Polymers for Fibers

8 Polymers for Fibers JOSEPH ZIMMERMAN Wilmington, DE, USA 8.1

249

INTRODUCTION

8.2 POLYMERIZATION AND ITS CONTROL 8.2.1 Polyamides 8.2.1.1 Nylon 6,6 8.2.1.2 Nylon 6 8.2.2 Poly (ethylene terephthalate) 8.2.3 Polymers Used in High Performance Fibers 8.2.3.1 Poly ( 1,4-phenyleneterephthalamide) 8.2.3.2 Thermotropic polyarylates 8.2.4 Polyacrylonitrile and Polyalkenes

250 250 250 253 254 256 256 257 259

8.3 SPINNING 8.3.1 Introduction 8.3.2 Threadline Tension 8.3.3 Draw Resonance 8.3.4 Structure Development in Spinning 8.3.4.1 Introduction 8.3.4.2 Nylon 6 and nylon 6,6 8.3.4.3 Poly(ethylene terephthalate) 8.3.4.4 Isotactic polypropylene 8.3.4.5 Gel-spun polyethylene 8.3.4.6 Liquid crystalline solutions and melts

260 260 260 263 263 263 265 266 267 268 269

8.4

271

DRAWING

8.5 STRUCTURE-PROPERTY RELATIONSHIPS 8.5.1 Modulus 8.5.2 Tensile Strength 8.5.3 Compressive Strength 8.5.4 Creep

272 272 274 277 277

8.6 FIBER DURABILITY 8.6.1 Introduction 8.6.2 Tensile Fatigue 8.6.3 Flex and Abrasion 8.6.4 Degradation

278 278 278 279 279

8.7

281

REFERENCES

8.1

INTRODUCTION

Almost any polymer can be converted into a fiber if its molecular weight is above some critical minimum value (e.g. somewhat above the chain entanglement point), if it melts before decomposition occurs or if it is soluble in some solvent which does not degrade it too rapidly. In fact many thousands of polymers have been evaluated as fibers in the laboratories of the the world but only a few have succeeded in becoming of commercial importance based on their balance of properties and cost. In apparel end uses, properties sought (e.g. modulus) are in some respects similar to those of natural fibers (cotton, wool, silk and flax) but with superior strength, durability, processibility and end use performance. To achieve this kind of performance requires not only proper selection of the polymers but the engineering techniques to develol? the potential inherent in them. The major PS 7-1*

249

250

Generic Polymer Systems and Applications

synthetic polymers developed for apparel fibers include nylon 6,6 and nylon 6, poly(ethylene terephthalate) (PET) and polyacrylonitrile (PAN). A higher performance polyamide fiber based on bis(4-aminocyclohexyl)methane and dodecanedioic acid (PACM-12), which yielded fabrics with excellent aesthetic qualities and wash and wear wrinkle performance, was commercialized by Du Pont but was eventually withdrawn because of cost considerations. This exemplifies the major importance of cost as a polymer and fiber property. In industrial end uses, strength, modulus, dimensional stability, fatigue resistance and durability in exposure to heat and light have been among the properties frequently used as criteria for acceptance (along with cost). Thus nylon 6,6, nylon 6 and PET were originally selected for such end uses as tire reinforcement, industrial fabrics, ropes and cables and ballistic fabrics. Subsequently, technology has developed for preparation of high performance fibers with higher levels of strength, modulus and flame resistance. For example, the development of technology for polymerizing aromatic polyamides (aramids) and spinning them from solution, facilitated by liquid crystal formation in some cases, has resulted in property levels two to ten times those of the incumbent fibers. This has, in turn, spurred new efforts to achieve even higher property levels by capitalizing on the new principles discovered in these developments with new polymeric structures which are even stiffer than the aramids or which form liquid crystalline melts. In this chapter, some important aspects of the polymerization of some of the major polymers used in fibers will be described, including control requirements for molecular weight and end groups. Key polymer chemical and structural factors and how they relate theoretically and practically to fiber properties will also be described, and it will be shown what is involved in engineering some of these polymers into the desired fibers (e.g. spinning and drawing). At this point, some of the units commonly encountered in the description of polymer and fiber properties will be briefly discussed, conforming in general with the SI system. The linear density of a fiber is the mass per unit length and the SI unit is the tex which is the mass in grams of 1000 m of yarn. The older unit, which is still used in many places, is the denier which is the mass in grams of 9000 m of yarn. The unit of force is the Newton (N), which is 10 5 dyn, while the unit of stress is the Pascal (Pa), which is N m - 2, so that 1 Pa = 10 dyn cm - 2. Viscosity is given in Pa s so that 10 poise = 1 Pa s. Specific stress is given in N tex -1 or dN tex -1, cf the older unit of g denier- 1 (1 g denier- 1 = 0.8826 dN tex -1). To relate specific stress to stress requires the introduction of fiber density so that to convert from GPa (10 9 Pa) to dNtex- 1 , one must multiply by ten and divide by the density in gcm- 3 •

8.2· POLYMERIZATION AND ITS CONTROL 8.2.1

Polyamides

The two major polyamides used for synthetic fibers are nylon 6,6 and nylon 6. Aspects of the polymerization and control of the former will be discussed first.

8.2.1.1

Nylon 6,6

The intermediates used for preparation of nylon 6,6, copoly(hexamethylenediamine/adipic acid), are hexamethylenediamine and adipic acid. 1 The repeat unit of the polymer is -NH(CH2)6NHCO(CH2)4CO-. The diamine, which melts at 40.87°C, is normally used in the form of a concentrated aqueous solution. The dibasic acid is used in its pure solid form (m.p. = 152.1 °C). In the preparation of the polymer, the first step is preparation of a salt from precisely stoichiometric quantities of the intermediates. While the aqueous salt solution is prepared so that it has a concentration of about 500/0, the equivalence is more readily determined by measurement of the pH of a 90/0 salt solution prepared by dilution of the main salt batch. The pH at equivalence is about 7.6 for this concentration and a shift in pH of only 0.1 results in a change in end group balance of about four equiv. 10- 6 g. As will be shown below, this small change in pH can result in unsuitable changes in the concentration of amine (-NH 2) end groups in the polymer and fiber for those end uses where a high degree of uniformity of dyeing with acid dyes is required. The extent to which such changes can affect molecular weight will be described below. In carrying out a typical polymerization of nylon 6,6, the salt solution is subjected to evaporation at the boil, possibly at elevated pressures, until concentrations 2 60% are achieved. The concentrated salt solution is then heated in a reactor so that temperature increases gradually and pressurereaches, typically, 1.73 MPa (250 psi). As water evaporates and temperature increases from about

Polymers for Fibers

251

212°C to 275 °c, the molecular weight of the polymer reaches about 4400. Further reaction is achieved by a gradual decrease in pressure to atmospheric and then holding of the polymer under these conditions for about one hour. At this point the polymer is not quite equilibrated but molecular weights are in the range of 15000 to 17000. The process has been devised to remove all of the liquid water present in the salt solution as well as almost all of the potential water of reaction present in the form of carboxyl (-C0 2 H) and amine (-NH 2 ) end groups with only minimal loss of the diamine (HMD) whose atmospheric pressure boiling point is 200°C. The finished polymer is then extruded in the form of a ribbon or strand, quenched with water and cut to form chips suitable, after drying, for remelting in a subsequent spinning operation. Alternatively the polymer may be conducted directly to a spinning machine without prior solidification. If higher molecular weights are needed, the molten polymer may be heated under vacuum for some additional period of time or else polymer chips can be heated below their melting point under nitrogen. With the latter process, it is more difficult to control molecular weight precisely. The polyamidation reaction is represented by equation (1). The equilibrium constant is shown in equations (2a) and (2b) where A is the concentration of amide groups (-CONH-) and P is the product of functional end groups (-C0 2 H x -NH 2 ). For typically high extents of reaction and molecular weights, the value of the equilibrium constant (K a ) at 280°C is about 300 ± 50. 2 ,3,4 The uncertainty arises because of some difference of opinion concerning the equilibrium water content of a nylon melt under 1 atm of steam. For one of the quoted values 1 (0.160/0 or 89 mol 10- 6 g) and the experimental value for equilibrium end group product of 3000 equiv. 2 10- 12 g, K a is calculated to be about 260 when the amide concentration of about 8800 equiv. 10- 6 g is inserted in equation (2a). Since the amide group concentration at high conversions is almost constant as molecular weight varies, and since the water concentration in the melt at a given temperature depends only on the water vapor pressure, the equilibrium value of P is proportional to water pressure. Deviations from this first power relationship are observed at high steam pressures and low molecular weights. -C0 2H Ka

=

+

-NH 2

~

-CONH-

+

H 20

[ -CONH-] [H 20]j[-C0 2H] [ -NH 2] K a = A[H 2 0]jP

(1) (2a) (2b)

Amidation at high conversions is an exothermic reaction with a heat of reaction believed to be about 25-29 kJ mol- 1 (6-7 kcal mol- 1). 2 However, the heat of vaporization of H 2 0 from nylon, based on equilibrium regain data 5 for nylon 6,6 extrapolated to polymerization temperatures, is in the same range (endothermic). As a result, the equilibrium product of ends for some given water vapor pressure does not change to any great extent as temperature is varied since the reduction in the value of K a with increasing temperature is almost exactly offset by the reduced H 2 0 content at a given steam pressure. The polyamidation reaction follows second order kinetics at conversions up to about 90% and is not accelerated by catalysts. 6 ,7, 8 However, at the higher conversions where molecular weights become of practical interest, the reaction becomes third order and it is catalyzed by -C0 2 H end groups. In the presence of H 2 0, the rate equation is given by equation (3) where k a is the amidation rate constant and kh is that for hydrolysis. At the later stages of polymerization, the amide concentration is almost constant. If the polymerization is carried out at some fixed water vapor pressure, khA [H 2 0] is constant and equal to kaPeq where Peq is [-C0 2 H] [-NH 2 ] at equilibrium. The rate equation is then given by equation (4). If the end groups are not equal in concentration, but if the difference in ends, [-C0 2 H] - [-NH 2 ] = D, is constant, the rate equations can be readily integrated. 2 This may not be the case for nylon 6,6 where degradation reactions can cause a slow decrease in the value of D with time. While there is no unanimity in reports of the activation energy for the polyamidation reaction, a reasonable estimate 9 for nylon 6,6 is about 88 kJ mol- 1 (21 kcal mol-I). Thus in the normal temperature range for polymerization the rate increases by about 40% for a 10°C temperature rise. This activation energy is significantly higher than that for diffusion, as determined from melt viscosity measurements (see below). Therefore the rate is not affected by stirring. Furthermore, because of the highly favorable equilibrium constant, the rate under an atmosphere of steam is not affected by the rate of removal of water bubbles up to a molecular weight of about 18000. Since the reaction is catalyzed by -C0 2 H at high conversions, it is not surprising that other catalysts have been found. Among these are hypophosphite salts 10 and phosphonic acids. 11 -d[-C0 2H]jdt -d[-C0 2H]jdt

ka [-C0 2H]2[-NH 2] =

ka [-C0 2H](P

-

Peq )

kh [-C0 2H] A [H 20]

(3) (4)

252

Generic Polymer Systems and Applications

The molecular weight distribution of linear nylon 6,6 follows the most probable distribution as derived theoretically by Flory.12 The distribution is shown in equation (5) where W x is the weight fraction of polymer chains having x monomer units and and p is the extent of reaction or the fraction of the original end groups which have reacted. For this distribution, the weight average degree of polymerization (Pw ) is 1+ P times the number degree of polymerization (Pn ) and the latter is equal to (1-p)-1. The corresponding average molecular weights are calculated by multiplication of the respective value of P by Wn the molecular weight of the repeat unit. For an AA-BB polymer such as nylon 6,6 [i.e. one based on a diamine (AA) and a dicarboxylic acid (BB)], one may use an average value of p for the two kinds of end groups if the ends are not too badly unbalanced. In this case an average value of W r is used (i.e. 113 g for nylon 6,6). This is the same value as is used for nylon 6, an A-B polyamide. Since p is usually 0.99 or greater for cases of practical interest the ratio of the two molecular weight averages is about two. In this distribution, the weight fraction of polymer is a maximum at about Mn • In a typical example, p=0.993, Pn =143, Mn =16140 and Mw =32170. (5)

The polymer contains -NH 2 and -C0 2H end groups. In addition, there may be stabilized or non-functional end groups which are introduced deliberately or produced by degradation; such end groups will be designated as E. An example is acetic acid which is sometimes added at less than 1 mol % to control molecular weight. The total end group concentration T = [-C0 2H]+ [-NH 2] + [E]. For the case of the linear polymer, each chain has two ends so that the concentration of chains is T/2. If end group concentrations are expressed in equivalents (10- 6 g), M n = 2 x 106 /T. Thus for a polymer where T= 124, M n = 16140. For linear polyamides, viscosity measurements of dilute or moderately concentrated solutions can be used to measure the molecular weight once a calibration curve is obtained. M n can be determined from end group titrations,2 if the concentration of stabilized ends is known, or from osmometry. Viscosity measurements are related more closely to M w • However, if the ratio of the two averages is always close to two, as is generally the case for linear polyamides, viscosity will also relate to Mn' Since viscosity is related to molecular weight by the Mark-Houwink equation 11 = KM rx , the calibration will generally be a log-log plot whose slope is lI.. One common solution viscosity measurement used for characterization is inherent viscosity (l1inh), typically at a concentration of 0.5 g of polymer in 100 ml of solvent, where l1inh = In I1r/c and I1r is the relative viscosity of the solution compared to the solvent (e.g. m-cresol). A typical value of l1inh for nylon 6,6 is about one for Mn of about 15000. If inherent viscosity is determined at several concentrations and extrapolated to zero concentration, the intercept is called the intrinsic viscosity or limiting viscosity number and designated with the symbol ['1]. For this kind of solution viscosity measurement, lI. is about 0.72. 13 Another method commonly employed is to measure the relative viscosity (RV) of an 8.4% solution of polymer in 900/0 formic acid. In this case, the log-log plots of viscosity vs. molecular weight may change slope abruptly when the molecular weight reaches the chain entanglement point for this concentration. Beyond this point, the slope will be about 3.4 for the linear polymer. Typical values of relative viscosity for nylon 6,6 are in the range of 30 to 70 where the lower ranges are used for textile yarns and the higher fqr industrial yarns. For example, an RV of 41 corresponds to M n of about 15 000· while an RV of 60 corresponds to about 19000. If trifunctional units are introduced into the polymer deliberately or as a result of degradation reactions, the polymer is branched and the molecular weight distribution is broadened 14 , 15 so that M w / M n becomes greater than two. Thus solution viscosity will be higher for a given Mn • The effects on the molecular weight averages are shown in equations (6a) and (6b) where b is the concentration of trifunctional units. 2 However, since the radius of gyration of a branched polymer chain is smaller than that of a linear chain of the same weight, the viscosity of a branched polymer having a given Mw will be lower than that of the corresponding linear polymer. For a relatively flexible polymer such as nylon 6,6, the solution viscosity will depend 16 on the appropriate function of (1- 3b/T)°·5 rather than 1- 3b/T or Mw • Thus, for example, a polymer with 100 ends (10- 6 g) and ten branches would have a solution viscosity equivalent to a linear polymer with about 84 ends (10- 6 g). Mw Mn

=

4

x

10 6 1T(1

3b1T)

(6a)

2

X

10 6 1T(1

biT)

(6b)

The viscosity of molten linear nylon 6,6 also depends on Mw but since, in all cases of practical interest, operation is at molecular weights well above the entanglement point, the value of lI. in the Mark-Houwink equation is about 3.4. This means that a variation of ± 1% in M w will result in

