time–temperature-dependent behavior of a substituted poly(paraphenylene): tensile, creep, and...

Time–Temperature-Dependent Behavior of a SubstitutedPoly(Paraphenylene): Tensile, Creep, and DynamicMechanical Properties in the Glassy State

DERRICK DEAN,1* MARK HUSBAND,1 MARK TRIMMER2

1 Materials Directorate of Wright Laboratory, Wright-Patterson AFB, Ohio 45433

2 Maxdem Incorporated, 140 E. Arrow Highway, San Dimas, California 91773

Received 3 September 1997; revised 11 June 1998; accepted 16 June 1998

ABSTRACT: The linear viscoelastic behavior of a poly(paraphenylene) with a benzoylsubstituent has been examined using tensile, dynamic mechanical, and creep experi-ments. This amorphous polymer was shown to have a tensile modulus of 1–1.5 Msi,nearly twice that of most common engineering thermoplastics. The relaxation behavior,which is similar to that of common thermoplastics, can be described by the WLFequation. Outstanding creep resistance was observed at low temperatures, with rub-bery-like behavior being exhibited as the temperature approached Tg. Physical agingwas shown to interact with long-term creep, rendering time–temperature superpositioninvalid for predicting the long-term properties. The effect of physical aging on the creepbehavior was characterized by the shift rate m. © 1998 John Wiley & Sons, Inc. J Polym SciB: Polym Phys 70: 2971–2979, 1998Keywords: poly(paraphenylenes); viscoelastic; creep; physical aging

INTRODUCTION

High-performance thermoplastic polymers arefinding an interesting number of applications,ranging from composite matrix materials to auto-mobile engine components. These applicationstypically require high thermal stability, highstrength and stiffness, as well as processability.Poly(paraphenylenes) and their derivatives arerecognized as an emerging family of heat-resis-tant, high modulus resins that may meet theserequirements. The synthesis of poly(paraphen-ylenes) has, in fact, been the subject of muchresearch over the past twenty years or so becausethey are expected to exhibit a number of desirable

properties, such as high mechanical strength andstiffness, good electrical conductivity and goodthermal and thermooxidative stability. Unfortu-nately, most synthetic attempts have yieldedpolymers that were intractable1 or actually lowmolecular weight oligomers.

Alternative approaches to make poly(para-phenylenes) substituted by such groups as longchain alkyl,2 aryl,3 carboxyl,4 and ester groups5

have yielded soluble polymers, but with low tomoderate molecular weights (e.g., intrinsic viscos-ity less than 1.0 dL/g). Many of the synthetichurdles have recently been overcome, resulting inthe preparation of high molecular weight, substi-tuted poly(paraphenylenes).6–8 These polymers,possessing para-linked phenyl units with benzoylsubstituents randomly distributed along thebackbone are known collectively as polyX-(TM).This interesting molecular architecture results inpolymers that are amorphous, exhibit thermo-plastic behavior, and possess a number of inter-

Correspondence to: D. Dean, Advanced Technology Group,BF Goodrich, 9921 Brecksville Road, Brecksville, OH 44141

* Current address: BF Goodrich, 9921 Brecksville Road,Brecksville, OH 44141Journal of Polymer Science: Part B: Polymer Physics, Vol. 70, 2971–2979 (1998)© 1998 John Wiley & Sons, Inc. CCC 0887-6266/98/162971-09

2971

esting properties. In addition to high thermal andthermooxidative stability, films and molded sam-ples exhibit isotropic (i.e., no orientation) tensileand flexural moduli of 1–1.5 Msi, comparable toliquid crystalline polymers which have been ori-ented along the machine direction. They also pos-sess surprisingly low glass transition tempera-tures and solubility in a number of common or-ganic solvents, making them readily processible.These properties make them attractive for a num-ber of applications, in both the military and civil-ian sectors, such as high compressive strengthcomposites, high modulus thermoplastic fibers,high performance films, abrasion-resistant coat-ings, and as additives to enhance the properties ofother polymers.

