EPR rubberyhase in PPR. So the higher cavitation ability of PPR-CO2 should alsoe strongly related to the looser MAF in the PP phase. Probably, the

ooser MAF would strongly assist the micro-voids to develop dueo its higher mobility. Besides, the elimination of inner stress afternnealing may also play an important role on the toughening ofPR.

. Conclusions

A novel and simple scCO2-assisted annealing technique isdopted to toughen polypropylene and its random copolymer with-ut sacrificing their stiffness and strength. PP with well-balancedechanical properties or PPR with superior toughness could be

btained by using this simple technique. This technique showsuch higher toughening efficiency compared to thermal annealing,

nd can be effective at relative low temperature.Microstructure characterization and toughening mechanism

xploration results reveal that scCO2-assisted annealing and ther-al annealing share the same toughening mechanisms, which

re cavitation-induced plastic deformation and diminished innertress. The higher toughening efficiency of scCO2-assisted anneal-ng compared to thermal annealing at the same temperature is

xplained by the plasticization effect of scCO2, which promotes theicrostructure evolutions to a higher extent during annealing.In view of both the importance of polymer toughening and

he promising application of scCO2 in polymer industry, this[

la level. (a) Untreated PP. (b) Annealed PP. MAF and RAF stand for mobile amorphoussize of micro-voids, macromolecules and lamellas are not to scale.

contribution provides a practical guidance to the industrial appli-cation of scCO2 and production of polymers with well-balancedmechanical properties.

Acknowledgments

This work was supported by the National Natural Science Foun-dation of China (Grant no. 51173112 and Grant no. 51121001) andthe Special Funds for Major State Basic Research Projects of China(2011CB606006).

Appendix A. Supplementary data

Supplementary data associated with this article can befound, in the online version, at http://dx.doi.org/10.1016/j.supflu.2014.01.003.

References

[1] Z. Yu, Y. Ou, G. Hu, Influence of interfacial adhesion on toughening ofpolyethylene-octene elastomer/nylon 6 blends, J. Applied Polymer Science 69(1998) 1711–1718.

[2] O. Okada, H. Keskkula, D.R. Paul, Fracture toughness of nylon 6 blends withmaleated ethylene/propylene rubbers, Polymer 41 (2000) 8061–8074.

[3] G. Theryo, F. Jing, L. Pitet, M. Hillmyer, Tough polylactide graft copolymers,Macromolecules 43 (2010) 7394–7397.

[4] Q. Fu, G. Wang, Polyethylene toughened by rigid inorganic particles, PolymerEngineering and Science 32 (1992) 94–97.

[5] J. Lee, A.F. Yee, Inorganic particle toughening I: micro-mechanical deformationsin the fracture of glass bead filled epoxies, Polymer 42 (2001) 577–588.

[6] H. Wu, X. Li, Y. Wang, J. Wu, T. Huang, Y. Wang, Fracture behaviors of isotac-tic polypropylene/poly(ethylene oxide) blends: effect of annealing, MaterialsScience and Engineering A 528 (2011) 8013–8020.

[7] Y. Lin, H. Chen, C. Chan, J. Wu, High impact toughness polypropylene/CaCO3

nanocomposites and the toughening mechanism, Macromolecules 41 (2008)9204–9213.

[8] H. Bai, F. Luo, T. Zhou, H. Deng, K. Wang, Q. Fu, New insight on the annealinginduced microstructural changes and their roles in the toughening of �-form

polypropylene, Polymer 52 (2011) 2351–2360.

[9] B. Na, R. Lv, Effect of cavitation on the plastic deformation and failure of isotacticpolypropylene, J. Applied Polymer Science 105 (2007) 3274–3279.

10] G. Yang, X. Hu, J. Su, C. Geng, W. Yao, Q. Zhang, Q. Fu, Significant reinforce-ment of poly(propylene carbonate): nanostructured polymer composites of

9 critica

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

2 C. Geng et al. / J. of Super

poly(propylene carbonate)/poly(methyl methacrylate) via a supercritical car-bon dioxide route, J. Supercritical Fluids 82 (2013) 200–205.

11] Z. Zhang, K. Huang, Z. Liu, Synthesis of high molecular weightNylon 46 in supercritical carbon dioxide, Macromolecules 44 (2011)820–825.

12] Z. Xing, M. Wang, G. Du, T. Xiao, W. Liu, D. Qiang, G. Wu, Preparation ofmicrocellular polystyrene/polyethylene alloy foams by supercritical CO2 foam-ing and analysis by X-ray microtomography, J. Supercritical Fluids 82 (2013)50–55.

13] J. Yang, L. Huang, Y. Zhou, F. Chen, M. Zhong, Multiwalled carbon nanotubesgrafted with polyamidoamine (PAMAM) dendrimers and their influence onpolystyrene supercritical carbon dioxide foaming, J. Supercritical Fluids 82(2013) 13–21.

14] E. Akbarinezhad, M. Sabouri, Synthesis of exfoliated conductivepolyaniline–graphite nanocomposites in supercritical CO2, J. SupercriticalFluids 75 (2013) 81–87.