Polymers for Fibers

253

about a 3.4% variation in viscosity. The activation energy of melt viscosity17 is estimated at about 59 kJ mol- 1 (14 kcal mol-I) for temperatures well above the polymer melting point of 265°C (e.g. ~ 280°C) so that the viscosity decreases by about 250/0 for a temperature increase of 10°C. In making such determinations, it is important to take into account any degradation which may occur and normalize for the differences in molecular weight, if any. At shear stresses above about 30 kPa (3 x lOs dyn cm - 2), melt viscosity begins to decrease with increasing shear stress and the polymer begins to show significant elastic behavior (e.g. as evidenced by a bulge at the exit of a spinneret capillary). A typical Newtonian (i.e. low shear rate) melt viscosity at 280°C for a polymer with Mn of 18000 is about 190 Pa s (1900 poise). If it is assumed that the controllable variables in polyamidation are the product of ends (P), the difference in ends ([-C0 2H] - [-NH 2] = D) and the stabilized ends (E), then the -NH 2 and -C0 2H end group concentrations are given 2, 17 by equations (7a) and (7b). The sum of these functional ends (S) is given by equation (8) while the total end group concentration (T) = S + E and Mn = 2 x 106 / T. From these relationships, for example, the degree of control in D and/or P needed to control the -NH 2 level to ± 1 or M n to ± 1% (and melt viscosity to ± 3.40/0) may be deduced. Alternatively, to achieve a certain Un along with some specified level of -NH 2, the required values of D, P and possibly E can be derived so that process conditions can be specified. 2[-NH 2 J

- D

=

2[-C0 2 H] S

8.2.1.2

=

[-C0 2 H]

D

+

+ +

[-NH 2 ]

(D 2 (D 2

=

+ +

4P)O.5

(7a)

4P)O.5

(7b)

(D 2

+

4P)O.5

(8)

Nylon 6

Nylon 6 can be considered to be the condensation polymer of e-aminocaproic acid and is thus an example of an A-B polyamide. Its repeat unit is -NH(CH 2)sCO-. In fact, this polyamide is almost always prepared from e-caprolactam in a process which is essentially an addition polymerization. The advantages of this monomer are low cost and relative ease of purification as compared with the amino acid. The pure monomer, which melts at about 69°C, does not polymerize when heated at elevated temperatures in the dry state. The key discovery18 that in the presence of both amine and carboxyl groups ring opening occurs readily, made possible the commercial processes for the polymerization of caprolactam. One way to accomplish this is by addition of e-aminocaproic acid. Another is to add a few percent of a nylon 6,6 salt. However, the latter results in production of a random copolymer with reduced melting point and crystallinity, which may not always be desirable. The simplest method is to carry out the polymerization in the presence of water which hydrolyzes some of the lactam to form -C0 2H and -NH 2 groups which then catalyze the addition polymerization reaction. This can be accomplished at atmospheric pressure, but the times required to achieve equilibrium monomer content can be greatly reduced by operation at elevated pressures and increased water concentrations. A major distinguishing feature of nylon 6 polymerization is the presence at equilibrium of a significant concentration of the cyclic monomer along with lesser amounts of cyclic dimer and higher oligomers. For example,t9 the equilibrium concentration of caprolactam at 250°C is about 7.80/0 and of the dimer is 1.13%. As the polymerization temperature is increased, the equilibrium monomer content increases moderately in accord with the exothermic heat of ring opening of 17-21 kJ mol- 1 (4-5 kcal mol-I). The reported equilibrium constants for the end group-amide equilibrium19,20,21 in nylon 6 range from about the same to somewhat larger than those for nylon 6,6 (e.g. 428 at 280 °C I9 ). However, the heat of amidation from these reports is about the same. It would appear that the presence of caprolactam tends to increase water content in the melt since the end group concentrations at fixed water vapor pressure decrease during the course of caprolactam polymerization as monomer concentration decreases towards equilibrium (see below). The ring-opening polymerization of caprolactam proceeds by a -C0 2H catalyzed addition of an -NH 2 group to the lactam ring. 22 ,23 The end group concentrations are in turn a function of the rates of hydrolysis of caprolactam and cyclization of aminocaproic acid and of the same reactions in the polymer. In the course of the polymerization under a steam atmosphere, there is initially an induction period as end group concentrations build up. As they do, the rate of polymerization increases and the end group concentrations reach a maximum in the region of the inflection point in the curve of monomer content vs. time.20,24 Some of the kinetic treatments assume both an

254

Generic Polymer Systems and Applications

uncatalyzed and catalyzed reaction19,21 and attempt to deal with all the reactions in detail. The activation energy reported for the polymerization 19 is about 78kJmol- 1 (18.7 kcalmol- 1) for the -C0 2H catalyzed reaction and 88 kJ mol- 1 (21 kcal mol- 1) for the uncatalyzed reaction, but the pre-exponential factor for the former is 24 times larger. Thus except at very small end group concentrations, the third order reaction will predominate so that the fractional rate of monomer conversion 2 will be approximately proportional to [-C0 2H] [-NH 2]. A review of advances in polymerization engineering of nylon 6 can be found in refs. 25 and 26. Molecular weights of nylon 6 are generally in the same range as those of nylon 6,6. For nylon 6, a commonly used viscosity measurement for molecular weight involves use of a solution of 1 g of polymer in 100 cm3 of 96% H 2 S0 4, A value of 2.7 has been quoted 27 for polymer with M n = 20000 and the value of rx in the Mark-Houwink equation was 0.7. For melt viscosity, rx was 3.5 and its low shear rate value was about 140 Pa s (1400 poise) at 280°C. The activation energy of melt viscosity was about 60 kJ mol- 1 (14.3 kcal mol- 1), about the same as for nylon 6,6. The onset of nonNewtonian behavior occurred at a shear stress of about 30 kPa. For nylon 6 prepared as described above, the molecular weight distribution of the final polymer is the same as for linear nylon 6,6 or other typical condensation polymers with M wi M n = 1 + P or about two. However for this AB polymer, in contrast to the AA-BB polymers, it is possible to obtain a narrower molecular weight distribution by addition of a bifunctional 'stabilizer' such as a dicarboxylic acid. z8 As the lesser end group (e.g. the -NH z group) is decreased toward zero concentration (e.g. by continuing polymerization under vacuum), Mwl M n approaches 1.5. The factor which makes the AB polymers unique in this regard is the fact that there can only be one BB unit in a chain since that chain will have B end groups on both sides. For a more practical situation where the A (-NH 2) groups have not decreased to zero, the situation is described by equation (9).2,17 Thus, for example, if 40 equiv. 10- 6 g of BB units are added and the concentration of A ends is reduced to 10 equiv. 10- 6 g, Mw/M n = 1.60. In such a case, it should be possible to increase Mn by about 250/0 for a given melt viscosity. If monofunctional impurities are present, this will have the effect of moving the polydispersity back towards the normal value. M w /2M n

8.2.2

=

[A

+

3BB] [A

+

BB]/[A

+

2BB]2

(9)

Poly(ethylene terephthalate)

The classical route for the polymerization of poly(ethylene terephthalate) (PET) -OCH2CH202CC6H4CO- involves the use of ethylene glycol (EG) HOCH 2 CH 2 0H and dimethyl terephthalate (DMT) Me0 2CC 6H 4CO zMe as starting intermediates. These are reacted in an 'exchanger' in the presence of a catalyst to eliminate methanol and form bis(2-hydroxyethyl) terephthalate (BHET), which is referred to as monomer. 29 If the mole ratio of EG to DMT is two or a little larger,high conversion to BHET is achieved. If, for economic reasons, the mole ratio is less than two, the 'monomer' contains some higher 0ligomers. 3o Typical effective catalysts for this reaction include manganese(II) acetate [Mn(OAc)2], Zn(OAc)2, zinc formate and Pb(OAc)z. The reactivity of the two methyl ester groups appears to be equa1 30 ,31 so that the reaction can be represented as a third order reaction of EG, methyl ester groups and catalyst, or a pseudo-second order reaction at constant catalyst concentration. Temperatures used range from about 150°C to over 200 °C as the reaction proceeds. The reported activation energy of this reaction varies somewhat with the catalyst and the author. Examples are 63 kJmol- 1 (15 kcalmol- 1),32 36 kJ mol- 1 (8.6 kcal mol- 1)33 and 52 kJ mol- 1 (12.4 kcal mol- 1)34 for zinc acetate catalyst and 52.3 kJ mol- 1 (12.5 kcal mol- 1)32 for lead oxide. Monomer can be synthesized by direct reaction of EG with terephthalic (T) acid with lower ratios ofEG to acid than required in the synthesis from DMT and without the use of special catalysts. This is made possible by the availability of high purity T acid. However, relatively high temperatures (e.g. 275°C) are needed to initiate the reaction which is catalyzed by -C0 2H groups. Alternatively, the intermediates are added to low molecular weight polymer at temperatures above 300°C. 29 The activation energy of this reaction has been reported to be 101 kJ mol- 1 (24.2 kcal Il101- 1)35 or 71.6 kJ mol- 1 (17.1 kcal mol- 1).36 The reaction which converts BHET to polymer involves a catalyzed interchange reaction of -OH end groups with terminal ester groups to form an ester link and EG which must then be removed. Since the reactants and products are chemically the same, it is not surprising that the heat of reaction is about zer0 35 ,37 so that the equilibrium constant does not change with temperature. For the same reason, the equilibrium constant (defined in equation 10) is near unity. In

Polymers for Fibers

255

equation (10), E represents the concentration of ester links. In one report,37 K e increased with increasing molecular weight because of abnormally high monomer content compared to that predicted from the most probable distribution. This indicates a lower reactivity for the -OH end of the monomer. Thus for about 400/0 end group conversion it was about 0.4-0.5 while for 84% conversion it was 1.1. Other reports are 0.55-0.59 29 and 0.50 32 and do not necessarily agree with the hypothesis of unequal reactivity for the monomer. If one considers the net heat of reaction, which includes the heat of vaporization (~Hv) of EG from the polymer (perhaps 59-71 kJmol- 1 or 14-17 kcal mol- 1), the reaction is endothermic. At a given vapor pressure of EG, the equilibrium -OH end concentration should decrease as the one half power of EG in the polymer melt, whether due to pressure or temperature changes. Since typical values of M n for PET are in the range 20000 to 40000, -OH end group concentrations range from::; 100 to ::;50 equiv. 10- 6 g, depending on the amount of other end groups (e.g. -C0 2H) present as a result of degradation or incomplete esterification in the direct process. These can range from 10-30 equiv. 10- 6 g or higher, depending on the thermal history of the polymer. For example, if we let the ester concentration be 10400 with -OH eq =70 and assume that K e =0.5, the EG concentration is calculated to be 0.118 (all concentrations in mol 10- 6 g). Typical pressures in finishing PET polymer are in the range 0.5-3 mmHg. In the absence of end groups other than -OH, and since (-OH)eq varies as the square root of pressure, if one wishes to maintain equilibrium molecular weight at ± 1.5% (to control melt viscosity to ± 5%), it is necessary to control pressure to ± 30/0 or ±0.06 mm for a 2 mm average pressure. The requirements are somewhat less stringent if other ends are present. Ke

=

2(E)(EG)/[-OH]2

(10)

Examples of catalysts for the polymerization of PET are Sb 20 3 and organic titanate esters. The use of the latter tends to produce a yellow color and this can limit their usefulness for apparel fibers. In spite of the fact that reaction rate depends on catalyst concentration, it also depends strongly on the rate of removal of ethylene glycol and hence on stirring rate, rate of surface generation and polymer film thickness.29,38-41 This results from the relatively low equilibrium constant (0.5-1) and contrasts with the situation for the aliphatic polyamides. Thus the time required to reach some goal molecular weight can decrease by about two orders of magnitude as film thickness decreases from 0.1 to 0.0025 mm. As a result, most values of the activation energy of polymerization must be considered as apparent values since both diffusion and chemical reaction enter into the process. One value quoted 40 for the activation energy for the Sb 20 3-catalyzed polymerization is 59 kJ mol- 1 (14 kcal mol- 1), while for the uncatalyzed reaction, values reported range from 96 42 to 188 40 kJ mol- 1 (23 to 45 kcal mol- 1). A study of the direct esterification process 43 involving computer simulation reports 93 kJ mol- 1 for the reaction of a -C0 2H end with a -OH end to form water and an ester group but only about 12 kJ mol- 1 for the ester exchange reaction to produce EG. From the high polymerization rates observed in thin film experiments, it is not surprising that the rate of interchange between a mixture of two compatible polyesters is also quite high. For example, in two studies of mixtures of normal and deuterated PET,44,45 both containing Sb 20 3, it was found that interchange was quite rapid and that 45 the block length was reduced to half its initial value after only 10 s at 280°C and to only' 4% of the initial value in 3 min. This reaction also occurred in the solid state. In one study,45 the activation energy of this reaction was surprisingly high (152 kJ mol- 1 or 36 kcal mol- 1) for such a fast reaction. A relatively high value (130 kJ mol- 1) was also found in an earlier study46 by observing the rate of molecular weight redistribution. However, in a recent study47 wherein a low molecular weight ester or an alcohol was mixed with PET at temperatures below its melting point and the molecular weight monitored, activation energies reported were only 45 and 91 kJ mol ~ 1 for ester-ester exchange and alcohol-ester exchange respectively. The molecular weight of PET is conveniently tracked by solution viscosity. For example, in one report,48 the Mark-Houwink relationship for intrinsic viscosity [11], determined in trifluoroacetic acid against viscosity average molecular weight, is [11] =4.33 x 10- 4 M~·68 so that [11] =0.56 corresponds to Mv = 37 400. In most cases, the molecular weight distribution for high molecular weight PET is normal (i.e. most probable) so that solution viscosity can be related to M n as well. For this distribution, Mv / Mn = 1.87. Thus, from another relationship49 for [11] determined at 25°C in 50:50 1,1,2,2-tetrachloroethane:phenol, [11]=2.1 x 10- 4 M~·82. With this set of parameters; Mn = 20000 corresponds to [11]=0.706 which· is not in good agreement with the other set given above. It is not clear to what extent the differences in the two solvents are responsible. Still another set of data 50 is consistent with the relationship [11] = 6.762 x 10- 4 M~·69 for [11] measured in 50: 50 trifluoroacetic acid: dichloromethane at 30°C. This relationship predicts [11] = 0.63 for Mn = 20000

256

Generic Polymer Systems and Applications

and is closer to the first set. In a comparison of several solvents,51 the values of r:x in the Mark-Houwink relationship (vs. My) ranged from 0.648 for phenol/tetrachloroethane (PCTE) to 0.723 in pentafluorophenol with corresponding values of K decreasing from 7.44 x 10- 4 to 3.85 X 10- 4. For My = 37400, the corresponding values of [11] for these solvents are 0.68 and 0.78. Finally, in a study of fractionated PET, r:x for PCTE was 0.73 and [11] for the polymer with Mw = 40 000 was about 0.7. It would be useful to have these discrepancies resolved. As expected for a linear polymer above the critical molecular weight for entanglement, the melt viscosity of PET increases as the molecular weight increases by approximately the power of 3.5. In one study,53 the flow was found to be Newtonian up to a shear stress of 9.65 x 104 Pa and have a power law exponent of about 0.73 between that stress and 4.14 x 10 5 Pa. The activation energy of viscosity was 56.5 kJ mol- 1 (13.5 kcal mol- 1). In another study,50 the activation energy reported was about the same but the power law exponent was 0.86 at 285°C and non-Newtonian behavior was already evident at shear stresses of about 4 x 104 Pa. In the latter study, polymer with [11] = 0.72 (Mn = 24400) had a melt viscosity at 290°C of 460 Pa s (4600 poise), while polymer with [11] = 0.96 (Mn = 37000) had a viscosity of 1950 Pa s (19500 poise). PET polymer and fibers contain a small concentration of cyclic oligomers, a major fraction of which is cyclic trimer. In one report,54 the total oligomer content was 1.920/0 of the fiber weight and 56% was trimer. There is a tendency for trimer to diffuse to the surface of the fibers so that> 97% of the surface oligomers are trimer. These can sometimes cause problems in fiber processing.