Very little work on the structure-property re-lationships of these substituted poly(paraphen-ylenes) has been published. A recent study byConnolly et al. examined the dielectric and thedynamic mechanical relaxation behavior of poly-(paraphenylenes) substituted with benzoyl and4-phenoxybenzoyl groups groups as well as acopolymer of benzoyl-1,4-phenylene and 1,3-phenylene.9a They found that the temperaturedependence of viscoelastic and dielectric relax-ation times were well described by the Williams-Landel-Ferry and Vogel-Fulcher equations, re-spectively. They explained the temperature de-pendence of the relaxation times by simple freevolume considerations; the temperature sensitiv-ity of the relaxation times was shown to correlatenot with the total free volume but with the relativechange in free volume with respect to the free vol-ume at Tg. Recent work by Simon et al., has utilizedpositron annihilation spectroscopy and thermallystimulated current to extend the work of Connollyet al. in investigating the molecular mobility ofthese substituted poly(p-phenylenes).9b

We have initiated a program to systematicallycharacterize the structure-property-processingrelationships of these polymers in both the solidand molten state. We report here on the effects oftime and temperature on the tensile, creep anddynamic mechanical properties of benzoyl-substi-tuted poly (1,4-phenylene) (Fig. 1) in the glassystate. As mentioned earlier, its outstanding me-chanical properties makes this polymer a poten-tial candidate for structural applications. In suchuses, the mechanical behavior of the material as afunction of time and temperature are of majorconcern. In particular, creep is known to be adrawback to use of most thermoplastics, espe-cially amorphous, thermoplastic resins which

have no chemical cross links or crystalline regionsto retard the creep process. Polysulfones areknown to exhibit very good creep resistance. Thishas been attributed to the repeating phenyl rings,which create both steric hindrance to rotationwithin the molecule and electronic attraction ofresonating electron systems between adjacentmolecules.10 We also expect good creep resistanceto be exhibited by the substituted poly(para-phenylenes), for similar reasons.

Creep studies of engineering thermoplasticshave all shown that as the temperature ap-proaches Tg, the creep rate increases. However,efforts to use time–temperature superposition topredict long-term creep behavior from short-termtests have been shown to be inaccurate, due tothe interaction of physical aging during creep.The concept of physical aging in polymers canbe described by considering free volume and en-thalpy.10 Above Tg, both quantities are at ther-modynamic equilibrium. When a polymer iscooled below Tg however, these quantities are nolonger at equilibrium. Although all long-rangemotion is retarded, the polymer continues to movetoward equilibrium through small scale changes,losing free volume and enthalpy in the process.This process is known as physical aging.

The reduction in free volume reduces the poly-mer mobility and increases the relaxation times.A distinguishing characteristic of physical agingis that it is thermally reversible, that is, its effectscan be erased by heating the polymer above Tg.This process is often referred to as rejuvenation.During the process of aging, the polymer contin-ually loses heat, although the magnitude is small.This loss of heat is commonly known as enthalpyrelaxation. When the material is heated above Tg,it absorbs the heat lost during aging. This gives

Figure 1. Chemical structure of PX™-1000 resin.

2972 DEAN, HUSBAND, AND TRIMMER

an endothermic peak near Tg in the DSC scanwhich is not seen in an unaged or rejuvenatedsample.

Struik and others have studied the phenome-non of physical aging in most common poly-mers,11–12 with studies of its effect on propertiesof newer engineering thermoplastics being re-ported recently.13–16 The physical aging behaviorof polyetherimide (PEI) has been studied by sev-eral authors. Matsumoto has shown that time-temperature superposition is invalid for predict-ing long-term creep properties of PEI, as well aspolycarbonate,16 with the actual data differingfrom the predicted data by a factor of 10.

Echeverria et al. have used creep and differentialscanning calorimetry (DSC) to study the physicalaging behavior of PEI also.17 Andrade plots of thecompliance versus the cube root of time were foundto be linear at short times with slope b decreasingwith increasing aging time to a constant value onceequilibrium is reached, while the enthalpy changeincreased linearly with aging time, leveling off asequilibrium is reached. They determined that thetime scales to reach equilibrium for enthalpy andfor mechanical measurements were the samewithin experimental error.

Brennan and Feller used DSC, dynamic me-chanical analysis, and room temperature tensiletests to study the effects of physical aging on thethermal and mechanical behavior of PEI.18 Theydemonstrated that PEI has a pronounced b re-laxation which exhibits a significant decreasein magnitude with aging, paralleling dramaticchanges in room temperature stress-elongationbehavior.

The physical aging behavior of PEEK, bothsemicrystalline and amorphous versions, has alsobeen extensively studied by a number of research-ers,15,19,20 using creep, DSC, and DMA. Ogale andMcCullough found that the degree of crystallinitydoes affect the amount of aging that the polymerundergoes,15 while D’Amore et al.20 found thatinterpreting physical aging above Tg in amor-phous PEEK was complicated by cold crystal-lization, which the polymer undergoes justabove Tg.20

EXPERIMENTAL

Materials

The benzoyl-substituted poly(1,4-phenylene) wassupplied by MAXDEM, INC. (PX™-1000 resin).

