Thermomechanical history effects on rigid PVC microstructure and impact properties

Thermomechanical History Effects on Rigid PVCMicrostructure and Impact Properties

Louise-Anne Fillot,1,2 Philippe Hajji,1 Catherine Gauthier,1 Karine Masenelli-Varlot1

1INSA de Lyon, GEMPPM UMR CNRS 5510, Bat Blaise Pascal-5eme etage, 7 av. Jean Capelle,69621 Villeurbanne Cedex, France2Additifs Matieres Plastiques, ARKEMA, Centre de Recherche Rhone Alpes, Quai Louis Aulagne, BP 35,69191 Saint-Fons Cedex, France

Received 21 August 2006; accepted 18 October 2006DOI 10.1002/app.25688Published online in Wiley InterScience (www.interscience.wiley.com).

ABSTRACT: Rigid PVC microstructure strongly dependson processing conditions: under both shear and heat influ-ence, gelation process occurs, and the resulting morphol-ogy can be characterized by the so-called gelation level pa-rameter. But thermomechanical history also affects severalother features of the microstructure. In this work, two dif-ferent aspects mainly related to (i) orientation and (ii) mo-lecular mobility are pointed out, and their respective ef-fects on poly(vinyl chloride) (PVC) impact properties aredescribed. Charpy impact tests have been carried out on atypical extruded window profile formulation, evidencing astrong anisotropy effect in those extrudates. It turned outthat modulated differential scanning calorimetry was ableto evidence macromolecular orientation. Then, the effectsof heat treatments on impact properties were investigated.

Below the glassy transition temperature (Tg) of PVC, phys-ical ageing results in a decrease of the impact performance,without affecting PVC anisotropy. The reason of this per-formance loss is the reduction of the molecular mobility,as evidenced by dynamical mechanical analysis experi-ments. After performing a heat treatment above Tg,another important decrease of the impact performance wasobserved. This decrease was attributed to (i) partial disori-entation of the PVC macromolecular chains and (ii) reduc-tion of the molecular mobility. � 2007 Wiley Periodicals, Inc.J Appl Polym Sci 104: 2009–2017, 2007

Key words: polyvinyl chloride; annealing; orientation;impact resistance; modulated differential scanning calo-rimetry

INTRODUCTION

Commonly used in many applications, poly(vinylchloride) (PVC) differs from other thermoplastics bytwo main specificities. The first one concerns its for-mulation, which contains a lot of different additives,allowing PVC to cover a wide range of physical andmechanical behaviors. The second one is the largedependence of its microstructure on processing con-ditions: during processing, PVC microstructure isfragmented into micronic particles, and under theinfluence of shear and heat, the interfaces of thesemicronic domains, known as ‘‘primary particles,’’progressively disappear.1–3 At a smaller scale, thisprocess is combined with a partial melting of thePVC crystallinity.4 The resulting morphology is usu-ally characterized by the so-called gelation level pa-rameter, gelation being fully achieved when primaryparticles structure has completely disappeared. Gela-tion level assessment is very important because gela-

tion state deeply affects both physical and mechani-cal properties of PVC.5–8 Several different gelationassessment methods are well described in the litera-ture,4,9–13 and in a previous work,14 the gelation lev-els of a PVC formulation subjected to different ther-momechanical histories were characterized. It em-erged from this work that gelation process is mainlygoverned by a thermal component that can be accu-rately measured by DSC: this thermal component isrelated to the highest temperature ‘‘seen’’ by PVCcrystallites, and in the following part, it will bereferred to as ‘‘DSC real melt temperature.’’ Thanksto DSC technique, in the case of PVC extrudates hav-ing different mechanical behaviors, it is possible toassess whether or not the gelation state is at the ori-gin of the observed difference. However, to evaluatethe intrinsic performance of additives such as impactmodifiers, one can wonder whether the comparisonat same gelation level is the only requirement. Inother words, do other thermomechanical historyeffects exist that are not taken into account by gela-tion level measurements? Orientation effects andmolecular mobility are pointed out.

Indeed, during a typical PVC extrusion process,shear stresses and drawing loads are expected togenerate a certain orientation level of PVC macromo-

Correspondence to: L.-A. Fillot ([email protected]).