8.2.3 Polymers Used in High Performance Fibers Since the early 1970s there has been increasing commercial and experimental interest in fibers having tensile strengths at least twice as high as nylon and polyester industrial yarns and moduli five to twenty times higher. Most of these fibers involve the use of ring-containing polymers which are rod-like in character and which form liquid crystalline solutions or melts. A notable exception to this trend is ultra-high molecular weight highly oriented polyethylene fibers which are prepared via a 'gel-spinning' process and which get their very high modulus fronl very high crystallinity and orientation rather than from chain stiffness, and their high strength from the use of very high draw ratios. Carbon fibers formed from polyacrylonitrile (PAN) are another exception. Some of the procedures used to prepare several of the stiff chain polymers which have been of recent interest and their properties are described below. 8.2.3.1

Poly( 1 ,4-phenyleneterephthalamide)

Poly(1,4-phenyleneterephthalamide) or PPD-T (-NHC 6H 4NHCOC 6H 4CO-) is the aramid on which the high strength, high modulus fiber Kevlar® is based. In general, such aramids are prepared by reaction of an aromatic acid dichloride with the diamine in an amide solvent, possibly containing a halide salt as LiCI or CaCI 2 • In one procedure, 55solid terephthaloyl chloride (TCI) was added to a cold solution of p-phenylenediamine (PPD) in the amide solvent which consisted of a mixture of hexamethylphosphoram'ide (HMPA) and N-methylpyrrolid-2-one (NMP) in a 2/1 ratio. The stirred reaction mixture became thick and gel-like in about 5 min and the polymerization was nearly complete at this point. The maximum 11inh (0.5 g polymer in 960/0 H 2 S0 4) was obtained at a concentration of each reactant of about 0.25--0.3 moll- 1 (e.g. 6.9 11inh). Following the polymerization, the polymer was washed to remove the solvent and HCI produced in the reaction and then dried. Since the aromatic diamine is a weak base, the solvent (present in excess) preferentially bonds the HCI to prevent significant protonation of the amine end group which would inhibit the polymerization. HMPA must be handled with care to avoid inhalation and skin contact since it has been found to cause cancer in laboratory animals. Another solvent system disclosed 56 is NMP containing CaCl z . A cold (- 20°C) solution containing 5.51 parts of PPD, 7.07 parts of CaCl 2 and 87.42 parts of dry NMP containing only 143 p.p.m. of H 2 0 was combined with a stoichiometric amount of TCI in a mixer and then in two twin-screw mixers in sequence. The highly exothermic reaction of the amine and acid chloride groups raised the prepolymer temperature to about 50°C. After about 15 min in the second screw mixer, a polymeric 'crumb' containing 11 % PPD-T was obtained with 11inh of about six. For rigid rods, the value of r:x in the Mark-Houwink equation should be 1.7-2. The fact that r:x is significantly less than two for PPD-T in 960/0 H 2 S0 4 indicates a polymer chain with a finite persistence length. However there is not universal agreement on the exact solution viscosity

Polymers for Fibers

257

parameters for PPD-T. In one report, 57 r:J. was given as about 1.07 with [17] = 7.6 corresponding to

M w = 40 000 and [17] = 8.2 corresponding to M w = 43000. In another,58 r:J. was 1.09 with 17inh of 4.4 corresponding to M w = 24 300, 5.12 to 35000 and 6.0 ([17] about 8.2) to 43500. This same report

concluded that Mw/ Mn was about two for 17inh of 2-3 but increased to 3.2 for 17inh of 6.0 ([17] = 8.2). It was also concluded that the persistence length in this solvent was about l5 nm. However, other reports 57 ,59-61 concluded that the persistence length was 24,41, 45 and 130 nm respectively so that the exact rod length for PPD-T in H 2S0 4 is somewhat in doubt. In another study,62 r:J. was reported to be 1.36 for 96% H 2S0 4 and polymer with [17] = 5.9 (17inh estimated at 4.6) had a value of Mw = 40 100, higher than in refs. 57 and 58 (about 33600), and [17] = 8.2 a value of 51 000. The average value of Mw from these reports for [17]=8.2 (17inh=6.0) is about 46000. Some studies 63 ,64 of the molecular weight distribution of PPD-T have given values of M w/ Mn ranging from about 1.8 to 2. Another report has given 1.8 to 3.0. 65 From a theoretical standpoint, the molecular weight distribution of a uniform condensation polymer should be the 'most probable', even in the absence of opportunities for equilibration.2,66 In the event of non-uniformity the distribution will be broader but it takes a fairly large degree of non-uniformity to get M w / M n as large as three. If we imagine that some (e.g. 20%) of the polymerizing PPD-T spends enough time more than the average polymer so that its molecular weight is twice as large as the main polymer, then Mw/M n =2.16. This can be calculated 2 by averaging Mwand the reciprocal of Mn. In another example, if 50/0 of the polymer has a very low molecular weight (e.g. 20% of the main polymer), then M w/ M n = 2.3. If we simply assume that the distribution is the 'most probable', then M n for polymer with [17] = 8.2 would be about 23000 based on the average value for Mw given above. However, this must be considered an estimate. In the spinning of PPD-T, concentrated solutions (e.g. 16-19%) in ca. 1000/0 H 2S0 4 are used. Because of the rod-like character of the polymer, there is a critical concentration for phase separation to form liquid crystals. The theory developed by Flory67 leads to the relationship shown in equation (11), where 4>c is the critical volume fraction of the polymer having chains with an aspect ratio (rod length/diameter) of R a . For PPD-T, 55 the critical concentration appears to increase linearly with the reciprocal of 17inh and is about 80/0 (by weight) for 517inh and about 100/0 for 1.817inh. In general, the viscosity of solutions of rigid rod polymers increases rapidly with concentration (c) in the region below the critical concentration for liquid crystal formation. In one report 68 for solutions of PPD-Tin 100% H 2S0 4 at 24°C, zero shear rate viscosity above the entanglement region (but below 4>c) was found to increase as (CM w )6.8. Thus for this relationship it might be expected that in going from the critical concentration (e.g. 8%) to (say) 19%, the viscosity, in the absence of liquid crystal formation, would increase by a factor of about 360. Even if the rods were not quite so stiff, so that viscosity increased as the fourth power of concentration (for example), the viscosity would increase by a factor of 32. In either case, processibility would become very difficult or impossible. Fortunately, with the onset of anisotropy or liquid crystal formation, viscosity starts to decrease69, 70 so that at the concentrations used for spinning it is in a conventional range typical of melt spinning, or lower. The implication is that the degree of anisotropy continues to increase as long as viscosity decreases with increasing concentration. In one rheological study71 with PPD-T obtained by dissolving Kevlar® fibers in H 2S0 4 at 60°C, viscosity of solutions up to about 8% were found to be Newtonian in the shear rate range of 10- 2 to 10s- 1 (i.e. up to about 10 3 Pa shear stress). On the other hand, 12% solutions decreased in viscosity starting at a shear rate of about 10- 1 s-1 (shear stress of about 10 Pa). This is another characteristic of anisotropic solutions and is due, at least in part, to the alignment of the liquid crystals under shear. A similar study72 with 14% and 16% solutions at 60°C showed an onset of non-Newtonian behavior at a shear stress of about 10 3 Pa. (11)

8.2.3.2

Thermotropic polyarylates

Fully aromatic p-linked polyesters or copolyesters can be synthesized which melt at temperatures below those at which they are thermally unstable. In such cases the polymer melts are anisotropic and are liquid crystalline (nematic) and thus have many similarities to the nematic liquid crystalline solutions of PPD-T. They orient efficiently in spinning and thus do not require drawing. However, they frequently cannot be spun at the molecular weight levels required for achievement of high tensile strengths. In this case they are spun at lower molecular weights and the oriented as-spun fibers are then subjected to solid phase polymerization during which they strengthen. Typical intermediates used for these polyesters include the diacetates of hydroquinone (HQ),

258

Generic Polymer Systems and Applications

phenylhydroquinone (PhHQ), chlorohydroquinone (CIHQ), p-hydroxybiphenyl (BP) and aromatic dicarboxylic acids such as terephthalic acid (T) and 2,6-naphthalenedicarboxylic acid (N). In addition, p-hydroxybenzoic acid (PHB) and 2,6-hydroxynaphthoic acid (PHN) have been used. Homopolymers such as poly(p-hydroxybenzoic acid) or poly(p-phenylene terephthalate) are too high melting to be melt processed without degradation, so it is necessary to use random copolymers in almost all cases (the polymer PhHQ-T is a possible exception since it melts at about 344°C). For example the homopolyester HQ-T melts at about 600°C 73 but the melting point decreases at the rate of about 37.5°C per 10 mol % modification with PhHQ. Poly(3,4'-dihydroxybenzophenone terephthalate) (3,4'-DHBP-T) also forms anisotropic melts. 74 For the most part, the polymerization involves the reaction of an acetoxy group with an aromatic carboxyl group to form the desired ester linkage and eliminate acetic acid, which is distilled from the mixture. In a typical polymerization procedure 75-77 the diacetate of the aromatic diol is combined with terephthalic acid or another aromatic diacid. Alternatively, the acetate of a hydroxy acid may be used. An excess (e.g. 4-5 mol 0/0) of the diacetate may be needed to compensate for some loss during the polymerization. Sodium acetate (e.g. 0.020/0) may be used as a catalyst. After a thorough purge with an inert gas (e.g. nitrogen or argon), the reactants are heated progressively to the polymer melting temperature or higher, and acetic acid is collected. After about 2-3 h (for example), the batch is placed under vacuum and subjected to progressively lower pressure for a period of time (e.g. 15 min to 1h) and to a final pressure level (e.g. 0.5-8 mmHg) which depends on the desired molecular weight. Molecular weight of the polymer is generally assessed by measurement of llinh for a solution containing 0.5 g of polymer in 100 ml of the solvent. An example of a solvent used is the mixture hexafluoroisopropanol (HFIP; hexafluoropropan-2-01)/CHCI 3 (50/50 by volume). In some cases where the polymers tend to be less soluble, concentrations of 0.1 g/100 ml solvent are used and the solvents selected include pentafluorophenol (PFP) or phenol/p-chlorophenol/tetrachloroethane (25/40/35). For a given polymer, llinh for the more dilute solutions will be higher at the same molecular weight. There does not seem to be much in the way of reliable data for solution viscosity vs. molecular weight for many of these polymers. Typical llinh of polymers which have been conventionally melt-spun range from about 1 to 2.5. In one report,73 PhHQ-T with llinh of 2.3 had Mw of 27000 as measured by low angle laser light scattering. In this same report, melt viscosity at 380°C and 4 s- 1 shear rate was given for as-spun and heat-treated fibers and these ranged from about 10 Pa s for a fiber with 3.1 dN tex -1 (3.5 g denier- 1 ) tensile strength to 104 Pa s for a 17.7 dN tex - 1 (20 g denier - 1) heat-treated fiber, approximately a fourth power relationship. However, it is possible that these data are affected by the thermal degradation during the measurements at such a high temperature. Our own data indicate that molecular weights are considerably lower for a given llinh (0.5 g 100 ml- 1 in HFIP/CHCI 3) than the result quoted just above. Thus we believe that llinh values of two and five correspond to M n = 10000 and 22000 respectively and r.J. is about 1.14. The effects of liquid crystallinity on the rheology of thermotropic melts is qualitatively similar to those seen with anisotropic solutions of aramids, namely a lower viscosity than would otherwise be expected for the degree of chain stiffness and molecular weight and the onset of non-Newtonian behavior at lower shear stresses than for similar isotropic polymers. For example,78 copolymers of PET and p-hydroxybenzoic acid increase in melt viscosity at 275°C as PHB content is increased to 30 mol 0/0. Beyond this concentration, further addition of PHB leads to a decrease in viscosity of ~ 1 order of magnitude up to about 70 mol % PHB. This drop coincides with the onset of liquid crystallinity. Beyond about 60-70 mol 0/0, viscosity increases again and this may be a result of the presence of crystalline blocks of poly(PHB). For example, at 80 mol % the number average block length of a statistical random copolymer would be five repeat units and thus likely to be high melting. Near the minimum viscosity region (60 mol % PHB), the flow is non-Newtonian at shear stresses of about 103 Pa or lower and the power law exponent (shear stress vs. shear rate) is only about 0.25. It is possible that this low power law exponent is related to some residual crystalline structure in the melt (see below). In contrast, the same study shows that PET is Newtonian until a shear stress of about 5 x 104 Pa. Other studies with this kind of polymer79 have indicated a significant reduction in viscosity when the polymer is subjected to a higher temperature just before the rheological measurement, possibly as a result of melting of small crystallites which may be present in the melt. However, a more recent report 80 with a different sample of the same polymer composition saw no evidence of residual crystallinity or effects of preheating. These same authors had previously observed the preheating effects with another sample which they had attributed to blockiness of the copolymer. They also found that the new, possibly more random, sample showed

Polymers for Fibers

259

essentially Newtonian melt viscosity at 260°C and 280 °C at shear stresses as high as 2 x lOS Pa in contrast to previous samples which were always non-Newtonian. In still another study,81 the composition containing 60 mol % PHB was non-Newtonian at temperatures from 250°C to 285 °C but there was little temperature dependence of viscosity and this was attributed to a decrease in degree of anisotropy as temperature was increased, which tended to offset the normal activation energy factor at a constant degree of anisotropy. Another graphic example 82 of the effect of liquid crystallinity on melt viscosity is demonstrated with the polymer 3,4'-DHBP-T modified with 2.5 mol % resorcinol. This polymer has a melting temperature of about 276°C and an anisotropic-isotropic transition of about 359 °C (it is evidently not as stiff as some of the conventionally p-oriented polymers, many of which show no transition of this sort before they decompose). Polymer with Y/inh = 1 (0.1 % in PFP) had a melt viscosity at 10 2 S-1 shear rate of about 30 Pa s and a power law exponent of 0.58. At 360°C and 400 °C, viscosity increased more than an order of magnitude to about 500 Pa s but the power law exponent did not increase even when the polymer became isotropic. This was attributed to long relaxation times in such stiff chain polymers.