15] L. Li, T. Liu, L. Zhao, W. Yuan, Effect of compressed CO2 on the melting behaviorand ��-recrystallization of �-form in isotactic polypropylene, J. SupercriticalFluids 60 (2011) 137–143.

16] H. Yokoyama, L. Li, C. Dutriez, Y. Iwakura, S. Kenji, M. Hiroyasu, S. Sono, H.Okuda, Horizontally and vertically aligned polymeric nanosheets: CO2-inducedmorphological changes of block copolymer thin films, Macromolecules 41(2008) 8626–8631.

17] D. Ferrer-Balas, M. Maspoch, A. Martinez, E. Ching, R. Li, Y.-W. Mai, Fracturebehaviour of polypropylene films at different temperatures: assessment of theEWF parameters, Polymer 42 (2001) 2665–2674.

18] H. Huo, S. Jiang, L. An, J. Feng, Influence of shear on crystallization behavior of the� phase in isotactic polypropylene with �-nucleating agent, Macromolecules37 (2004) 2478–2483.

19] G.R. Strobl, M.J. Schneider, Direct evaluation of the electron density correlationfunction of partially crystalline polymers, J. Polymer Science: Polymer PhysicsEdition 18 (1980) 1343–1359.

20] G.R. Strobl, M.J. Schneider, Model of partial crystallization and melting derivedfrom small angle X ray scattering and electron microscopic studies on low-density polyethylene, J. Polymer Science: Polymer Physics Edition 18 (1980)1361–1381.

21] J. Cao, Z.Y. Zhao, R.N. Du, Q. Zhang, Q. Fu, Effect of � nucleating agent on mechan-ical properties of PP/POE blends, Plastics, Rubber and Composites 36 (2007)320–325.

22] O.K. Muratoglu, A.S. Argon, R.E. Cohen, Microstructural processes of fracture ofrubber-modified polyamides, Polymer 36 (1995) 4771–4786.

[

l Fluids 87 (2014) 83–92

23] C. Hedesiu, D.E. Demco, R. Kleppinger, G.V. Poel, W. Gijsbers, K. Remerie,Effect of temperature and annealing on the phase composition, molecularmobility, and the thickness of domains in isotactic polypropylene stud-ied by proton solid-state NMR, SAXS, and DSC, Macromolecules 40 (2007)3977–3989.

24] F. Luo, C. Geng, K. Wang, H. Deng, F. Chen, Q. Fu, New understanding in tuningtoughness of �-polypropylene: the role of �-nucleated crystalline morphology,Macromolecules 42 (2009) 9325–9331.

25] R. Hingmann, J. Rieger, M. Kersting, Rheological properties of a partiallymolten polypropylene random copolymer during annealing, Macromolecules28 (1995) 3801–3806.

26] Z. Wei, P. Song, C. Zhou, G. Chen, Y. Chang, J. Li, W. Zhang, J. Liang, Insight into theannealing peak and microstructural changes of poly(l-lactic acid) by annealingat elevated temperatures, Polymer 54 (2013) 3377–3384.

27] M. Bonnet, K. Rogausch, J. Petermann, The endothermic annealing peak ofpoly(phenylene sulphide) and poly(ethylene terephthalate), Colloid and Poly-mer Science 277 (1999) 513–518.

28] A.M. Jonas, T.P. Russell, D.Y. Yoon, Time-resolved SAXS studies of morphologicalchanges in cold crystallized poly (ethylene terephthalate) during annealing andheating, Colloid and Polymer Science 272 (1994) 1344–1351.

29] B.B. Sauer, W.G. Kampert, E. Neal Blanchard, S.A. Threefoot, B.S. Hsiao,Temperature modulated DSC studies of melting and recrystalliza-tion in polymers exhibiting multiple endotherms, Polymer 41 (2000)1099–1108.

30] T. Liu, S.Y.M. Bonnet, I. Lieberwirth, D. Rogausch, J. Petermann, DSC and TEMinvestigations on multiple melting phenomena in isotactic polystyrene, J. Mate-rials Science 35 (2000) 5047–5055.

31] H. Xu, P. Cebe, Heat capacity study of isotactic polystyrene: dual reversiblecrystal melting and relaxation of rigid amorphous fraction, Macromolecules 37(2004) 2797–2806.

32] Q. Zia, D. Mileva, R. Androsch, Rigid amorphous fraction in isotactic polypropy-lene, Macromolecules 41 (2008) 8095–8102.

33] Y. Men, G. Strobl, Evidence for a mechanically active high tempera-ture relaxation process in syndiotactic polypropylene, Polymer 43 (2002)2761–2768.

34] M. Hoyos, P. Tiemblo, J.M. Gómez-Elvira, The role of microstructure, molar

mass and morphology on local relaxations in isotactic polypropylene. The �relaxation, Polymer 48 (2007) 183–194.

35] S. Humbert, O. Lame, J. Chenal, C. Rochas, G. Vigier, New insight on initiationof cavitation in semicrystalline polymers: in-situ SAXS measurements, Macro-molecules 43 (2010) 7212–7221.