Journal of Applied Polymer Science, Vol. 104, 2009–2017 (2007)VVC 2007 Wiley Periodicals, Inc.

lecular chains. The benefits of molecular orientationon mechanical properties are well known since theeighties, leading for instance to the development oforiented PVC pipes.15,16 To study the effects of suchorientation phenomena, Hitt and Gilbert17,18 devel-oped a specific stretching device, and were able toevidence an increase in the falling weight impactperformance with the draw ratio. However, theeffects of orientation in more conventional processes(extrusion, injection, milling) are not well describedin the literature. For extrusion process, if an orienta-tion of macromolecular chains along the extrusiondirection is generally assumed, neither the impor-tance of this orientation, nor its effect on impactproperties is discussed. Nevertheless, a study of Yar-ahmadi19 revealed an anisotropy in extruded sheetstested in low speed tension.

In addition, mechanical properties of polymers areknown to be dependant on molecular mobility.20 Thismolecular mobility, often associated with ‘‘free vol-ume,’’ influences both viscoelastic behavior and yieldstress. As the impact performance of a polymerdepends on its yielding ability, a significant influenceof the molecular mobility on impact properties isexpected. At room temperature, PVC is in the glassystate, which is known to be an ‘‘out of equilibrium’’state: indeed, in nearly all the cases, cooling rate is toofast to allowmacromolecules rearrangements in a ther-modynamical equilibrium configuration. Neverthe-less, further rearrangements occur upon time, and thecloser to Tg the temperature, the faster the rearrange-ments. This process called ‘‘physical ageing’’ leads to areduction of the molecular mobility. For rigid PVC, theinfluence of physical ageing on the impact perform-ance has been reported in the literature,8,21,22 highlight-ing a decrease of the impact performance after physicalageing heat treatments at 60–708C during a few dozenof hours. Molecular mobility also depends on the cool-ing rate that determines the macromolecules’ glassystate configuration. However, the cooling rate influ-ence on rigid PVC microstructure and impact proper-ties is poorly documented.

During extrusion process, molecular mobility isaffected by the cooling rate implied at the end ofthe process (calibration section), and later, it can bereduced if PVC is stored a long time or used athigh temperature. Sometimes, PVC is also sub-jected to various heat treatments devoted to pre-vent shrinkage or to relax ‘‘internal stresses.’’ Inthis context, it appears fundamental to determineto what extent orientation, heat treatments, andcooling can affect PVC impact properties. This wasthe aim of the current work. Moreover, the signa-tures of the microstructural features involved wereinvestigated so as to propose (at the end) appropri-ate tools capable of explaining the origin of a givenimpact performance.

EXPERIMENTAL

Materials

The PVC formulation used in this study was a stand-ard lead-stabilized window profile recipe detailed inTable I. All ingredients were mixed together up to1108C using a high speed mixing station (Papenmeier75/150). The resulting dryblend was extruded at fivedifferent melt temperatures ranging from 180 to 2008Cusing a conical twin-screw extruder (Krauss MaffeiKMDL 25), the size of the die being 70 � 2.5 mm2 andthe screw speed set at 30 rpm. After extrusion, PVCsheets were cooled through a calibration section regu-lated at 58C by water and air circulation.

The major part of this work was performed onextruded PVC sheets whose formulation is given in Ta-ble I. This PVC formulation was also milled, injectionmolded, and extruded at low temperature and lowshear, to assess the thermomechanical history effects onthe orientation state. To produce the low shear/lowtemperature extrudate, the dryblend was extrudedusing a granulating device (Andouart monovis extruder40.20) set at 1608C. Part of these granules was injectionmolded using a Demag D85 NCIIIK equipment set at1958C. Finally, the dryblend was processed at 1908C,using a two-mill system (Collin) having 15 cm diameterrolls, the relative speed of the two rolls being V1/V2

¼ 1.1 and the gap between the rolls set at 0.35 mm.An annealing heat treatment was performed on

extruded sheets, consisting of a-1 h isothermal holdingat 1308C (i.e., above glass transition temperature (Tg)� 858C). The cooling condition effects were also stud-ied: after the above-mentioned annealing heat treat-ment, samples were cooled in icy water (iw), in air(air), or in the closed oven (ov), coming back slowly toambient temperature (return to ambient temperaturewas achieved in about 2 h). Physical ageing was per-formed on extruded sheets by holding the samples at708C (i.e., below Tg) for 23 h, followed by a coolingstep performed under air. The references of the differ-ent materials and their description are summarized inTable II.