8.2.4

Polyacrylonitrile and Polyalkenes

Polyacrylonitrile (PAN), linear polyethylene (PE) and isotactic polypropylene (PP) are also significant polymers for the fiber industry. Their monomers are all H 2C=CHX, where X is CN, H and Me respectively for the three polymers. The principles of their polymerization are covered elsewhere in these volumes. However some key points will briefly be mentioned. They all polymerize via addition polymerization. The polymerization of PAN is a free radical addition polymerization initiated by such materials as azobis(isobutyronitrile), benzoyl peroxide and persulfates. 83 Comonomers are frequently used to enhance solubility in spinning solvents such as dimethylformamide and the dyeability of fibers. Acrylic fibers contain at least 85 % of acrylonitrile (AN), while modacrylic fibers contain less than 850/0 but more than 35% of AN. Some expressions which have been proposed for the relationship between [Y/] and molecular weight for PAN are 83 [Y/] = 2.33 x 10- 4 M~·7S and [Y/] = 2.43 x 10- 4 M~·7S so that M w= 100000 corresponds to [Y/] = 1.31. Typical fibers have [Y/] values of about 1.5. Polyethylene and polypropylene are polymerized anionically with the help of Zeigler type catalysts,84, 8S which are generally based on TiCl 3 and metal alkyls or alkyl halides. The linear polymers produced by this type of process tend to be relatively highly crystalline with densities in the range of 0.95 for PE and 0.90 for PP. The molecular weight distributions of typical polyalkenes tend to be much broader than those encountered with condensation polymers with Mw/M n values ranging from three to 15. In recent work with PE, polymers with M w as high as 4 x 106 have been spun into fibers. Based on various data in the literature,8s-87 the relationship [Y/] = 6.3 x 10- 4 M~·70 seems to fit much of the data so that [Y/] = 10 corresponds to M w= 106 . Some differences between the various sources may be a result of differences in the molecular weight distributions. A relationship proposed for polypropylene 84 for Y/inh (0.1 % in decalin at 135°C) is Y/inh = 1.00 6 4 X 10- M~·80 so that Y/inh = 6 corresponds to M w= 10 • This relationship seems to apply to data given in ref. 86 as well. Another relationship8s gives [Y/] = 1.62 x 10- 4 M~·77. Mw/M v for commercial polypropylene is about 1.3. 8s The melt viscosity of high density polyethylene at a given molecular weight is lower than that for polyamides because of the absence of hydrogen bonding. However, higher molecular weights are generally used for PE, in part because molecular weight distributions are broader. Non-Newtonian behavior begins at shear stresses of about 103 Pa or less. 89 The activation energy of viscosity90, 91 is about 28-29 kJ mol- 1 (6.8-7 kcal mol- 1), lower than that for the polyamides, again because of the absence of hydrogen bonding. While a power law relationship is generally used for shear stress vs. shear rate data in the non-Newtonian region, an alternative which has been used 90 over the entire range of shear rates is to plot log (Y//Y/o) vs. shear stres~ to obtain a straight line. Polypropylene also becomes non-Newtonian at about 103 Pa shear stre~s but the onset stress decreases and the degree of viscosity reduction increases as the molecular weight distribution broadens. 89 ,92 As expected, viscosity increases as about the 3.5 power of molecular weight (M w or M v )92 as it does for polyethylene and other linear polymers. The activation energy for viscosity of PP is about 46 kJ mol- 1 (11 kcal mol- 1).91 As an example,89 the absolute value of zero shear melt viscosity at 200°C for high density polyethylene with M w= 1.68 X lOS (Mw/M n =84) is 1.9 x lOS Pas, while that for polypropylene with M w= 4.44 X lOS (Mw/Mn = 10.4) is 5.7 x 103 Pas.

260

Generic Polymer Systems and Applications

8.3 SPINNING 8.3.1

Introduction

Fibers can be prepared from their parent polymers by melt, dry or wet spinning. In melt spinning, the molten polymer is transported to a spinning pack containing a filtration me.dium and a spinneret containing small orifices (e.g. 0.2-0.6 mm in diameter), extruded, pulled away under some tension and simultaneously cooled and forwarded via a series of rolls or godets to a windup. The cross sectional area is reduced by (for example) a factor of20 to 100. At some point in the process, spinning finishes are applied which contain lubricants and antistats. The spun fiber bundles are then subjected to a drawing operation where the orientation of the polymer chains is increased to a degree which depends on the level of tensile properties desired and possible for the particular polymer and spun fiber. This type of process is conventionally used for nylon 6,6 and nylon 6, PET and polypropylene. In dry spinning, a solution of the polymer (e.g. 20 to 30% solution of PAN with M w of 80000 to 170000 in DMF at 80 to 150 °C)S3 is filtered, deaerated and extruded through a column containing air at 230 to 260°C. As the solvent (e.g. DMF) is evaporated, the filaments solidify. Since this is a diffusion-controlled process, the skin of the filament solidifies first. As a result, upon further contraction of the filament, a non-round (dog bone) cross section is formed. The filaments leaving the column have 10-50% residual solvent which is removed in a subsequent wash draw operation. In wet spinning of PAN, a filtered, deaerated solution is extruded into a coagulant bath whose composition and coagulation rate are selected to give an optimum, smooth fiber structure. S3 For PAN filament yarns, an air gap between the spinneret and the quench bath is frequently used. The use of such an air gap (in the so-called dry jet-wet spinning process) is involved in the process for preparing high performance fibers, such as Du Pont's Kevlar® aramid fibers, from anisotropic solutions. In such cases the spun fibers are already highly oriented and require no further drawing step. Another variant of this type of process is the gel-spinning process developed for producing very high strength polyethylene fibers. S'6, 93 In this process, ultra-high molecular weight polyethylene in dilute solution (gel) is spun, extracted and drawn to very high draw ratios.

8.3.2 Threadline Tension The development of structure in the spun fibers depends greatly on the tension developed in spinning. For most fibers, spun fiber orientation will increase directly with tension, while fiber crystallinity mayor may not be strongly affected. In the case of PET, crystallinity is very low at low to moderate spinning speeds but increases rapidly as the spinning speed increases above about 3500 m min - 1. On the other hand, the crystallinity of nylon 6,6 spun fibers is fairly high at all spinning speeds. In spinning of anisotropic solutions of aramids, orientation becomes high at relatively low spinning tensions. The details of the development of tension in spinning can be found (for example) in refs. 89, 94 and 95. Some of the issues involved will be described by considering some limiting cases. The polymer or solution issues from the spinneret with a velocity Vo. If the cross sectional area of the spinneret capillary is A o and the volume flow rate is w, then Vo = wlA o . In fact this relationship holds for all positions in the threadline and w is a constant for a normal uniform threadline. This general relationship is shown in equation (12). v

=

w/A

(12)

The polymer chains just emerging from the spinneret are generally under a considerable elastic stress as a result of elongational flow in the convergence region at the entrance to the capillary and this stress is maintained to a large extent in conventional capillaries of relatively low length to diameter (LID) ratios (e.g. 1-5). On exit from the capillary, these elastic stresses are removed almost immediately and the elastically extended chains retract at a rate which depends on the relaxation time (r) of the polymer (e.g. 10- 3-10- 5 s). Because of this the velocity is reduced and a bulge appears just below the spinneret, typically at a distance of 1-2 spinneret diameters. In many cases the tension in the attenuating threadline is not sufficient to prevent the almost complete retraction of the polymer chains. In other cases (e.g. at higher spinning speeds), the bulge is considerably reduced from its free fall value. Typical area expansion ratios range from 1.5-2.5. In consideration of spinning dynamics, starting from the point ofmaximum cross sectional area A l' the end of the attenuation process is taken to be the point or region where 'solidification' occurs because of crystallization, immersion in a quench bath or approach to the glass transition temperature Tg • The velocity at this point V 2 may be somewhat less than the velocity of the feed roll since it is possible

Polymers for Fibers

261

for some solid state attenuation to take place to an extent which depends on the tension level and the modulus of the quenched fiber. This attenuation can have an important effect on the level of spun fiber orientation. The increase in velocity (V2/Vl) or the decrease in cross sectional area (A 1 /A 2 ) is referred to as the 'attenuation factor' (AF). This is larger than the 'spin stretch factor' or SSF (A o/ A 2 or V2/VO) by the degree of bulging a under the tensions which exist in the process under consideration. The stresses which contribute importantly in a typical spinning process include the following components: rheological (Ir), inertial (Ii) and aerodynamic or drag (fa). For most cases, gravitational and surface tension effects are negligible (except possibly for very slow spinning of large diameter filaments). 95 If only rheological effects are involved, the force (F = fA) along the threadline is constant but stress increases as the cross sectional area decreases. This would apply to spinning at relatively low speeds and particularly with high viscosity polymers. At higher speeds the other two components increase the threadline force. The inertial force is a function of mass flow rate and the velocity increase as shown in equation (13), where p is the. density, while the inertial stress is given by equation (14). In the cases where the attenuation factor is relatively large, the inertial stress is simply equal to pv 2 • If v is in units of em s - 1 then Ii is in dyn em - 2 (for p in g cm - 3). Thus for a speed of 600mmin- 1 (1000 ems-I) and a melt density of 1.0,1i=106 dyncm- 2 or 10 5 Pa, while at 5000 m min -1, Ii = 2.5 X 106 Pa. As will be shown below, the inertial stress becomes a dominating influence at very high spinning speeds. The aerodynamic force is proportional to V1.2 and the distance from the spinneret to the first power. 95 While this component can become quite large by the point at which the fibers reach the feed roll, it is probably not of major significance for the region above the solidification point. 95 pw(v pv(v

vd -

pvA(v 2

pv (1

Vt)

-

Vt)

(13)

Vt/v)

(14)

In discussion of the rheological contribution to threadline tension, the case of isothermal spinning where the threadline is subsequently solidified essentially instantaneously (e.g. by a quench bath or very rapid crystallization induced by orientation) is considered first. The stress developed in the attenuation process depends on the elongation rate (dv/dx), where x is the distance along the spinning path, and the elongational viscosity (1]e). The quench takes place at a point x = xc. For an incompressible polymer melt, it might be expected that 1]e = 31]0 (where 1]0 is the zero shear viscosity at the same temperature). Many of the theoretical treatments of spinning assume that elongational flow is Newtonian so that viscosity is independent of elongation rate. 94 - 97 However, some measurements 9S ,99 have indicated that YJe/YJo is much greater than three in the case of nylon 6 and may be as high as 10-15. The activation energy deduced for elongational flow 96 ,99 in several spinning studies was found to be significantly lower than that found for shear viscosity. In addition, there are a number of indications that elongational flow may not be Newtonian (see ref. 89, for example). For the Newtonian assumption, the rheological stress vs. elongation rate is given by equation (15). By use of the relation F r = irA and differentiation of equation (12), the result in equation (16) is obtained. Since, in the limit of insignificant inertial or drag forces, F is constant, integration of equation (16) between the limits x = 0 and Xc gives the results for F andk (the stress just before the quench point) shown in equations (17) and (18), where AF is again the attenuation factor. It can also be seen that In (A) decreases linearly with x or, conversely, In(v) increases linearly with x as shown in equation (19).

.£ Fr F

k In (Ao/A)

(15)

fle dv/ dx

(16)

-(flew/A)(dA/dx) =

(fleW/Xc) In (Ao/A c)

(fie/Xc) In (AF)

(17)

(fleVc/xc)ln (AF)

(18)

In (v/v o)

(19)

=

(x/xc) In (AF)

With equation (18), some of the effects can be quantified. Let it be assumed from ref. 98 that YJe of nylon 6 at 230°C is 1.5 X 104 Pa s, and that the attenuation factor AF is 10 and Xc is 50 cm. For a speed of 300 m min - 1, the rheological stress just above the quench point (fc) is about 3.5 x 10 5 Pa while the inertial stress is about 6.4% of this value. The maximum value of dv/dx is about 23 s - 1. An increase in AF from 10 to 20 would increase fc by about 30% while reducing Xc to 10 cm would increase k by a factor of five. It must be mentioned at this point that an attenuation factor of 20 in

262

Generic Polymer Systems and Applications

isothermal spinning of a Newtonian fluid can result in severe instability or 'draw resonance', as will be discussed below. The stress at the bulge maximum is Ie/AF. It can also be seen that in linear coordinates the slope of the velocity vs. distance curve (dv/dx) is a minimum near the spinneret and a maximum just above the quench point and increases linearly with velocity. If we increase the spinning speed V c but let Xc increase in proportion to the speed, thenk will not change. It is evident that introduction of cooling would increase the k by increasing the viscosity. This will be discussed further below. Considering the other extreme, where the contribution of the inertial stress to the total is much larger than that of the rheological stress, for large values of AF, equation (20) is a good approximation for the isothermal case and it can be seen that dv/dx increases as v2 • Integration of equation (20) gives the result in equations (21a-e). For fie = 1.5 X 10 2 Pa s (1.5 X 10 3 poise), density = 1.0 and 4 6 V c = 5000 m min -1 (8333 cm s -1), k is about 6.2 X 10 Pa and dv/dx is about 4.1 x 10 . From equation (21c), it is calculated that, for the example just given, an attenuation factor of ten occurs over a distance (xc - x o) of only about 1.6 cm. If the viscosity were 100 times as high, as assumed in the examples above, the distance would increase to about 160 cm. Obviously, the necking observed in high speed spinning 100 is not consistent with viscosities of this magnitude and therefore with a Newtonian model of the melt, especially when one considers the decrease in temperature which occurs in typical melt spinning. 1'fe dv / dx

pv 2

v- 1

-1

vo

Vc

-1

Vo

V- 1

Vc

(20)

-1

-1

+

pX/1'fe

(2Ia)

pX c /1'fe

(2Ib)

p(X c

-

X)/1'fe

(2Ie)

Next to be considered are the effects of cooling on the development of threadline tension. In a typical process, a stream of air is blown across the filaments descending in the spinning chimney. Under these conditions, heat is lost to the quench gas mainly by convection. The rate of cooling95 ,101 varies directly with the difference in temperature between the filament and the quench air at its surface (T - Ta). It also increases with some function of filament speed v", where n appears to be of the order of 0.3-0.5. It decreases with increasing throughput w. Taking all of these factors into account, the temperature vs. distance relationship has the shape of an exp( -x) approach to equilibrium with the slope depending mainly on wm , where m is of the order of ca. 0.5. For a given w, d(T- Ta)/dx increases only slowly as v increases by a factor of ten. 98 The rate of cooling will also vary inversely with the polymer specific heat. Examination of these relationships indicates that, over a moderate temperature range in the cooling process (e.g. to about 100°C below the spinning temperature), a linear relationship between l/T and x is obtainable when T is given in absolute temperature (K), as indicated in equation (22), where the cooling rate constant kc is mainly a function of w. When this holds (e.g. if the polymer crystallizes in this temperature range), it is frequently possible to obtain analytical solutions of the equations for threadline tension if there is an Arrhenius viscosity vs. temperature relationship (i.e. all temperatures are well above Tg ). Thus, when using equation (16), it is necessary to integrate dX/11e where l/fle = (l/fle,o) exp( -~E*kcx/R). With these assumptions, the maximum stressfc increases directly with the cooling rate and the activation energy for viscous flow as well as the other factors shown in equation (18). (22)

Next the possibility that elongational flow is non-Newtonian over at least the higher ranges of elongation rates or stress levels is considered. If this happens, it may be for the same reasons that shear flow becomes non-Newtonian, i.e. elastic extension of the polymer chains, orientation of liquid crystals or other structural anomalies of the melt. Handling of the rheological equations for spinning of a power law fluid at low speeds (i.e. insignificant inertial component) is straightforward and has been discussed in the literature. 89,102,103 In this situation, the apparent elongational viscosity is given by f/(dv/dx) so that it is preferable to call the constant relatingfto (dv/dx)" something other than viscosity. The relationf= q(dv/dx)" will be used. Using the same approach as for the isothermal Newtonian example given above, ,gives the result shown in equations (23H26), where p=n- 1 -1. Closer inspection of these equations shows that the area reduction pattern shifts to a linear one for A/A o vs. x as n decreases to 0.5 (p = 1). For still smaller values of n (larger p), area reduction becomes initially rather slow and then speeds up as x/xc approaches unity. For example, if n = 0.2 (p = 4) and Ao/A c or AF= 10, A is still about 3.2A c when x=0.99x c'

263

Polymers for Fibers -(A 1/n-2)dA/dx

Ag (Ag [(Ao/Ac)P

-

-

AP AP)/(Ag

=

F1/n/wql/n

(23)

pF1/nx/wql/n

=

-

(25)

x/xc

A~)

(A/Ac)P]/[(Ao/Ac)P

(24)

-

1]

=

(26)

For the situation where inertial stresses are much higher than the rheological stresses, the left hand side of equation (20) may be modified with the power law relation as shown above. Integration leads to an attenuation equation exactly the same as equation (26), except that for the inertial case p = (2In) -1. For example, if n = 0.25 (p = 7) and AF = 10, A is still about 3.7 Aowhen xlx o = 0.999x c ' Thus it is seen that introduction of inertial stresses incrementally increases the critical value of n for neck formation just above the solidification pqint but that the values are still quite low and the viscosity at the high shear stresses involved is reduced by orders of magnitude; cf the Newtonian value for the temperature which exists in the vicinity of the neck. This is an especially severe decrease when it is considered that, when air quenching is used, the temperature is well below the extrusion temperature and viscosity would be expected to be much higher than that at the spinneret.