Synthetic details are published elsewhere.6 Thepolymer used in this study had an approximateMn of 17,000. Samples used for dynamic mechan-ical analysis were compression-molded rectangu-lar bars with dimensions of 1.5 mm thick, 9.5 mmwide, and 43 mm long. Samples were molded in aCarver press at 285°C under 5000 pounds of pres-sure for 15 min. Dogbone specimens used for ten-sile testing were cut from compression-moldedsheets of the polymer. Specimens had a gagelength of 25.7 mm, width of 1.5 mm and a thick-ness of 3 mm. All samples were annealed at 200°C(ca. 25°C above the Tg) for 30 min, to erase anyprevious thermal history, then quenched to roomtemperature before testing. Some samples werephysically aged before testing. This was done byfirst heating the sample to 200°C, quenching tothe desired aging temperature, and annealing atthat temperature for specified amounts of time.

Instrumental Methods

Differential scanning calorimetry experimentswere performed on a Perkin-Elmer DSC-7 using ascan rate of 10°C/min. Sample weights were 5–10mg. Tensile tests were performed on an MTSmodel 312.21 test frame equipped with an envi-ronmental chamber. Tests were performed attemperatures ranging from 25 to 150°C. An MTSextensometer was used to measure the strain.Tensile fracture surfaces were examined using aJEOL-840 scanning electron microscope (SEM).Dynamic mechanical analysis (DMA) was per-formed on a Rheometrics RDS 770 rheometer inthe torsion rectangular mode at temperaturesranging from 250–230°C. Strain sweeps wereperformed at various temperatures to determinethe linear viscoelastic range. All of the DMA ex-periments discussed here were performed at astrain of 0.1%. Flexural creep tests were per-formed on a Perkin-Elmer DMA-7 at tempera-tures ranging from room temperature to 150°C.The linear viscoelastic range for each tempera-ture was determined by performing isothermalcreep tests at various loads.

RESULTS AND DISCUSSION

Tensile Behavior

Room temperature tensile and flexural tests haveshown PX™-1000 samples to have isotropic mod-ulus values of 1 to 1.5 Msi, nearly double the

TIME-TEMPERATURE-DEPENDENT BEHAVIOR OF PPL 2973

value for most engineering thermoplastics, andtensile strengths of 20–25 Ksi. It is noted thatmodulus values twice as high as those reportedhere as well as substantially higher tensilestrength values have been attained utilizing care-fully machined samples.7b We did not optimizethe processing conditions for the samples utilizedin this study. We have studied the tensile defor-mation behavior of this polymer at various tem-peratures from 23 to 150°C (15°C below its Tg).The results are summarized by the stress-straincurves in Figure 2, with modulus and strengthvalues displayed in Table I. The tensile behavioris typical of a semiflexible backbone, thermoplas-tic polymer, exhibiting brittle failure with rela-tively low elongation at room temperature, withsome yielding and necking seen at 120°C, fol-lowed by rubbery-like behavior and ductile failureat the higher temperatures. The SEM fracto-graphs in Figure 3 show crazing and void forma-tion, typical of brittle failure at 75°C, and ductileflow, due to chain slippage at 130°C.

Dynamic Mechanical Studies

Dynamic mechanical experiments involve apply-ing an oscillatory strain to a sample while moni-

toring the resultant stress, which consists of bothin-phase and out-of-phase components.22 Thesestresses can then be used to calculate the in-phase (G9) and out-of-phase (G0) components ofthe modulus. Figure 4 shows the shear storagemodulus G9, and the shear loss modulus, G0, as afunction of temperature. The dramatic drop inmodulus indicates the onset of Tg. The point atwhich the G9 and G0 curves intersect is taken asTg. This value, ca. 165°C, corresponds well withthat reported in the literature for this polymer.7–9

This surprisingly low Tg is apparently related tothe nature of the side group, both its bulkiness aswell as its regioplacement along the backbone.

Figure 2. Tensile properties of PX™-1000 samples asa function of temperature.

Table I. Tensile Properties As a Functionof Temperature for PX™-1000

Temperature(°C)

Modulus(Mpsi)

Strength(Kpsi)

23 1.23 19.5775 1.00 17.7

100 0.75 13.7130 0.40 11.6140 0.40 10.6

Figure 3. (a) SEM of tensile fracture surface at 75°C,showing crazing and void formation. (b) SEM of tensilefailure surface at 130°C, showing considerable chainslippage and ductile flow.