Measurements

An accurate measurement of the real maximum melttemperature reached during the processing stage can

TABLE IWindow Profile Formulation (Quantities Given in phr)

PVC Lacovyl1 S110P (K ¼ 67) 100Lead one pack (lead-basedstabilizer þ lubricants þ process aid)

5.45

Durastrength1 320(acrylic impact modifier)

7.5

Titanium dioxide 5Coated calcium carbonate 6

2010 FILLOT ET AL.

Journal of Applied Polymer Science DOI 10.1002/app

be successfully assessed by DSC10,23: this tempera-ture referred to as ‘‘DSC real melt temperature’’ willbe used to characterize the gelation state of eachsample. The assessment of the ‘‘real melt tempera-ture’’ as well as the characterization of PVC aftervarious heat treatments were performed on a Q1000differential scanning calorimeter from TA Instru-ments. Twenty-milligram samples were subjected toa temperature scan from 50 to 2458C at a constantheating rate of 258C/min (optimized conditions tolimit PVC degradation). The real melt temperaturesof the five extruded sheets are given in Table III.

The same apparatus (Q1000 from TA instruments)was used in its modulated mode (MDSC) to separatethe reversible part from the nonreversible part of thetotal heat flow, the reversible part being associatedto heat capacity changes, and the nonreversible partto fusion/crystallization phenomena or enthalpicrelaxations. The mean heating rate in modulatedmode was 28C/min, the modulation amplitude beingset at 0.218C during a period of 40 s.

Mechanical behavior at small strains was character-ized with a mechanical spectrometer consisting of aforced oscillation torsion pendulum, working in thetemperature range of �1708C to 4308C and frequencyrange of 5 � 10�5 to 5 Hz. The storage (G0) and loss(G00) moduli of the complex shear modulus (G*) andthe loss factor (tan F ¼ G00/G0) were measured as afunction of temperature, between �1208C and 1508C,for a fixed frequency of 0.1 Hz. The dimensions of thesamples were about 30 mm � 7 mm � 2.5 mm.

Impact performances were determined according toa ‘‘modified’’ version of the ‘‘Charpy BS 0.1’’ standard.The only modification concerned the notch radius thatwas either 0.25 mm or 1 mm instead of 0.1 mm. Thisincrease of the notch radius shifts the ‘‘brittle/ductile’’transition and it allows the discrimination of samplesthat would appear ‘‘brittle’’ in BS 0.1. Impact perform-ances were evaluated at 238C with a noninstrumentedCharpy pendulum, the pendulum energy being 4 Jand the corresponding impact speed close to 2.9 m/s.The dimensions of the sample were 50 mm � 6 mm �2.5 mm. Samples were notched with a pneumatic de-vice (Notchvis from Ceast), the depth of the notchbeing 1.8 mm and its radius either 0.25 or 1 mm. Theenergy dissipated during impact solicitation wasrecorded, and at least 10 samples were tested for eachreference. Fracture was arbitrarily considered to be‘‘brittle’’ when the sample was broken into two pieces,and ‘‘ductile’’ (hinge break) otherwise. Afterwards, the‘‘impact performance’’ was calculated as the meanvalue of both ‘‘brittle’’ and ‘‘ductile’’ breaks strengths.

RESULTS AND DISCUSSION

Evidence of the PVC anisotropy

To evidence the magnitude of the anisotropy exist-ing in extruded sheets, charpy samples were takenin two different directions: samples whose lengthwas parallel to the extrusion direction were denomi-nated ‘‘a,’’ and samples whose length was perpen-dicular to the extrusion direction ‘‘b.’’ Both serieswere taken from the extrudates (EXT T1–T5 series).After samples notching (with r ¼ 0.25 mm), impactperformance was evaluated with the charpy pendu-lum and results are presented in Figure 1.