8.3.3

Draw Resonance

As we have mentioned above, isothermal spinning can lead to severe instabilities and large fluctuations in fiber diameter when some critical attenuation factor is reached and when there is a discontinuity introduced in the form of rapid quench or solidification of some sort. For Newtonian fluids, theoretical analysisl04-106 has indicated that the critical attenuation factor is about 20. For power law fluids, theoretical analysis 107 ,108 has predicted that the critical value of AF will decrease as n decreases. Examples are AF = 2.8 for n = 0.3 and 4.8 for n = 0.5. Cooling of the threadline tends to stabilize it against draw resonance 110 since much higher values of AF are achieved in commercial spinning. On the other hand, some studies indicate that cooling can increase the severity of draw resonance in viscoelastic fluids (see ref. 111 for a discussion). It is fortunate that, in the air gap spinning process used for spinning of p-aramids, relatively low stretch factors are needed to achieve high orientation so that draw resonance can be avoided. In one experimental study of draw resonance,110 the critical draw down ratio was about 13 for polypropylene and the severity of the diameter oscillations increased more rapidly with increasing draw down when the air gap was 12 cm compared to 6 cm. In another studyl12 the critical draw down ratio was about ten. There are some data113,114 indicating that the power law exponent n for elongational flow of PP is about 0.4-0.5. One of these studies113 shows an effect of spinneret capillary dimensions on the critical stretch ratio for draw resonance with longer LID ratios giving higher critical ratios. In another study,115 the critical stretch ratio was also affected by capillary flow variables and reached values as low as five, the theoretical value for n = 0.5. 109 Finally, in the spinning of H 2S0 4 solutions ofPPD_T,116 the critical jet stretch for draw resonance was found to increase from about 12 to 21 as the size of the air gap was reduced from 2 cm to 0.35 em. Effects of capillary dimensions were also oqserved. If there was a bulge formed below the spinneret under the spinning conditions used, the critical attenuation factor would have been larger than the numbers quoted for critical draw ratio. In certain cases, it is possible to encounter severe draw resonance even when there is no external rapid quench. Thus it has been found that high molecular weight liquid crystalline polyarylates, which spin well at low molecular weights, manifest draw resonance at quite low spin stretch factors and this represents a formidable barrier to realizing the tensile strength potential of these polymers. 77 It was found that proper selection of polymer composition and spinning below the melting point (but above the freezing point) of the polymer eliminates draw resonance and permits attainment of higher strength fibers.

8.3.4 Structure Development in Spinning

8.3.4.1 Introduction The two major structural features of spun fibers which are of concern are degree of crystallinity and orientation. While the latter depends strongly on threadline tension for all of the flexible chain polymers, crystallinity can also be affected by threadline tension for polymers such as PET where

264

Generic Polymer Systems and Applications

orientation has a profound effect on crystallization rate. Beyond these structural parameters, the absolute uniformity of orientation, which has an inevitable relationship to orientation level in most cases, is an important determinant of the strength potential of a fiber having a given molecular weight. For spun fibers (e.g. those spun at high speeds) which will not undergo further drawing, more subtle features such as crystallite size and intercrystalline spacing can have important effects on fiber and fabric properties. In those situations where the spun fibers will be drawn, the orientation and crystallinity of the final yarn will be more affected by the details of the drawing process. However, orientation non-uniformities in the spun yarn cannot generally be overcome and propagate into the drawn yarn. There are several methods commonly used for evaluation of degree of crystallinity. One involves the use of wide angle (WAXS) and small angle (SAXS) X-ray diffraction. The former can be used to estimate degree of crystallinity and crystallite size from the sharpness of the diffraction spots as well as the structure and lattice parameters of the unit cell of the crystallite. However, there are significant differences of opinion on how to interpret the patterns and it is not certain that absolute values of crystallinity can be obtained from WAXS with which all experts will agree (see ref. 17 for a brief discussion of the situation for nylon 6,6). SAXS intensity is believed to be related to the degree of chain folding, 11 7 while the long period spacing is related to the length of the chain folds or crystallite to crystallite repeat distance. Another method which is commonly used to assess degree of crystallinity is fiber density, which is useful because the non-crystalline regions of the polymer or fiber have a lower density than the crystalline regions. In this method, the average specific volume (reciprocal density) of the fiber (V) is considered to be the weighted average of the specific volumes of the crystalline (Ve ) and noncrystalline (Va) components as indicated in equation (27) where fe is the weight fraction of crystalline material and is, of course, equal to 1 -fa (the weight fraction of amorphous or non-crystalline material). The densities of the components can be deduced from X-ray measurements and are, for example, 1.24 g cm - 3 for the crystals of nylon 6,6 118 and 1.09 for the amorphous regions. However they can also be deduced from IR measurements and in the case of nylon 6,6119 the respective densities were determined to be 1.22 and 1.07. Thus a fiber density of 1.14 is 36% crystallinity based on the X-ray parameters and 50% based on the IR. There are similar uncertainties with PET where the crystal density seems to depend somewhat on the temperature at which the fiber was exposed or annealed. Thus quoted densities for PET crystals range from 1.455 120 to 1.51,121 while the amorphous density is given as about 1.335. 120 The conformation of the methylene groups in PET can be in the trans or one of two gauche conformations. It is believed that the crystals are all trans, while the non-crystalline regions can be largely gauche or contain variable amounts of trans depending on the 'amorphous orientation'. IR measurements 121 give 1.326 for the density of amorphous regions if they are 100% gauche and 1.430 if they are 100% trans (which probably cannot be achieved). If these values are correct, then an amorphous fiber with density of 1.335 should be 900/0 gauche. (27)

IR spectroscopy can also be used for the assessment of crystallinity. Thus for nylon 6,6,119,122 a band at 936 cm - 1 is ap. indicator of crystalline regions while 1138 cm - 1 is an amorphous band. The band at 3305 cm - 1, the N-H stretching band, gives strong absorption when the electric field vector is perpendicular to the chain direction and thus can be used as a measure of orientation with polarized IR (IR dichroism). For PET,121, 123,124 a crystal index is given by the ratio of absorbance at 868 cm -1 to that at 1410 cm -1. The use of FTIR spectroscopy gives information on the conformations. The trans is given by the 973 cm -lor 1473 cm -1 band, while the gauche is given by the 898 cm - 1 or 1454 cm - 1 band. Orientation can be determined by X-ray diffraction, measurement of birefringence, IR dichroism or sonic velocity. X-ray essentially gives a measure of crystallite orientation by determination of the sharpness of an azimuthal trace of the intensity of a diffraction pattern. The angular displacement between the points whose intensity is half that at the maximum is called the orientation angle and it varies inversely with orientation. The Hermans orientation factor 125 is given by equation (28), where cos 2(J is the mean square cos orientation angle. Probably the most accurate and convenient method for assessing the orientation of spun fibers is to use the birefringence B, which is the difference of the refractive index of polarized light whose electric vector is either parallel or perpendicular to the fiber axis (i.e. B = nil - n.i). The orientation factor (fo) can also be defined as B/B m where B m is the birefringence of a perfectly oriented fiber. The birefringence of a semicrystalline fiber is the weighted average of the birefringence of the crystalline and non-crystalline regions of the fiber. The orientation factors for the two regions may be determined if their birefringence for perfect orientation and

Polymers for Fibers

265

the percent crystallinity are known, and efforts have been made to do so. However, there are some differences of opinion as to the true values. One report 126 gives Be = 0.29 and Ba = 0.20 for the crystalline and amorphous birefringences of PET for perfect orientation. Another report12 7 gives 0.22 and 0.19 for the two values but this seems rather low for Be since values of this magnitude are sometimes observed in real fibers. A theoretical calculation of the birefringence of an 'ideal' PET fiber 128 gives 0.2360 for the value. Another gives 0.24-0.25 for Be .129 The birefringence of a fiber of given orientation increases with the refractive index of the fiber. Thus it is believed to be about 0.08 for perfectly oriented nylon 6,6, while values as high as 0.064 are observed in commercial fibers. In general, polymers with aromatic rings will tend to have higher birefringence for a given orientation level because they have higher refractive indices so that the birefringence of Kevlar® aramid fibers has been reported to be about 0.44-0.46 130 or 0.60-0.75. 131 fo

=

(3cos 2 ()

-

1)/2

(28)

Sonic velocity is sometimes used to judge orientation factors with the relationship shown in equation (29)132 where Vu is the velocity of sound in an unoriented sample and v is that for the sample. See also refs. 133 and 134 for applications and discussion of the method, in particular the need to take into account the effects of varying crystallinity. (29)

8.3.4.2

Nylon 6 and nylon 6,6

In discussion of structure development in the spinning of nylon 6 and nylon 6,6, it must first be pointed out that, as with most polymers, the temperature at which crystallization occurs in a cooling melt is considerably lower than the melting point Tm . In addition there are considerable differences in the rate of crystallization which depend on polymer composition and molecular weight. Typically there is a maximum in the curve of crystallization rate vs. temperature, which results from an increase in the rate of formation of nuclei and decrease in the mobility of the polymer chain segments as temperature is reduced. For example, one relationship proposed 135 is Tmax = 0.82Tm (temperatures in K). This predicts a Tmax of about 134°C for nylon 6 and 168 °C for nylon 6,6. Another relationship136 introduces the effect of the glass transition temperature Tg as shown in equations (30a-c) and is believed to correlate with data for a wider range of polymers. Thus the data given for nylon 6 quote Tm = 500 K, Tg = 325 K and Tmax = 411 K (138°C), while for nylon 6,6, Tm = 545 K, Tg = 325 K and Tmax = 414 K (141°C). These data are not in full accord with the previous relationship and other data. 94 Moreover, effects of orientation can become critical, particularly with PET (see below). Tg (1 + Y) Tm (30a) Y

Tmax

=

(30b)

TITg Tg

-

50

(30e)

The birefringence development in amorphous polymers is not very pronounced, except at very high deformations. Thus in the spinning of nylon 6,99 birefringence development becomes appreciable only when the temperature drops to 120°C or below, as crystallinity develops. This development of crystalline orientation can result from crystallization of oriented chains, with a subsequent deorientation of the remaining amorphous chains. In addition, it can result from the effect of tension on the crystalline fiber to promote drawing. This process will be influenced by the aerodynamic forces, which increase directly with the path length. Birefringence increases directly. with wind-up speed or spinline stress. 96 ,99,137 In one study137 with nylon 6,6 of 0.95[11J, birefringence increased from about 0.0035 at a stress of 5 x 10 5 Pa to about 0.018 at 2 x 106 Pa. On-line measurements showed no significant birefringence development until about 60 cm from the spin- . neret, where the filament temperatures were about 160°C or lower. Nylon 6,6 crystallizes in the threadline to form the normal triclinic structure 118 ,137 but nylon 6 crystallizes much more slowly, much of the crystallization occurring on the bobbin in a poorly defined pseudohexagonal structure. 137 This is in accord with studies which indicate that nylon 6,6 crystallizes faster by about an order of magnitude. 138, 139 Thus in one study141 with nylon 6, birefringence development started about 40 cm from the spinneret but did not reach very high levels, on-line, at stresses of 2 x 106 Pa by the time the final yarn velocity was reached (B about 0.0018).

266

Generic Polymer Systems and Applications

However, the final yarn birefringence was about 0.024 after equilibration and crystallization of the spun yarn in the time period before the yarn reached the wind-up and after lagging. In another study,96 similar results were found and, for example, a nylon 6 fiber with a limiting value of B = 0.002 in the quenched threadline had, after conditioning, a value of B = 0.02, or about a tenfold increase. The diffusion of H 20 into the fibers during the conditioning period is to a large extent responsible for the crystallization during the lag period since the Tg of the fiber is thereby reduced to room temperature or below. The crystallization and orientation increases which occur on the bobbin are accompanied by a lengthening of the yarn, which can cause practical problems of loosely wound packages. The magnitude of the birefringence increase on lagging of nylon 6 increases with increasing spinning speed up to about 3000 m min - 1 but then decreases in the range of 4000-7000 m min - 1. 142 The birefringence of the lagged yarn increases rapidly with spinning speed up to about 3000 m min - 1 to a level of about 0.033 but then increases much more slowly to about 0.04 at 7000 m min - 1. The increase in length of spun yarn as water is absorbed also occurs with nylon 6,6 as spinning speed is increased. However, its magnitude is not as great as with nylon 6. Nevertheless, at the intermediate speeds where the effect is greatest, steam conditioning of the yarn below the quench chimney is needed to provide stable packages. In the high speed spinning of nylon 6, spun fibers contain a mixture of two crystal forms, called ex and y.143. 144 The former is the stable monoclinic crystal, while the latter, whose structure is not fully understood, is a less stable form whose relative abundance increases with spinning speed up to about 5000-6000 m min - 1. Annealing or hot drawing of the spun yarn causes conversion to the ex form to an extent which depends on the draw ratio. In general, nylon 6,6 fibers are in their stable, triclinic ex form when examined at room temperature. A high temperature, hexagonal crystal form can be observed at temperatures above about 160°C but it converts rapidly to the triclinic form at room temperature. For both fibers, the increase in orientation which results as spinning speed and spinline tension are increased is accompanied by an increase in tensile strength of the as-spun fibers and a corresponding decrease in elongation to break. Thus, in one series of experiments with nylon 6,144 tensile strength increased from about 200MPa (about 1.75dNtex- 1) at 500mmin- 1 to about 450MPa at 5000 m min -1, while the elongation to break decreased from 600% to about 100%. In this research, further increases in spinning speed resulted in a decrease in tensile strength which was attributed to non-uniformities across the cross section of the filaments resulting from the relatively small but significant temperature gradients from the center to the skin of the filaments. For moderate levels of orientation in nylon 6,6, the birefringence of spun yarn corresponds to an effective draw ratio R s (vs. a completely unoriented fiber) as indicated by equation (31).140 This relationship should also apply fairly well to nylon 6, which has about the same refractive index. Rs

8.3.4.3

=

1

+

22.2B

+

284B 2

(31)

Poly( ethylene terephthalate)

The development of structure in spinning of PET is more complex than that for the nylons. A major reason for this is the slow crystallization rate (including a relatively long induction period) for unoriented PET and the dramatic effect of orientation on the rate. For example,145.146 at 120°C, the half time for crystallization of unoriented PET has been reported to be in the range 840 to 210 s. Differences between observers may reflect differences in the molecular weight of the samples. For example,147 the crystallization rate of PET decreased by a factor of about six as Mn was increased from 19000 to 39 100. For polymer with M n of 21 000, a rapid decrease in half time was observed148 as the birefringence of the fibers increased. For B = 0.005, it was 20 s, for B = 0.027 it was 0.7 s and for B = 0.080 it was less than 0.01 s. This exponential increase in rate with increasing orientation correlates with the increase in density and in trans methylene content of 'amorphous' yarn. 123 . 149 A detailed study of the relationship between amorphous orientation and crystallization rate is given in ref. 150. The literature on structure formation and property development in the spinning of PET is extensive and has grown rapidly in recent years with the advent of high speed spinning. Previously, yarns were spun at relatively low speeds and drawn. The spun yarns were amorphous because the residence times at temperatures where crystallization could occur were too short for the birefringence levels obtained at these speeds. The advent of partially oriented yarns (POY) needed for high speed false twist texturing to provide improved stability vs. storage time and improved operability at higher texturing speeds pushed speeds into the 3000-3500 m min - 1 range but these yarns were still of low crystallinity. However, as research and commercial development pushed to still higher speeds, it was found that crystallinity continued to increase with beneficial effects on shrinkage, stability and