Connolly et al. have shown that the size of thesubstituent dictates the free volume and thus theTg, with the larger substituent polymer (phenoxy-benzoyl versus benzoyl) having less free volumeand a lower Tg. In some related work, Wang et al.have synthesized related versions of poly(benzoyl-1,4-phenylenes) in which they maintain strictcontrol of the regioplacement of the benzoylgroups, resulting in higher Tg values (170–210°C) compared to values in the range of 150–165°C for samples with random placement of thependant group.8b A very weak, broad secondaryrelaxation has been reported for this polymer byDMA, in the temperature range of 2100 to260°C.9 We have observed a similar relaxation,barely distinguishable from the experimentalnoise, but it is reproducible and totally disappearswhen temperature scans are run at frequencies of6.28 rad/s or higher. It is noted that both Connollyand Simon have observed a much more pro-nounced relaxation peak in this temperaturerange, using dielectric relaxation spectroscopy.9

In order to further characterize the relaxationbehavior of this polymer, a series of isothermalfrequency sweeps were done from 250 to 230°C.For clarity, only data for temperatures between136 and 226°C are shown in Figure 5. The storagecompliance J9 is plotted versus frequency. Thebehavior shown is that typical of a thermoplasticpolymer. At low frequencies, the polymer exhibitshigh compliance, rubbery-like behavior and athigher frequencies it exhibits low compliance,glassy-like behavior. It can also be seen from thisfigure that at a given frequency the storage com-pliance increases with temperature. It should be

noted that the shear moduli are approximately 13

that of the tensile moduli, as seen in most poly-mers.22 While the data in Figure 5 gives muchinformation for the range of frequencies studied,it is useful to understand the behavior over amuch broader frequency range. This can be done,using time temperature superposition.

Time-temperature superposition is based onthe fact that the viscoelastic behavior at one tem-perature can be related to that at another tem-perature by a change in the time or frequencyscale only, assuming that the material behavior isthermorheologically simple.22 For dynamic me-chanical data, this implies that by selecting areference temperature, one can shift higher tem-perature curves to lower frequencies and lowertemperature curves to higher frequencies,thereby constructing a master curve. The mastercurve covers a much wider range of frequenciesthan those experimentally accessible. A mastercurve for the PX™-1000 material is shown inFigure 6. For most polymers, the master curveobtained by time-temperature superposition ofdata obtained in the vicinity of Tg is well de-scribed by the Williams-Landell-Ferry (WLF)equation: log at 5 OC1(T 2 Tr)/C2 1 (T2 Tr). This is also true in our case. The mastercurve has been fitted to the WLF equation and aplot of the shift factors, at, used to construct themaster curves, versus temperature is shown inFigure 7. An activation energy for this viscoelas-tic relaxation can also be calculated, using theequation: DHa 5 2.303RC1C2T2/(C2 1 T2 Tr)

2, where R is the universal gas constant.22

The value we obtain, 139 kcal/mol, agrees well

Figure 5. Storage compliance versus frequency atvarious temperatures for PX™-1000 samples.

Figure 4. DMA temperature scan (at a rate of 0.1rad/s) of PX™-1000 resin, indicating a Tg of > 165°C.


with that reported by Connolly et al.,9 as well asvalues obtained from molecular dynamics simu-lations.23 This activation energy value is nearly3.5 times that for polystyrene.22

CREEP BEHAVIOR

In order to fully characterize the creep behavior ofa polymer, tests must be conducted over a rangeof temperatures and load levels. In this study, wewanted to ensure that all of the loads chosen werein linear viscoelastic range. The linear viscoelas-tic behavior of polymers can be described by theBoltzmann superposition principle, which allowsthe state of stress or strain in a viscoelastic bodyto be determined from knowledge of its entiredeformation history.24 The basic assumption isthat during viscoelastic deformation in which the

stress is varied, the overall deformation can bedetermined by summing up the strains due toeach load step.

This concept can be demonstrated rather easilyby performing creep tests at different stresses.The linear viscoelastic range will be the range inwhich the creep compliance remains independentof the stress. An example is shown in Figure 8,which shows the creep compliance as a function ofstress at the various temperatures of interest inthis experiment. At the lower temperatures thelinear viscoelastic range is rather large, but be-comes smaller as the temperature increases.

Figure 9 shows creep strain versus time forPX™-1000 samples at various load levels at120°C, and is representative of what we haveobserved at the other temperatures. The curvesdisplay an initial elastic deformation in responseto the application of the load, followed by the

Figure 6. Storage compliance master curve, Tref

5 161°C.

Figure 7. WLF plot of shift factors used to constructthe master curve shown in Fig. 6.

Figure 8. Creep compliance versus stress at varioustemperatures.