TABLE IIReferences (in Bold) and Descriptions of the Different

Materials

G Granulated sample—lowshear/low temperature

INJ Injection molded sampleM Milled sampleEXT Samples extruded at five different

temperatures: T1, T2, T3, T4, T5EXT Without heat-treatmentAg-EXT Physical ageing heat treatment <Tg

An-EXT Annealing heat treatment >Tg

An(iw)-EXT Cooled in icy waterAn(air)-EXT Cooled in airAn(ov)-EXT Cooled in oven

TABLE IIIReal Melt Temperature (8C 6 0.18C) of Extruded Sheets

Measured by Conventional DSC

EXT-T1 186.8EXT-T2 189.5EXT-T3 190.9EXT-T4 196.3EXT-T5 203.5

Figure 1 Charpy impact performances (notch radius¼ 0.25 mm) of extruded series as a function of the gelationlevel—specimens taken parallel or perpendicular to theextrusion direction.

THERMOMECHANICAL HISTORY EFFECTS ON PVC STRUCTURE/PROPERTIES 2011


The charpy impact performance exhibits a maxi-mum at intermediate DSC real melt temperature(i.e., at intermediate gelation level), which is in ac-cordance with many results reported in the litera-ture.6,7,8,24 Concerning the comparison of the sam-ples taken parallel or perpendicular to the extrusiondirection, a great difference is noticed. At intermedi-ate gelation level, 100% of the breaks of the ‘‘par-allel’’ series (a) are ‘‘ductile,’’ whereas for the ‘‘per-pendicular’’ series (b), 100% of the breaks are‘‘brittle,’’ whatever the gelation level. This resultreveals a strong anisotropy in extruded sheets, andthe effect of orientation on impact performanceseems to manifest even to a greater extent to that ofgelation.

Structural order in polymers is classically studiedby X-ray scattering,25 but it should be pointed outthat this technique only focuses on the crystallinepart of the material. To evidence both amorphousand crystalline oriented parts, birefringence meas-urements should be preferred.26 However, this tech-nique can only be applied on transparent formula-tions, i.e., it does not cover applications such as PVCwindow profiles or pipes. Another technique thatcould evidence orientation in rigid PVC is DSC.Illers27 and Gray and Gilbert28 reported the existenceof an exotherm immediately above the glass transi-tion. According to Illers, this exotherm, referred toas ‘‘cold crystallization peak’’ (CCP), emerges if PVChas been quenched from the molten state, generatinga quasi-amorphous glassy state that tends to crystal-lize if a sufficient amount of heat is brought. Morerecently, Yarahmadi19 observed a similar peak onDSC thermograms of rigid extruded PVC. In thatcase, the peak was attributed to the crystallization ofamorphous parts oriented during the process. Thus,CCP could be considered as a signature of orienta-tion, but as the peak emerges close to the glass tran-sition region (Fig. 2), the analysis of the thermo-

grams is difficult. The advantage of modulated DSC(MDSC) is precisely to separate the glass transitionfrom other phenomena. Therefore, MDSC could beadvantageously used to investigate CCP. The MDSCthermogram of an extruded sheet (EXT-T3) is re-ported in Figure 3, MDSC thermograms being simi-lar whatever the gelation level.

It can be noticed that the signal of Figure 3 associ-ated with the total heat flow is rather different fromthe signal of Figure 2 obtained in conventional DSC.This can be explained by the very low heating rateused in MDSC (28C/min), which leads to higher re-solution. On the reversible part of the MDSC signal,a clear glass transition is evidenced close to 858C.On the nonreversible part, four peaks (three endo-therms and one exotherm) are observed. To identifytheir origin, the formulation was also evaluated inits powder state (dryblend). The nonreversible partof the signal is given in Figure 4.

The three endotherms are present in both proc-essed and unprocessed PVC, whereas the exothermonly exists in processed PVC. Hence, the three endo-therms can be attributed to the components of theformulation, and after further investigations, itturned out that they were attributed to the fusion of

Figure 2 Thermogram of processed PVC (EXT-T3)obtained by conventional DSC at a heating rate of 258C/min(Exo down).

Figure 3 MDSC thermogram of processed PVC (EXT-T3),the total heat flow being the sum of the reversible and thenonreversible parts (Exo down).

Figure 4 MDSC thermogram (nonreversible part of thesignal) of unprocessed PVC (Exo down).