Polymers for Fibers

267

tensile properties. At 6400 m min -lor higher, as-spun yarns with many of the properties of drawn yarns, but with some advantages (e.g. dyeability), could be produced. 151 The subject of spinning of PET, and especially high speed spinning, is reviewed thoroughly in several chapters of ref. 95 142 ,143,152-154 and in refs. 155-160. Here the key points of these extensive researches with respect to development of structure and properties will be summarized. . Fiber birefringence increases directly with spinning speed. In at least one report 154 there is a fairly sharp upturn in the curve of B vs. speed at about 4000 m min - 1, where significant crystallinity starts to develop. There will obviously be differences between different sets of data if polymer molecular weights are not the same. In one set of data 158 with polymer of 0.62[1]], B was 0.01 for 1600 m min - 1, 0.03 for 2800 m min - 1 and 0.058 for 4500 m min - 1. Data for B at 6500 m min - 1were in the range of 0.12-0.14. 154 In general, the calculated amorphous orientation factor increased much more slowly with increased spinning speed than did the crystalline orientation factor as determined from X-ray diffraction. Density increases slowly with spinning speed up to about 3500 m min - 1 and is only about 1.340-1.343 at this point with the fiber~ still being of very low crystallinity. The content of crystalline trans conformation starts to increase appreciably143 only at about 4000 m min - 1. As spinning speed is increased to 6000 m min -1, density rises to about 1.39, reflecting fairly high levels of crystallinity. X-ray measurements of crystallinity indicate an increase from essentially zero for a birefringence level of 0.02 to about 40% for B of about 0.12. 151 Other data 143 with different calculationprocedures indicate crystallinity levels of only 24 % for yarns of this latter level of orientation. One manifestation of the non-crystalline character of the yarn spun at lower speeds is a very high shrinkage when the yarn is immersed in boiling water. This shrinkage (BaS) increases with spinning speed up to about 3000 155 or 4000 143 mmin- 1 with maximum values ranging from 60-70%. The absolute amount of shrinkage measured for samples containing some nuclei can depend on how rapidly the sample is heated, which can influence the relative rates of crystallization and deorientation. It should be noted that 60% shrinkage corresponds to undoing a draw ratio of 2.5, while a shrinkage of 70% corresponds to undoing a draw ratio of 3.3. This gives a minimum measure of the degree of extension of the amorphous polymer chains in the yarn spun at 3000-4000 m min - 1. As spinning speed increases to 5000 m min -1, shrinkage decreases to about 2%. This is lower than typical shrinkages for drawn yarn, which range from 7-10%. As spinning speed is increased, the fiber temperature at which crystallization occurs increases (i.e. less supercooling is required to produce the required nuclei). 143, 160 Thus at 4750 m min -1 the crystallization temperature has been estimated 143 at 186 ac, while at 6000 m min -1 it is 218 ac. The density of the crystals also appears to increase, from 1.488 to 1.501 at the two speeds respectively. Crystal sizes of high speed spun yarns tend to be larger than for drawn yarns (~ 5 nm vs. about 4.5 nm). Thus yarns spun at 6400 m min -1 had a crystal size of 7.2 nm. 154 Long period spacings are also much higher ( ~ 30 nm vs. 11.5 nm). This is believed to relate to the 4-7 fold higher dye rate for the yarns spun without drawing (i.e. because of greater distances between crystallites). However dye rate decreases with increasing speed so that amorphous orientation may be importantly involved as well. As with other fibers, tensile strength increases and elongation to break decreases with increasing spinning speed. For example, in one case 158 tensile strength was found to increase from 1.6 to 3.3 dN tex - 1 as spinning speed increased from 2000 to 4500 m min - 1, while elongation to break decreased from 1990/0 to 680/0. In another example,154 as speed increased from 2800 to 6000 m min - 1, tensile strength increased from 2 to about 3.5 dN tex - 1, while elongation decreased from 1500/0 to about 40%. At still higher speeds, tensile strength deteriorated somewhat, presumably because of orientation gradients across the filaments. Yarn spun at 3200 m min ~ 1 or lower had a stress-strain curve which showed a well-defined yield point at which there was a decrease in stres& and a subsequent flat region preceding a sharp upturn in stress. 154 This is typical of lower speed yarns which are to be drawn (see Figure 1). The yield could still be observed in yarn spun at 4575 m min - 1, but the subsequent flat region had almost disappeared. At higher speeds, the shapes of the stress-strain curves showed no stress decrease and were similar to those of drawn yarns except that the initial moduli were somewhat lower, but with an increasing trend in modulus as spinning speed was increased.

8.3.4.4 lsotactic polypropylene In the spinning of isotactic polypropylene with air quenching, highly crystalline yarns are obtained, generally in the monoclinic (Lt) form. 142, 161 For a sample with [1]] = 2.24 (M w = 2.74

268

Generic Polymer Systems and Applications 200

150

~

&

100

50

100

150

200

Elongation (%)

Figure 1 Typical stress-strain curve for undrawn PET yarn spun at moderate speeds; drawn at 50°C

x 105), extruded at 230°C and collected at low speeds (50mmin- 1), significant birefringence development started at a distance of about 100 cm from the spinneret when the temperature had decreased to about 100°C and crystallization began. 161 The relationship between spinline stress and spinning speed was less than a first power dependence up to 600 m min - 1, the maximum speed studied. This again provides evidence of non-Newtonian elongational viscosity with a power law exponent much less than unity. Birefringence increased with increased spinline stress from about 0.008 at 100 KPa to 0.02 at 10 MPa. The degree of crystallinity of the spun yarn increased only to a small extent with increasing spinline stress from about 53% at 10 KPa to 570/0 at 10 MPa. In other experiments 142 birefringence for yarn spun at 230°C increased rapidly with spinning speed to a level of about 0.015 at 1000 m min -1 but then increased relatively slowly to about 0.022 at 4000 m min -1. Density was almost unaffected by speed for this spinning temperature, changing from 0.900 to 0.901 over this speed ra..1ge. X-ray diffraction showed that the crystalline orientation factor also increased rapidly with speed to a level exceeding 0.8 but that the calculated amorphous orientation factor increased gradually and reached a level of only about 0.3 at speeds as high as 6000 m min - 1. The stress-strain curves for spun yarns prepared at 550 m min -1 or higher 161 no longer showed the yield point and stress drop followed by a flat region that existed in yarns spun at 50 m min - 1 and which are found with PET yarns up to much higher speeds. In this study, tensile strength of spun yarn increased from about 50 MPa to 200 MPa as spinline stress increased from 25 KPa to 10 MPa, while elongation to break decreased from about 14000/0 to 400%. Initial modulus ranged from about 500 MPa to 2500 MPa over this range of spinning stresses.

8.3.4.5

Gel-spun polyethylene

One of the interesting developments in recent years is 'gel-spinning' of ultra-high molecular weight polyethylene to form very high strength, high modulus fibers after drawing at high draw ratios. In this technology86, 162 -164 polyethylene with Mw preferably greater than 106 is spun as a dilute solution (2-10%) in a solvent such as decalin or paraffin oil. Spinning temperatures are 130-170 °C. The fibers are wet-spun with an air gap, preferably with a low spin stretch factor, into a quench bath (e.g. H 20). The solvent is extracted (e.g. with hexane or trichlorofluoroethane) and the fibers are dried and then hot drawn (e.g. 100-140°C) at very high draw ratios to give tensile strengths of 3-4 GPa (31-42 dNtex- 1) and moduli up to about 130 GPa (1400 dNtex- 1). Alternatively, the drying can be accomplished simultaneously with drawing. An explanation which has been advanced86, 165 for the high drawability of these spun fibers (in addition to the high molecular weight) is that the polymer contains only a few entanglements, just enough to provide coherence, and therefore it is relatively easy to disentangle'them in drawing so that strengthening can occur. One/definition of the preferred spun yarn structure for this process 86 is a crystal orientation factor less th~n 0.1, crystal index less than 0.75 and microporosity less than 10%. It should be noted that polyethylene has the potential for reaching very high crystallinity levels so that 900/0 or higher has been attained in model experiments by hot drawing. 166

Polymers for Fibers

269

8.3.4.6 Liquid crystalline solutions and melts It has been seen above that very high melt draw down ratios can be obtained in melt-spinning of flexible chain polymers without achievement of high birefringence. The elastic extension of the polymer chains before crystallization depends on the threadline stress and the modulus of the polymer chains in the melt. For natural rubber, draw ratios of three and five result in birefringences (B) of only about 0.005 and 0.02 respectively,167 the relationship between B and draw ratio curving upward as the draw ratio is increased. On the other hand, birefringence of semicrystalline polymers initially increases rapidly with draw ratio and then more slowly at draw ratios above three or four, in general accord with theories of affine deformation of such polymers (see Figure 2 for nylon 6,6). Thus, orientation factors are already quite high at draw ratios of five or six. For example, a nylon 6,6 yarn drawn to a draw ratio of six from the completely unoriented state has a birefringence of 0.061 and is quite strong. Similarly a polyester fiber with an initial birefringence of 0.008, when drawn at 200°C to a draw ratio of 5,168 had a birefringence of 0.2483, representing a calculated orientation factor of over 95% (see Figure 3). When a rod-like polymer such as a p-oriented aramid is dissolved in H 2 S0 4 at concentrations of 15-20%, the formation of highly anisotropic solutions consisting mainly of nematic liquid crystals provides two key benefits. The first is the major reduction in melt viscosity discussed above. The second is the high efficiency of orientation at modest spin stretch factors (SSF) through affine deformation of the liquid crystalline melt. Thus in the air gap spinning of H 2 S0 4 solutions of PPD_T 169 high orientations were attained with SSF as low as 1.5 and as high as six (attenuation 0.07

0.06

0.05 Q)

u cQ)

0' C

:E ~

iii

0.04

0.03 0.02 0.01

2

4

3

5

6

7

Draw ratio

Figure 2 Birefringence vs. draw ratio for nylon 6,6 (B o = 0)

0.3

./

.

~

~

0.2

0.1

2

3

4

5

6

Draw ratio

Figure 3

Birefringence vs. draw ratio for PET (B o =0.OO8), drawn at 200°C (data from ref. 168)

270

Generic Polymer Systems and Applications

factors may have been somewhat larger to the extent that there was a bulge below the spinneret prior to attenuation in the air gap) and tensile strengths of the fibers were very high (e.g. as high as 24-26.5 dN tex - 1 or 27-30 g den - 1) at only modest molecular weights, in the range of 20000-25000. In this process air gap lengths of 0.5-2 cm were typically used. The attenuated fibers travelling at (for example) 275 m min -1 were quenched in very cold water (e.g. 1°C), washed, neutralized (e.g. with NaHC0 3 or NaOH solutions) and dried (e.g. on heated rolls at 160°C) at relatively low tensions (less than about 0.25 dN tex '- 1). The dried fibers were highly oriented as indicated by an X-ray orientation angle of about 11 0. The fibers were crystalline and had an apparent crystallite size of less than 5.2 nm (e.g. 4 nm). They were also characterized as having a high degree of radial orientation as indicated by electron diffraction patterns of thin oblique sections of the fiber. Heat treatment under tension of such fibers under nitrogen at, for example, 525°C resulted in a substantial increase in fiber modulus and a further reduction in orientation angle to 9°. Crystallinity increased as indicated by an increase in fiber density from 1.44 to 1.45. This can be compared with one estimate 170 of the crystalline density of PPD-T as being 1.48 g cm - 3. The melting point of a 20% solution of PPD-T in H 2S0 4 as measured by DTA 171 has been reported to be about 75°C, but on cooling the crystallization temperature is about 45 °C so that supercooling does occur. It is not certain what effect the rate of cooling will have on this crystallization temperature. However, since there is little time for cooling in the air gap, crystallization of the oriented spinning solution occurs very shortly after entrance into the quench bath, simultaneously with initial coagulation of the fiber surface by water diffusion into the fiber and H 2S0 4 diffusion out. Cooling occurs faster than coagulation 172 since the thermal diffusion coefficient is several orders of magnitude higher than that for diffusion of water or acid. Examination of Dupont's Kevlar® fibers by various research workers has led to differences of opinion concerning the degree of crystallinity of the as-spun (Kevlar® or Kevlar® 29) and heattreated fibers (Kevlar® 49 or fibers annealed by the researcher). Some X-ray studies give percent crystallinity values of about 61 % 173 or 68% 174 for as-spun fibers and 71-75% for heat-treated fibers. Others believe that the fibers are fully crystalline 175 ,176 with periodic defect layers and prefer a paracrystalline model with statistical lattice disorder over a longer range. Still others 177 question the ability to determine the crystallinity of these fibers by existing techniques. The need to invoke intercrystalline regions for wholly aromatic polymers has been supported by studies of mechanical properties vs. stress and by sorption experiments. 178 ,179 This writer tends to agree with the need for the concept of non-crystalline regions because of the moisture regain levels of the various fibers as well as the effects of stress on mechanical properties. For example, Kevlar® 49 fibers have an equilibrium H 20 content of about 3.7% at 55% RH (relative humidity),180 which is in the same range as that for nylon 6,6 having a crystallinity level of about 50%. The regain for the as-spun fibers is even higher, which indicates the expected effect of annealing. Since it is clear from the literature quoted above that the polymer chains are not folded, the lack of perfect registry may be a result of deviations in the paths of the polymer chains, which are not completely rigid. In any event, this writer prefers the use of the term non-crystalline regions rather than amorphous regions for these areas whose inter-amide hydrogen bonds are more readily replaced by H 20, since it is more likely that they are some kind of liquid crystal because of the chain stiffness, concentration and temperature. Steric effects such as trapped entanglements may be responsible for the lack of perfect order. Electron and optical microscopic examination of Kevlar® fibers combined with selective etching experiments show the existence of a pleated, fibrillar structure with periodicities in the range of 200-1000 nm. 172 ,173, 175, 181, 182 In the spinning of thermotropic (liquid crystalline) polyester melts, the efficient development of orientation is similar to that seen in the spinning of liquid crystalline aramid solutions. In one study183 of the spinning of a copolymer prepared from 60 mol% of p-acetoxybenzoic acid, 20 mol% of 2,6-naphthalerie diacetate and 20 mol% of terephthalic acid and having an 1Jinh of 5.4 (0.1 % in PFP at 60°C) and spun at 340°C, modulus and tensile strength increased steadily with increasing spin stretch factor up to a value of about 16 and then levelled off or decreased somewhat (tensile strength). The X-ray orientation angle decreased from 180° for no draw down to 8.5° for 16 SSF. Thus it is true, in general, that as-spun thermotropic polyarylate fibers are highly oriented. 184, 185 Their crystallinity levels vary greatly depending on the specific polymer composition used. Since it is frequently true that the as-spun fibers have only moderate tensile strength because of relatively low molecular weight, they are heat treated, preferably under nitrogen or other inert gas, to increase molecular weight and strength. 185 As a result, crystallinity will tend to increase, while orientation will increase only slightly, if at all. For example,185 spun fibers of an AA-BB thermotropic polyarylate, poly(chloro-1 ,4-phenylene terephthalate/2,6-naphthalate) (70/30), with an orientation angle of 21 ° and tensile strength of 5.8 dN tex -1 were subjected to prolonged heat treatment with no

Polymers for Fibers

271

applied tension and with staged increases in temperature (290°C maximum). The fibers increased in strength to 26 dN tex - 1, while the orientation angle decreased to 18°. In spite of this, the modulus did not increase from its as-spun level of 483 dN tex - 1. This may indicate a slight relaxation of the non-crystalline regions during heat strengthening. In one study with poly(phenyl-l,4-phenylene terephthalate),73 a direct correlation of tensile strength with molecular weight as indicated by melt viscosity was shown to exist. Since the melting point of these polyesters varies in expected ways with copolymer composition, it is likely that the heat of fusion varies similarly, and also the degree of crystallinity. There are reports 186 that prolonged heating of these polyesters can lead to a degree of blockiness but it is not clear whether this contributes in any way to increases in strength. The ability to change by copolymerization the heat of crystallization during quenching of the melt and hence the rate of crystallization has made it possible to spin77 while maintaining the spinneret below the melting point of the polymer (but above the freezing point) provided that the DSC heat of crystallization is less than about 10 J g - 1. This has permitted the spinning of high molecular weight polymers without draw resonance, as indicated above.