Figure 9. Creep strain for various load levels at120°C for PX™-1000 resin.


much slower viscoelastic deformation. PX™-1000samples exhibit excellent creep resistance at thelower load levels, with the creep strain becomingmore significant as the stress increases. It can beseen that the level of strain for the curve tested at0.6 Kpsi is roughly 0.6 times that shown in thecurve tested at 1 Kpsi, in accordance with theBoltzmann superposition principle. A representa-tive plot of the creep behavior from room temper-ature to 150°C is shown in Figure 10, whichshows creep compliance versus time, plotted dou-ble-logarithmically. The creep behavior of PX™-1000 resin is similar to that of other, well-studiedhigh-performance thermoplastics in that the levelof creep is very small at the lower temperatures,but as the temperature increases, the relaxationrates increase and a higher level of complianceresults. Examination of the creep curves at 130,

140, and 150°C, show that the creep rates aresimilar, however. At about 10,000 s as the phys-ical aging begins to interact with the creep, thecreep rate begins to decrease and, at 150°C, be-gins to level off. This fact, and the overall shape ofthe curves, especially at 150°C, are due to theonset of physical aging during the creep process.It is this phenomenon that causes time-tempera-ture superposition to inaccurately predict thecreep behavior over long time periods. We haveexamined the effect of physical aging using creepand DSC. By aging samples for varying times at130°C and subsequently performing creep testson them at a temperature lower than the agingtemperature, 120°C, we can see in Figure 11 thatas the aging time increases, the polymer becomesstiffer, resulting in lower compliance values. Incontrast to the curves for the samples that werenot aged prior to testing, the creep rates of thecurves in Figure 11 are different for each agingtime, continually increasing with increased agingtime, as you would expect if no aging were occur-ring and retarding the creep process. According toStruik,11 all polymers age in a similar way, andall mechanical relaxation times change in a sim-ilar way. Therefore, as the aging time, te, in-creases, the relaxations increase, resulting in ashift of the creep curve along the log time scale.This effect can be characterized by the shift rate,m, defined as

m 5 2d log a/d log te

where a denotes the shift factors and te denotesthe aging times. A double logrithmic plot of the

Figure 11. Creep compliance versus aging times forPX™-1000 samples aged at 130°C and tested at 120°C.

Figure 12. Log ate vs. logte for the creep curvesshown in Figure 11. The slope gives the shift factor, m.

Figure 10. Creep compliance at various tempera-tures for PX™-1000 resin.


shift factors from the data in Figure 11 versus therespective aging times is shown in Figure 12. Alinear fit of the data gives a slope, m, of 0.74,indicative of moderate aging behavior as definedby Struik.11 DSC experiments, shown in Figure13 show that as the aging time increases, the areaunder the peak near Tg, caused by enthalpic re-laxation, increases. This behavior is characteris-tic of typical, flexible backbone type polymers.Both creep and physical aging are known to occurvia small-scale molecular motions in such poly-mers. Although this particular polymer has a sig-nificantly more rigid backbone, we can attributethe effects we observe to changes in the free vol-ume as a function of aging.9,23 A more in depthstudy of the physical aging behavior of this andseveral related substituted poly(paraphenylenes)is currently underway.

SUMMARY

The linear viscoelastic behavior of a substitutedpoly(paraphenylene) has been examined in theglassy state, using tensile testing, dynamic me-chanical analysis, creep, and DSC. Although thispolymer has a rigid backbone, the Tg is surpris-ingly low, 165°C, and the relaxation behavior ischaracteristic of conventional, flexible backbonepolymers. Tensile tests indicate a modulus atroom temperature of 1–1.5 Msi, twice that of con-ventional engineering thermoplastic resins. Inaddition to displaying good creep resistance, long-term creep and DSC measurements indicate thatthe interaction of physical aging affects the long-term properties below Tg, with cessation of creepbeing observed after about 10,000 seconds for alltemperatures studied. Physically aging samples

before testing, however, results in a continualincrease in the creep rate during testing. Theeffect of physical aging on the creep behavior wascharacterized by its shift rate, m, which has avalue of 0.74, indicative of moderate aging, asdefined by Struik. It is noted that creep and phys-ical aging occur via small-scale molecular mo-tions. Given the rigidity of the backbone of thispolymer however, both the degree of creep andphysical aging are presumed to depend primarilyon the presence of free volume.

REFERENCES AND NOTES

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Figure 13. Enthalpy relaxation as a function of agingtimes for samples aged at 130°C.


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time–temperature-dependent behavior of a substituted poly(paraphenylene): tensile, creep, and...

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