2012 FILLOT ET AL.


ingredients of the one pack (stabilizers þ lubricants).As the exotherm only appears on the thermogram ofprocessed PVC, it is consistent to attribute it to theCCP mentioned in the literature. As MDSC providesa clear separation of the glass transition from theother phenomena, a precise analysis of the enthalpyvalue of the CCP is allowed. At this stage, one canwonder whether the value of this enthalpy couldreflect a quantitative orientation level. To check thisassumption, samples of same formulation obtainedby different processing routes were characterized.The nonreversible part of the heat flow is given inFigure 5. In this figure, it can be noticed that thegranulated sample (processed at very low tempera-ture with low shear) does not exhibit any CCP, sug-gesting that this kind of soft process does not gener-ate any significant orientation of the macromole-cules. On the contrary, for the three other processingroutes, CCPs are observed and enthalpy values canbe calculated. The observed trend is that the strongerthe assumed orientation, the higher the enthalpyvalue. Indeed, the strongest orientation was assumedfor the injection molded sample and the lowest formilled sheets.

By quantifying the cold crystallization phenom-enon, MDSC appears as a promising technique tocharacterize orientation. However, it was earlierunderlined that CCP can also emerge on quenchedPVC samples. In the current work, the samples proc-essed by the different routes were not quenchedfrom the molten state, but they were cooled slowlyfrom a partially molten state. Moreover, their crystal-linity degrees were found to be very similar, asassessed in a previous work14 by DSC and X-raymeasurements. Thus, the differences observedbetween the CCP enthalpies associated to different

processing routes can be mainly attributed to orien-tation effects. To correlate MDSC results to thosegiven by another experimental technique, an inter-esting perspective may be to study the birefringenceproperties of a transparent PVC formulation.

Effects of a heat treatment below Tg

(physical ageing)

To simulate the effects of long storage periods ofPVC products, accelerated physical ageing can beperformed by heat-treating PVC extrudates justbelow the Tg. In the current work, extruded sheetswere aged at 708C for 23 h. Charpy impact perform-ances of samples taken parallel or perpendicular tothe extrusion direction were evaluated before and af-ter physical ageing. Results are given in Figure 6.The influence of physical ageing on the nonreversi-ble part of the heat-flow measured by MDSC is illus-trated in the case of EXT-T3 in Figure 7.

After physical ageing, a significant decrease in theimpact performances occurs, but results still exhibita strong anisotropy. Moreover, the fact that CCP stillexists on MDSC thermogram after physical ageingsuggests that heat treatment below Tg does not bringenough mobility to macromolecules to disorientate.In addition, physical ageing is highlighted on MDSCthermogram by the typical physical ageing endo-therm.29 It appears, in the present case, superim-posed on one of the endotherms associated to thefusion of the one-pack ingredients.

The effects of physical ageing on molecular mobil-ity are often investigated by dynamical mechanicalanalysis (DMA). The loss factor (tan F) of one of the

Figure 5 MDSC thermograms (nonreversible part) of (1)milled, (2) extruded, (3) injected, and (4) granulated sam-ples (Exo down).

Figure 6 Charpy impact performances (notch radius¼ 0.25 mm) of extruded series (a) parallel or (b) perpendic-ular to the extrusion direction as a function of the gelationlevel, before and after physical ageing.



extrudates (EXT-T3) was determined before and afterphysical ageing. Results are given in Figure 8.Between �1208C and 308C, a broad peak associatedto the b relaxation of PVC is observed, and it is notsignificantly affected by the physical ageing heattreatment. On the contrary, the low temperature partof the second peak, related to the a relaxation andobserved between 308C and 1208C, is deeply affectedby physical ageing. The loss factor decrease corre-sponds to a reduction of the mobility of PVC macro-molecules, explaining the impact performance lossobserved in Figure 6.

Moreover, the existence of a shoulder on the lowtemperature part of the a relaxation of the nonagedPVC can be noticed. In the PVC literature, the pres-ence of a shoulder close to 508C is sometimes men-tioned but not really discussed.30 However, Diaz-

Calleja31 reported, in the case of quenched samples,the existence of a ‘‘relaxation’’ occurring between thea- and b relaxation of PVC. According to the author,this process is not a true relaxation but the result ofa quenched state that exhibits higher molecular mo-bility. This aspect will be discussed in the followingpart, the effects of a heat treatment above Tg as wellas the influence of the cooling rate being investi-gated.