8.4 DRAWING In general, spun yarns of flexible chain polymers have relatively low orientation and must be drawn to develop the mechanical properties desired for the given end use. An exception to this has been discussed above in connection with very high speed spinning. A common view of what happens in drawing is that coiled polymer chains or folded chain lamellae are extended until some tie chains between crystallites become taut. Up to this point, the stresses required are relatively low. At still higher draw ratios, it is necessary to pull polymer chains through the crystallites so that an increasing fraction of the tie chains become taut and this increases the tensile strength of the fiber and, frequently, the modulus. This process requires higher stresses and, usually, higher temperatures to achieve these properties at an acceptable level of process breaks. The force vs. elongation curve for undrawn fibers may have the shape shown in Figure 1, which is typical of PET fibers. In this case, there is an initial yielding which can be regarded as a failure of the undrawn fiber structure and this is followed by a relatively flat region which ends in the 'natural draw ratio'. It should be pointed out that the true stress is increasing in this region, since the fiber cross sectional area is decreasing. Beyond the natural draw ratio, the stress increases steeply until another yield region is reached, which finally results in tensile failure. The length of this second yield region depends on such factors as spun fiber orientation, molecular weight and fiber defects. Another type of undrawn yarn force vs. elongation curve is shown in Figure 4 which may apply to semicrystalline fibers tested above their glass transition temperature (e.g. nylon 6,6 at 25°C and 68% RH or dry above 50°C). In this case, there is no clear yield point or perfectly flat region even though similar processes are occurring. For fibers with behavior similar to that shown in Figure 1 a sharp neck is observed when drawing occurs, while with fibers which behave as in Figure 4, the deformation is more gradual. The natural draw ratio decreases with increasing spun yarn birefringence, as might be expected. For PET, for example,187,188 it decreases fro~ five for B = 0 to about two for B = 0.02. 3

2

2

3

4

Elongation (%)

Figure 4 Typical stress-strain curve for undrawn nylon 6,6 spun at moderate speeds; drawn at > 50°C

272

Generic Polymer Systems and Applications

The initial yielding which occurs with an undrawn yarn appears to be a non-Newtonian creep or flow process which acts in accord with the Eyring absolute rate theory.189 For high stress levels, typical of those encountered in drawing or in tensile failure, the rate of flow increases exponentially with stress and l/T(K). When an undrawn fiber is stressed, it creeps at a relatively slow rate and then there is a sudden increase in rate, as if the fiber has failed. This is the onset of drawing. Data for nylon 6,6 190 at various stress levels are in accord with the theory in that log(time to initiate draw) decreases linearly with increasing stress. Similarly, expected effects of strain rate and temperature were observed for nylon 6,10 191 ,192 and for PET. 187 ,193 Effects of temperature were consistent with an activation energy for the flow process for PET of 230 kJ mol- 1 (55 kcal mol- 1). One consequence of this exponential dependence of flow rate on stress is that the yield stress (measured at some given rate of testing) decreases linearly with increasing temperature, in the absence of changes in the fiber structure as temperature is increased (e.g. changes in crystallinity or crystal structure or loss of orientation). It should be noted that the draw yield stress for as-spun yarns increases with increasing spun yarn birefringence. 142, 154 These theoretical aspects of the absolute rate theory will be discussed in more detail below in connection with the subject of tensile failure. In carrying out practical drawing procedures, it is common to use various forms of 'draw assist' devices such as pins (heated or unheated), heated plates (flat or slightly curved) or heated tubes over which the yarn is conducted with helical wraps. One of the purposes of these devices is to reduce the tension developed in the drawing operation at points upstream of the draw zone. This is done, in part, to prevent slippage on feed rolls and/or to control the point or region where draw occurs. The decrease in upstream force F 1 vs. the draw force F 2 is given approximately by the belt friction or capstan equation (32), where J.1 is the coefficient of friction (e.g. 0.1-0.4), fJ is the angle of wrap in radians (e.g. 2n for one circular wrap) and

8.5 STRUCTURE-PROPERTY RELATIONSHIPS 8.5.1

Modulus

The modulus of a fiber is determined by the following structural factors: crystallinity, polymer chain stiffness, orientation (crystalline and amorphous) and intermolecular forces (e.g. H-bonds or covalent cross links). Consider first the modulus of a fiber crystal oriented in the direction of the fiber axis. In one theoretical study in which bond stretching and opening of valence angles were considered,196 the modulus of a nylon 6,6 crystal was calculated to be 157 GPa (1290 dN tex -1 for a crystal density of 1.22), while that of a PET crystal was 146 GPa (about 1000 dN tex -1 for crystal density of 1.46). Another set of calculations for various fiber crystals along with experimentally determined values 197 (e.g. from X-ray c spacing vs. stress) is shown in Table 1. The observed value of 153 GPa reported here for Kevlar® can be compared with another X-ray observation 199 for Kevlar® 49 of 200 GPa vs. 110 G Pa measured on an Instron tensile tester for this fiber. Thus it is seen that crystal moduli are in the range of 150-300 G Pa. When the actual moduli of various fibers are examined, It IS seen that they range from 0.01-0.1 dNtex- 1 (ca. 0.001-0.01 GPa) for elastomeric fibers which are mostly amorphous to

Polymers for Fibers

273

Table 1 Fiber Crystal Moduli a Fiber

Kelvar® Polyethylene PET

Nylon 6 (a)

a

Crystal moduli (GPa) Observed Calculated

153 235 108 165

182 316 95 244

H. Tadokoro, Polymer, 1984, 25, 158.

1000-2000 dN tex - 1 (100-200 GPa) for highly drawn polyethylene whose crystallinity ranges from 70 to > 90%. One oversimplified way of looking at the effect of crystallinity is to consider the crystalline and non-crystalline (or amorphous) regions to be in series so that the compliance (reciprocal modulus) is the weighted sum of the compliances of the components. Thus, for exampl~, if the crystal modulus is 160 GPa and that of the amorphous regions is 2.5 GPa, the modulus of the fiber will be 4.9 GPa for 50% crystallinity, 21.9 GPa for 90% crystallinity and 38.5 GPa for 95% crystallinity. However, it is possible that the modulus of the amorphous regions will increase with increasing crystallinity since the tendency of these regions to retract after drawing may be reduced. Therefore modulus may increase to a greater extent with increasing crystallinity than shown here. Similar calculations have been made with more complex models (e.g. series-parallel arrangements) but the general trends are similar. It should be noted that it is only with polyethylene that very high crystallinity levels have been attained. Thus far, with the other main fiber types from flexible chain polymers, crystallinity levels have probably not exceeded about 65 or 70%. It should also be mentioned that the observed moduli are generally measured at relatively high rates of elongation (e.g. 1% S - 1) and that they may be much lower at much slower rates as a result of stress relaxation or creep. The stiffness of the polymer chain has a profound effect on fiber modulus, particularly when crystallinity levels are not too high. Thus the modulus of drawn nylon 6,6 fibers is about 45 dN tex - 1 but it may decrease to as low as 10 dN tex - 1 for fibers relaxed in boiling water. The corresponding range for PET fibers is 25-160 dN tex - 1, reflecting the influence of the aromatic rings. The observed range for Kevlar® fibers, in which both the components are p-linked, is 199 465-777 dN tex - 1, while for Nomex® fibers, in which the components are m-linked and hence are not so stiff, the modulus is only about 88 dN tex - 1. The effect of orientation on modulus is also profound. When fibers are drawn, crystalline orientation increases with increasing draw ratio as described above. This, per se, should increase modulus by about cos 4 of the orientation angle (which decreases with increasing draw ratio). If this was the only factor, the modulus would become relatively insensitive to draw ratio when orientation angles became low (e.g. the difference in modulus between 10° and 5° would be only about 4-5%). However for PET, the modulus increases steadily with draw ratio. For example,123 the modulus increased from about 3 GPa at a draw ratio of two to 12 GPa at a draw ratio of five and 14 GPa at a draw ratio of 6.5. Similarly, for gel-spun polyethylene 162 the modulus increased approximately linearly with increasing draw ratio from about five to 90 GPa for the draw ratio range of 3-30. This behavior points to the combined effects of amorphous orientation and improved tie chain length distribution in increasing modulus. For polymers of moderate crystallinity, amorphous orientation is sensitively affected by retraction which may occur after drawing, either by relaxation at room temperature or by thermal treatments such as boil-off or heat setting of fabrics so that fivefold reductions in modulus can occur as indicated above. Relaxation at room temperature and ambient relative humidity of drawn yarn occurs more readily for yarns which are above their glass transition temperature under these conditions. Intermolecular forces can have important effects on the fiber modulus for flexible chain polymers. In the case of polyamides such as nylon 6,6, hydrogen bonding plays an important role. For example,20o nylon 6,6 yarns which have been preshrunk in boiling water have a modulus of about 40 dN tex -1 at 0% RH. The modulus decreases linearly with increasing RH to about 10 dN tex -1 at 1000/0 RH. Nylon yarns which contain radiation-grafted maleic acid to the extent of 1170 equiv.10- 6 g of -C0 2H groups have a dry modulus of about 105 dNtex- 1 because of increased H-bonding involving the -C0 2H groups. The modulus again decreases linearly with increasing RH and is only about 12 dN tex -1 at 100% RH. These effects are largely absent in the case of PET tested at room temperature since H-bonding does not contribute to the modulus. PS 7-J

Generic Polymer Systems and Applications

274 8.5.2

Tensile Strength

The factors which influence the tensile strength of a fiber include (1) the 'ultimate' tensile strength for a perfectly oriented fiber having the degree of crystallinity which can be practically obtained and where all of the tie chains are bearing the load equally; (2) the molecular weight (£In); (3) the draw ratio, which affects both crystalline and amorphous orientation as well as tie chain length distribution, the maximum value of which is influenced by £In; (4) the uniformity (particularly along the fiber) of thickness and orientation; and (5) the level of defects such as end groups and grosser heterogeneities. Over the years, several different approaches have been used to predict ultimate tensile strength. One method involves the determination of crystallinefc and amorphousfa orientation functions for a series of yarns and plotting them against tensile strength. Extrapolation to the point where fa = 1 (fc is close to unity for all of the higher draw ratio yarns) is used to determine ultimate strength. For nylon 6,6 201 this has led to a value of about 18 dNtex- l . For nylon 6 and PET, a similar approach has given 21 and 9.7 dN tex -1 respectively.202 In still another study,203 ultimate strength was determined to be 13 dN tex -1 for polypropylene and 8.8 dN tex -1 for PET. Based on recently reported achievements, these estimates appear to be low for at least some of the fibers. Perhaps the effects of molecular weight have not been adequately considered. Another semiempirical approach which has been applied to ultra-high molecular weight polyethylene 204 determines the relationship between tensile strength S and modulus M for yarns (M w = 1.5 x 106) drawn at various draw ratios (S = 0.105M o.77 ). Extrapolation to the theoretical modulus of 250-300 GPa gives for the ultimate strength 7.4-8.5 GPa or about 74-85 dN tex -lor slightly higher if the density is less than unity. One popular view of tensile failure is that the process involves the breaking of covalent bonds. If this is the case, then ultimate strength can be estimated by calculation of the force to break all of the bonds in a given cross section of the fiber. Thus with the assumption of a Morse potential for the bonds in question (equation 33 where E is the energy at interatomic distance rand D is the bond dissociation energy), differentiation with respect to r gives the force. The maximum force can be shown to equal 0.5 aD. With this approach, the ultimate tensile strength has been calculated to be 20.6GPa for PET 205 (ca. 150dNtex- 1) based on the maximum force for a C-C bond of 5.175 x 10- 9 Nand 3.98 x 10 12 chainsmm- 2). In another report 206 the calculated ultimate strengths are 15.2-16.2 GPa for PET, 17.2-19.6 GPa for nylon 6,8.1-9.2 GPa for polypropylene and 18.6-19.6 GPa for polyethylene. Even higher estimates have been made for these fibers. 207 It should be pointed out here that the generally accepted dissociation energies of the C-C, C-N and C-O bonds are about 347,293 and 351 kJmol- 1 (83,70 and 84kcalmol- 1) respectively. In some discussions of tensile failure, much lower energies are sometimes quoted but these are based on activation energies of degradation reactions which frequently do not involve homolytic bond cleavage (e.g. fJ elimination reactions for polyamides). This author estimates that the weakest bond dissociation energy for Kevlar® is 42 kJ mol- 1 (10 kcal mol-I) greater than that for the aliphatic polyamides; this is based on the TGA decomposition onset temperature, which is 100°C higher for Kevlar® than for aliphatic polyamides. EID

=

exp [ - 2a(r

-

ro )]

-

2exp [ - a(r

-

ro )]

(33)

Another approach to calculation of ultimate tensile strength involves a different view of the mechanism of tensile failure. This view considers failure to be a creep process involving motions of polymer chains through crystallites under the influence of the applied stress. One of the main reasons for favoring such a mechanism is the fact that activation energies for tensile failure (to be discussed below) correlate more closely with melting points than with bond energies. They are also similar to the activation energies for drawing at high draw ratios, which is a creep process. In view of the stress levels involved in typical failure experiments, the Eyring absolute rate theory of flow in its simple exponential form applies. 208 - 210 When a fiber is placed under a fixed load, integration of the equation for the exponential dependence of the creep rate on the stressfis simple and leads to the result shown in equation (34), where t b is the time to break under the fixed stress f, h is Planck's constant (6.626 x 10- 34 Js), k is Boltzmann's constant (1.38 x 10- 23 JK- 1), Tis in K, n is the number of jumps before there is enough stress concentration to make the process catastrophic, AF* is the free energy of activation for the failure process, A is the cross sectional area per chain, d 1 is the 'jump distance' (in cm) and is twice the distance to the top of the free energy barrier and WI is the fraction of chains bearing the load in a real fiber, where the tie chain length distribution is relatively broad. If AF* is in J mol- 1, R = 8.314. This equation predicts that the log of the time to break increases linearly with decreasing stress. It also predicts that, at some fixed breaking time, the

Polymers for Fibers

275

strength will decrease linearly with increasing temperature if there are no structural changes along the way. It also predicts that the extrapolated zero strength temperature To (at 1 s testing time, for example) is related to the extrapolated zero stress intercept a at room temperature on the log time axis. Equation (34) can be written in the form logt b = a - bfor logt b = a(1 - flft) whereft is the stress required to break the fiber in 1 s and a is given by equation (35). There is no good information concerning the value of n but changing it from one to ten will only have a small effect on the calculation of ~F* from the time dependence of breaking strength. At 25°C, 10g(kTlh) = - 12.8. Thus for 1 s break time, ~F*/2.303RTo is about 11.8-12.8 for n = 10 to 1, if To is determined from data over a relatively narrow temperature range. Similarly the zero stress intercept a is about (12-13) (ToIT - 1). (34) In(h/kT) + In(n) + tiF*/RT - fAd t /2kTw t a

10g(h/kT)

+

log(n)

+ tiF* /2.303RT

(35)

Data for lifetimes under load vs. stress are not necessarily unique for a given fiber type. For example, in Figure 5 are shown data for melt-spun polyethylene with Mn of about 2 x 104 which has been hot drawn at 110-120 °C 214 (line 1). In the same figure are also shown data 215 for gel-spun ultra-high molecular weight polyethylene (M w = 4 x 10 6 ) which has been hot drawn at a maximum temperature of 160°C (line 2). It can be seen that the combined effects of molecular weight and drawing temperature (crystallinity) have resulted in a significant increase in a and therefore in ~F*. Similar data have been obtained 209 , 210 for nylon 6,6 and Kevlar® and these have given values of a of about 15 and 28 respectively.