Effects of a heat treatment above Tg

Extruded sheets presenting different gelation levelswere annealed at 1308C for 1 h, and after this heattreatment, samples were cooled in icy water (iw), inair (air), or in oven, coming back slowly to ambienttemperature (ov). The charpy impact performancesof the series cooled in air were first evaluated, sam-ples being taken either parallel or perpendicular tothe extrusion direction. Results before and afterannealing are reported in Figure 9, as a function ofthe gelation level.

After annealing above Tg, charpy impact perform-ances decrease drastically. For instance, when Fig-ures 6 and 9 are compared, it can be noticed that theperformance decrease is more pronounced for an-nealing than for physical ageing. A first explanationof the performance decrease after annealing at 1308Ccould be related to the disorientation of the polymerchains. Indeed, after the annealing heat treatment, ashrinkage of about 2% along the extrusion directionwas measured. This shrinkage is a direct conse-quence of the macromolecular chains reorganizationallowed by the mobility increase provided by heat-

Figure 7 MDSC thermograms (nonreversible part) ofprocessed PVC (EXT-T3) (1) before and (2) after physicalageing (Exo down).

Figure 8 Evolution of the loss factor (tan F) with the tem-perature of processed PVC (EXT-T3) before and after phys-ical ageing.

Figure 9 Charpy impact performances (notch radius¼ 0.25 mm) of extruded series (a) parallel or (b) perpendic-ular to the extrusion direction as a function of the gelationlevel, before and after annealing and cooling in air.

2014 FILLOT ET AL.


ing above Tg. As MDSC yields information about theorientation state, annealed PVC was also character-ized by MDSC. The thermograms before and afterannealing are given in Figure 10.

After the annealing heat treatment, the CCP com-pletely disappeared. From this observation, it couldbe concluded that (i) PVC is totally disoriented or(ii) cold crystallization took place during the heattreatment at 1308C. In the latter case, this suppressesany further possibility of cold crystallization duringthe MDSC test. From Figure 9, it seems that someanisotropy still exists after annealing. But as both se-ries ‘‘parallel’’ and ‘‘perpendicular’’ exhibit fully‘‘brittle’’ breaks, it is difficult to conclude. To betterdiscriminate these series, the severity of the charpytest was reduced by increasing the notch radius

from 0.25 to 1 mm. Impact performances of annealedPVC evaluated in these latter conditions are given inFigure 11.

The results highlight the subsistence of anisotropyafter annealing at 1308C. This demonstrates thatsuch thermal heat treatment does not allow a com-plete disorientation of the PVC chains. At this stage,one can wonder whether the dramatic impact per-formance loss observed after annealing can only beexplained by a partial disorientation. After the heattreatment above Tg, as PVC has been cooled fromthe rubbery state, an influence of cooling rate on themolecular mobility is expected. To check thisassumption, viscoelastic properties and impact per-formances of annealed PVC cooled in three differentconditions were assessed. The DMA thermogramswere first analyzed by focusing on the previouslymentioned shoulder linked to quenching. The evolu-tion of the loss factor (tan F) with the temperature isgiven in Figure 12 for the three different cooling con-ditions, the reference extruded sample before heattreatment also being reported.

In the Figure 12, it can be noticed that the fasterthe cooling rate, the more pronounced the shoulder.Thus, the shoulder appears as being a signature of aquenched state exhibiting a high molecular mobility.It has been shown in the previous part that thereduction of this shoulder caused by physical ageingwas correlated to an impact performance decrease.As expected and shown in Figure 13, a similarbehavior is also observed when the reduction of theshoulder is caused by a slower cooling rate: thefaster the cooling rate, the more pronounced theshoulder, and the higher the impact performance.

Since the influence of cooling conditions on impactproperties has been clearly evidenced, the coolingconditions after the 1308C heat treatment has also tobe taken into account, together with the disorienta-

Figure 10 MDSC thermograms (nonreversible part) ofprocessed PVC (EXT-T4) (1) before and (2) after annealingand cooling in air (Exo down).

Figure 11 Charpy impact performances (notch radius ¼ 1mm) of annealed extruded series (a) parallel or (b) perpen-dicular to the extrusion direction as a function of the gela-tion level.

Figure 12 Evolution of the loss factor (tan F) with thetemperature of (1) EXT-T3, (2) An(iw)-EXT-T3, (3) An(air)-EXT-T3, and (4) An(ov)-EXT-T3.



tion, to explain the impact loss observed in Figure 9.Indeed, before annealing, the molecular mobility isdetermined by the cooling conditions that follow theextrusion, and after annealing, molecular mobility isdetermined by the cooling conditions that follow theannealing heat treatment. These two cases corre-spond, respectively, to the DMA thermograms 1 and3 of the Figure 12. In this figure, it can be seen thatthe shoulder before annealing (thermogram 1) ismore pronounced than the shoulder after annealingand cooling in air (thermogram 3). This means thatthe cooling rate of the extrusion process is fasterthan standard air cooling. As a consequence,annealed samples exhibit a lower molecular mobil-ity, and this reduced mobility could be part of theorigin of the significant impact performance lossobserved after annealing.

The impact performance of a PVC sample that hasbeen heat-treated above Tg seems to be affected by(i) PVC partial disorientation and (ii) reduced molec-ular mobility caused by cooling conditions. But inaddition to these two effects, a third one related toPVC crystallinity could be considered. Indeed, dur-ing the heat treatment above Tg, isothermal crystalli-zation occurs, as evidenced by DSC thermogramspresented in Figure 14.

In this figure, thermogram 1 corresponds to thePVC reference sample (EXT-T3) before annealing,and the two well-known endotherms related to pri-mary and secondary crystallites of processed PVCare observed.23 After annealing either at 1108C (ther-mogram 2) or at 1308C (thermogram 3), a third peakis superimposed on the endotherm related to sec-ondary crystallites, highlighting the isothermalgrowth of crystallites that melt at temperaturesslightly above the heat treatment temperature. Thisphenomenon has been reported in the literature byseveral authors,25,26,32 but few studies have evi-denced the influence of this crystallinity modification

on impact properties. In the work of Kim and Gil-bert,26 crystallites isothermal growth appears toaffect more the ability of PVC to resist to shrinkage,than the mechanical properties. However, one canwonder whether the isothermal growth of crystallitesdoes not affect the molecular mobility of PVC. InFigure 12, it can be seen that the height of the lossfactor peak decreases after annealing. Flores-Flores33

also evidenced in rigid PVC such a decrease after aheat treatment slightly higher than Tg, and he attrib-uted it to crystallites isothermal growth. If thisassumption looks quite consistent with the results ofthe present work, the decrease of the loss factorcould also result from the reorganization of theamorphous polymer chains (partial disorientation).Further investigations would be needed to definitelyconclude on the origin of the loss factor decrease af-ter annealing.

CONCLUSIONS

The sensitivity of rigid PVC microstructure and im-pact strength to its thermomechanical history hasbeen investigated. Whatever the gelation level, rigidPVC processed by extrusion may exhibit a strongimpact strength anisotropy effect, which is related tomacromolecular orientation occurring along extrusiondirection. This orientation phenomenon has been suc-cessfully highlighted by MDSC experiments, sinceoriented PVC chains exhibit a crystallization exother-mal peak upon heating (so-called cold crystallizationpeak). Moreover, the enthalpy of this exothermappears as being directly correlated to the extent ofmacromolecular orientation. When a rigid PVC extru-date is heat-treated below its Tg, a significant impactstrength loss is observed without significant modifica-

Figure 13 Charpy impact performances (notch radius ¼ 1mm) of (1) An(ov)-EXT-T3, (2) An(air)-EXT-T3, and (3)An(iw)-EXT-T3.

Figure 14 DSC thermograms of EXT-T3 (1) before and af-ter annealing for 1 h at (2) 1108C and (3) 1308C (Exodown).

2016 FILLOT ET AL.


tion of the PVC orientation state. This performancedecrease is associated to a reduction of the molecularmobility (typical physical ageing effect), as evidencedby DMA experiments in the low temperature part ofthe a mechanical relaxation. If rigid PVC is heat-treated above its Tg, an even more pronounced impactperformance loss is observed. The annealing heat-treatment above Tg provides a complex modificationof the microstructure that can be seen as a superimpo-sition of three possible effects: (i) a partial disorienta-tion of the PVC chains, (ii) a modification of the mo-lecular mobility depending on the cooling rate afterthe heat-treatment, and (iii) a reduction of the molecu-lar mobility induced by the isothermal growth of PVCcrystallites. Disorientation as well as molecular mobil-ity reduction contribute to the overall impact per-formance loss observed after having annealed PVCabove its Tg. Finally, it emerges from this work thatDSC could be advantageously used in its modulatedmode to assess the orientation state of PVC. And froma more general point of view, this work highlights thefact that impact performance of PVC is not only gov-erned by its gelation level, it also strongly depends onboth orientation state and molecular mobility.

The authors express their thanks to ARKEMA, and moreparticularly to its Plastic Additives group, for permissionto publish this work.

References

1. Hattori, T.; Tanaka, K.; Matsuo, M. PolymEng Sci 1972, 12,199.

2. Menges, G.; Berndtsen, N.; Opfermann, J. J Plast Rubber Pro-cess 1979, 4, 156.

3. Krzewki, R. J.; Collins, E. A. J Macromol Sci Phys 1981, 20,443.

4. Gilbert, M.; Vyvoda, J. C. Polymer 1981, 22, 1134.5. Terselius, B.; Jansson, J. F. Plast Rubber Process Appl 1983, 4,

291.

6. Terselius, B.; Jansson, J. F.; Bystedt, J. Plast Rubber ProcessAppl 1985, 5, 1.

7. Hajji, P.; Marchand, F.; Gerard, P.; Gauthier, C.In Proceedingsof the PVC Conference, Towards a Sustainable Future, Brigh-ton, UK, IOM Communication: London, 2002.

8. Cora, B.; Daumas, B.; Zegers, A. Plast Rubber Compos 1999,28, 165.

9. Gonze, A. Chimie et Industrie Genie chimique 1971, 104, 422.10. Gilbert, M.; Hemsley, D. A.; Miadonye, A. Plast Rubber Com-

pos Process Appl 1983, 3, 343.11. Potente, H.; Schultheis, S. M.; Gollner, M. Kunstoffe 1988, 78, 641.12. Summers, J. W.; Rabinovitch, E. B.; Booth, P. C. J Vinyl Tech-

nol 1986, 8, 2.13. Covas, J. A. Plast Rubber Process Appl 1988, 9, 91.14. Fillot, L. A.; Hajji, P.; Gauthier, C.; Masenelli-Varlot, K. J Vinyl

Additive Technol, to appear.15. Chapman, P. G.; Agren, L. Presented at the Proceedings of the

Plastics Pipes X Conference, Gothenburg, Sweden, Sept. 14–17,1998.

16. West, D. B.; Truss, R. W. J Mater Sci 2004, 39, 2789.17. Hitt, D. J.; Gilbert, M. Polym Test 2000, 19, 27.18. Hitt, D. J.; Gilbert, M. Plast Rubber Compos 2000, 29, 149.19. Yarahmadi, N.; Jakubowicz, I.; Hjertberg, T. Polym Degrad

Stab 2003, 82, 59.20. Perez, J. In Physique et Mecanique des Polymeres Amorphes;

Tec et Doc Eds.: Paris, 1992; 384 p.21. Zerafati, S.; Black, J. J Vinyl Additive Technol 1998, 4, 240.22. Rabinovitch, E. B.; Summers, J. W. J Vinyl Additive Technol

1992, 14, 126.23. Fillot, L. A.; Hajji, P.; Gauthier, C. J Vinyl Additive Technol, to

appear.24. Calvert, D. J.; Haworth, B.; Stephenson, R. C. Plast Rubber

Compos Process Appl 1991, 15, 229.25. Dawson, P. C.; Gilbert, M.; Maddams, W. F. J Polym Sci Part

B: Polym Phys 1991, 29, 1407.26. Kim, H. C.; Gilbert, M. Polymer 2004, 45, 7293.27. Illers, K. H. J Macromol Sci Phys 1977, 14, 471.28. Gray, A.; Gilbert, M. Polymer 1976, 17, 44.29. Hutchinson, J. M. Prog Polym Sci 1995, 20, 703.30. Perera, M. C. S.; Ishiaku, U. S.; Ishak, Z. A. M. Eur Polym J

2001, 37, 167.31. Diaz-Calleja, R. Polym Eng Sci 1979, 19, 596.32. Covas, J. A. Polymer 1993, 34, 3204.33. Flores-Flores, R. Comportement Mecanique du PVC et de Mel-

anges PVC/PMMA—Effet de Traitements Thermiques et de laReticulation Chimique, Ph.D. Thesis, INSA de Lyon: France, 1994.



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