8

234

Stress (GPo)

Figure 5

Time to break vs. stres~ for polyethylene fibers: (1) Mn =2 x 104 , maximum draw temperature = 11{}-120°C (ref. 214); and (2) M w =4 x 10 6 , maximum draw temperature = 160°C (ref. 215)

To calculate ultimate tensile strength (fu) at 1 s testing time, the value WI = 1 is set in equation (34) to get the result in equation (36), where p is the density of the fiber and d 1 and A are in nm and nm 2 , respectively. If a value for d 1 of about 0.13 nm 210 is assumed and the experimental values of a for nylon 6,6, Kevlar® and ultra-high molecular weight polyethylene (15, 28 and 7.8) are used, and A is set equal to an average value of about 0.2 nm 2 , as an approximation, this gives ultimate strengths ranging from 60 dN tex - 1 for polyethylene to 140 dN tex - 1 for Kevlar®. Precision of the estimates can be improved by using more accurate values for A. However it can be seen that the estimates based on the creep failure theory are lower than those based on breaking 'primary covalent bonds but much higher than those based on extrapolations to orientation factors of unity. For polyethylene, the values of strength achieved are now close to or even above this calculation. (36)

The chain cleavage theories also lead to a prediction of the time and temperature dependence of strength similar to that of the creep failure theory 211- 213 so that the predicted effects of temperature and testing time are about the same and the calculated activation free energies are about the same.

276

Generic Polymer Systems and Applications

The problem comes in interpretation of these values. Thus the range of values of ~F* quoted are 160-190 kJmol- 1 (38-45 kcalmol- 1) for nylon 6,6, 205-240 kJmol- 1 (49-57 kcalmol- 1), for Kevlar® and 105-126kJmol- 1 (25-30kcalmol- 1) for polyethylene. These can be compared with the bond energies for these fibers (293, 335 and 340 kJ mol- 1 or 70, 80 and 83 kcal mol- 1 respectively) and it can be seen that there is no correlation. There is a better relation to the extrapolated zero strength temperatures (380, 640 and 150 °e respectively) which, in turn, are related to the properties of the crystal (i.e. intermolecular forces and crystallite size). In spite of this, it must be noted that free radicals are produced when fibers are broken and this subject has been recently reviewed. 216 The effects of molecular weight (M n) on tensile strength result from two factors. One is the role of end groups as defects and this is especially important at lower molecular weights. The second is the ability to draw to higher draw ratios as M n is increased. One way to visualize this is to consider the ratio of extended polymer chain length to the end-to-end distance of a random coil. This ratio increases as the 0.5 power of molecular weight for a freely jointed chain and perhaps the 0.4 power for a more realistic situation. This has been invoked in a discussion of the strength of polyethylene 88 ,217 and it has been proposed that draw ratio is proportional to M~·5 and demonstrated 88 that strength is approximately proportional to M~·4. Data for nylon 6,6 show a similar relationship.218 However, it must be noted that the beneficial effects of increased Mn can be washed out by processing problems such as excessive melt viscosity, increased spun fiber orientation and its variability and melt fracture. These problems can be helped by dilution with a solvent as in the case of gel-spun polyethylene although other factors have been invoked to explain the success of this approach. The tensile strength of an assembly of fibers may be considerably less than the average of its constituent fibers. For example, one calculation for a large untwisted bundle of fibers having no time dependence of tensile strength 219 shows that the ratio R of mean bundle strength to mean filament strength decreases with increasing filament strength coefficient of variation ey. Thus R = 100% for ev = 0, 87 % for ev = 5% and 670/0 for ev = 20%. When a fiber bundle is twisted, tensile strength increases with increasing twist and then reaches a maximum. Beyond this twist level, effects of helix angle and non-uniform loading of the filaments result in decreasing strength. The initial increase in strength results from frictional stress transfer between filaments so that it is possible that a filament which breaks at some distance from the location where failure will untimately occur can still bear stress at that location. The term 'ineffective length' has been used for the distance over which a broken filament is considered to have zero stress 220 and this decreases as twist is increased. The twist level is given in terms of the 'twist multiplier' (TM), which is defined as turns per inch x (denier)o.5/73 or turns per cm x (tex)o.5/9.58. For a given fiber density TM is related to the twist helix angle. For Kevlar® aramid yarns, the maximum strength occurs when TM is about 1.1 (e.g. 2.5 turns per inch for a 1000 denier yarn).209 For strong mechanical interaction between filaments, the effect of nonuniformity on strength can be even more severe than that shown above. 221 A detailed analysis of Kevlar® 49 yarns 222 has indicated that the filament strength ev averages about 12%. In general it is theoretically predicted and observed that yarn or filament strength decreases with increasing gauge length of the sample being tested. Frequently, the strength decreases with log(length). The dependence on length incr~ases with increasing strength ev. Since ey is frequently less for a twisted bundle, the length dependence will also be less. For example,209 zero twist Kevlar® 49 yarns lose strength at the rate of 20% per decade of length in the range of 4 to 40 cm. For the same yarn with TM = 1.1, the strength loss is oilly 3.5-5% per decade. The effects of stress transfer are particularly important in composites or resin-impregnated strands. For example,209 with typical hard epoxy resins reinforced by Kevlar® 49 yarn, the ineffective length is quite short and tensile strengths (calculated on fiber weight) are higher than those of optimally twisted yarns when the fiber volume content is 40% or less. As the fiber volume content is increased above 45%, fiber tensile strength decreases steadily and by 60 vol % it is 85% or less of the value for 40 vol % and is on a steep downward curve. This author believes that this is a result of increasing delamination which occurs as a result of shear forces on the relatively brittle epoxy with a resulting increase in ineffective length. This is supported by experiments in which a much lower modulus, higher elongation urethane resin is substituted for the epoxy resin. In this case, the 40 vol % strength is somewhat lower but there is little decrease in strength until fiber volume content exceeds 60%. Unfortunately, the desirable thermal and other properties of epoxy resins frequently make it impossible to take advantage of this knowledge in practice. It is assumed that the extent of decrease in strength referred to above is affected by yarn and filament uniformity. Similar improvements in strength retention of resin-impregnated strands or composites at high fiber volume contents can be effected by use of twist but thi~ is not always practical.

Polymers for Fibers 8.5.3

277

Compressive Strength

When composites reinforced with an organic fiber such as Kevlar® 49 are tested in compression, one finds that the initial modulus is about the same as that for tensile testing. 180, 209 However, at about 0.3% compressive strain, yielding begins and the ultimate compressive strength is only about 18% of the tensile failure stress for the same composite. It has been proposed that the compressive strength fe is related to Tg for the fiber 209 ,223 and an approximate correlation has been established 223 thatfe is proportional to T~·85. For example,j~ for polyethylene with Tg of 148 K is about 19 MPa, while that for Kevlar® with Tg of about 648-673 K is about 440 MPa. Thus to double the compressive strength vs. Kevlar®, according to the proposed relationship, Tg (in K) must be increased by a factor of about 1.45 to at least 940-976 K (667-703°C). Compressive strength is time and temperature dependent as is tensile strength. However, ~F* for a given fiber is lower than that for tensile failure and so is the extrapolated zero strength temperature. Thus, for Kevlar® 49/epoxy composites,209 To is 400°C or less for compression and about 640 °C for tension. The extrapolated zero stress intercept log (t b ) for compressive fatigue should thus be about 15 or 16 for Kevlar® vs. about 26-28 for tension. However, there seems to be little information available on this point. We can also estimate the fiber characteristics needed for increased compressive stress from the absolute rate theory. For example, to double the compressive stress vs. Kevlar® it would be necessary to increase To/T-1 by a factor of two, from 1.26 to 2.52. This means an increase in To from 400 °C to 776°C, a little higher than the Tg prediction described above. 8.5.4

Creep

Fibers can be considered in terms of mechanical models consisting of springs and dashpots (viscous elements). These elements are not necessarily linear in strain or strain rate but are frequently considered to be linear. A simple Hookean spring elongates instantaneously when a stress S is applied and it has a stress which is linear in strain. It recovers instantaneously when the stress is removed. When a spring is in parallel with a dashpot (the Voigt element), elongation occurs gradually at a rate which depends on the relaxation time r which is the ratio of the viscosity 11 of the dashpot to the modulus E of the spring. Equilibrium is reached when the spring reaches the extension S/E. If a spring is in series with a Voigt element, application of a stress results in instantaneous elongation followed by creep of the retarded element (both elements are under the same stress). When elongation is plotted against log(time), the curve has a sigmoid shape with a central portion which is approximately linear for perhaps 1.5 decades of time. Frequently, creep curves are linear over many decades of time and this means that a model of the fiber must contain several retarded elements and has a distribution of relaxation times. An increase in temperature can reduce the viscosity of a dashpot to such a low value that the relaxation time becomes very short with respect to the time scale of the experiment and the element then acts as an unhindered spring. The creep rate will be linear in stress and vary inversely with spring modulus for linear elements. It turns out that for some fibers, creep rate is linear in stress up to some level and then levels ofT. This has not been explained but may be the result of changes in viscoelastic behavior with increasing stress when tie chains become taut. Creep data 224 for various cordage yarns produced by Du Pont (Type 707 nylon, Type 67 Dacron®, Kevlar® 29 and Kevlar® 49), all at 1.1 TM, and measured at 50% of the nominal breaking stress for each of the yarns, show that the plots of creep vs. log(time) are linear over four decades of time, from 0.1 to 1000 hours. In this chapter, 'creep' is defined as the elongation at various times

Table 2

Creep of Cordage Yarns at 24°C, 55% RH, Stress = 50 % of Breaking Stressa Yarn

T-707 Nylon T-67 Dacron® Kelvar® 29 Kelvar® 49

(dNtex- 1 )

(dNtex- 1 )

Modulus

Creep rate (0/0 decade - 1 )

4.3 4.2 9.6 9.6

49 101 459 750

0.27 0.135 0.052 0.020

Stress

aM. H. Horn, P. G. Riewald and C. H. Zweben, Oceans '77 Con!, 1977, 24E-1.

278

Generic Polymer Systems and Applications

minus that which has occurred in 1 min. The data are summarized in Table 2. It can be seen that there is a general trend of decreasing creep rate with increasing modulus. Exact comparisons are difficult because of the differences in absolute stress levels. Other data for the Kevlar® fibers show linear semi-log creep plots for as long as five or six decades of time for stresses up to 60°A. of the breaking stress. 225 Creep data for highly oriented ultra-high molecular weight gel-spun polyethylene 86 with a modulus of 1290 dN tex -1, measured at a stress level (3.35 dN tex -1) which is 10% of the breaking stress, have been plotted and show very low creep out to more than 10 h. However, over the next 1.5 decades or so (especially between 100 and 1000 hours), the semi-log plot shows a sharp upturn with an average creep rate of 0.76% decade -1. Data for melt-spun polyethylene with much lower M n (18000) at comparable stresses 226 show much higher creep rates with a sharp upturn occurring at about 100 s and an average creep rate in the next decade of about 4% decade -1. Irradiation (52 Mrad) in vacuum to produce cross linking did not affect the initial creep rate but eliminated the upturn at 100 s so that the average creep rate over six decades was about 0.50/0 decade-I. In this work, there was no beneficial effect of the radiation on an ultra-high molecular weight gel-spun sample and this was attributed to the dominating effects of chain scission which occurs as a result of irradiation.

8.6 8.6.1

FIBER DURABILITY Introduction

In consideration of the useful life of a fabric, tire cord, rope or other structure based on fibers, we are concerned with the ability to survive mechanical stresses of various kinds (tension, compression, flex, abrasion, shear, etc.) or to resist hostile environments for the expected use life of the structure (heat, light, high energy radiation, water, solvents, etc.). To a certain extent the resistance to these stresses is an inherent property of the polymer from which the fiber is made. However, in many cases, the durability of the fibrous assembly depends on such variables as molecular weight, end group concentrations, fiber and fabric thickness, filament tex, degree of crystallinity, orientation, cord or rope construction, lubrication, etc. The following sections will attempt to deal briefly with many of these factors.

8.6.2 Tensile Fatigue In the previous discussion of tensile strength, it was noted that the time to failure increases . exponentially with decreasing stress. It was shown that the magnitudes of the effects depend on the polymer, its molecular weight and crystallinity and on the testing temperature. In most cases, the results for cyclic variations in stress are similar to those for static loading except for small factors relating to the smaller time exposure for cyclic loading. There are some indications that cyclic loading can give somewhat greater effects if the minimum stress decreases to zero or a very low value. Thus with Kevlar® filaments, 209,227 reduction in the minimum stress to zero (vs. 7 dN tex - 1) resulted in up to two decades reduction in time to break. For Kevlar® this results in at most a 7% reduction in failure stress for a given testing time. It has been speculated 210 that this may be a result of the orientation and stress gradients which exist across a filament which may result in compressive effects when the stress drops to zero. In general, the absolute rate theory for tensile failure predicts that there should be little loss in tensile strength (as measured in conventional testing) until very near the time to break. The predicted relationship is shown in equation (37), where b is the slope of the line relating log(time) to stress,!o is the initial strength (e.g. at 1 s breaking time) and It is that for some time t. Data for several fibers228.210 (nylon 6,6, PET and Rayon) show that when fibers were cycled for 80% of their estimated time to break at 25 or 120°C, the strength losses ranged from 0 to 11 % vs. a theoretical prediction of about 8% for a fiber having an initial strength of 6.6 dN tex -1 and b = 2.27. It appears that in most cases, with the possible exception of polyethylene, tensile fatigue will not be an important factor when reasonable safety factors are used. For example, at a safety factor of three, the average lifetimes predicted for nylon 6,6 (a = 15), Kevlar® (a = 26-28) and ultra-high molecular weight polyethylene (a = 7.8) in the absence of chemical degradation would be about 300 years, 10 10 years and 2 days. The values of a are the extrapolated zero stress intercepts discussed previously. The calculation for polyethylene increases to about 12 days if a = 9. Thus, for many uses we are concerned mainly with

279

Polymers for Fibers

other stresses such as abrasion, flex and various chemical processes leading to loss in molecular weight. For example, tire cords in use are generally operated at safety factors of 5-10. On the other hand, tensile failure is a major factor in the durability of ladies hosiery, but mainly in the form of single high stress events. However, for a given tensile strength, higher elongation to break results in improved wear life,218,229 presumably because of a greater probability of escape from the snagging force before failure occurs. Similarly, plunger energy tests of tires involve tensile failure due to the single penetration event.

10 - it

8.6.3

=

- (llb)log(l

-

(37)

tlt b )

Flex and Abrasion

The strength loss which occurs in operating tires is frequently a result of compression of the twisted tire cords in a limited region of the tire. The two mechanisms invoked for this strength loss are filament flexing or interply abrasion at the tire operating temperature. At this point, the chemical degradation which can also occur will not be discussed. The relative importance of flex and abrasion in tire cord fatigue depends on the twist multiplier of the twisted cord as well as on the stiffness of the tire cord as influenced by the nature and penetration of the adhesives used to bond the cord to rubber. Use of a proper twist multiplier results in minimizing of the flex contribution 209 ,228 for cords which are not too stiff. Abrasion can also be a major contributor to the loss of strength of ropes and hence use of a proper twist multiplier and suitable lubricants can provide major improvements in durability.224,230 The flex life behavior of single filaments can be determined by bending them repeatedly around a 0.076 diameter wire with some fixed stress (e.g. 0.53 dN tex -1) applied. There is a general tendency for flex life to decrease exponentially with increasing stress and increasing filament tex. The flex life of Kevlar® filaments (1.67 dtex) and nylon 6,6 filaments (6.67 dtex) of similar molecular wei'ght are about the same. 209 While the intrinsic flex resistance of nylon is thus better, this is compensated for by proper tex selection. A twofold increase in filament tex for Kevlar® results in a three- to five-fold reduction in flex life, while a doubling of the applied stress results in a similar decrease. The abrasion resistance of typical melt-spun fibers is strongly affected by molecular weight. Cycles to failure in several standard abrasion tests approximately doubles as the number average chain length increases from 80 to 120 nm. 218 Abrasion resistance can decrease for yarns having a very high orientation level.209,218 One way to view these effects is that sensitivity to abrasion depends on the number of interfibrillar tie chains. These would be expected to increase with increasing molecular weight and decrease with increasing orientation or chain stiffness. Even for the aramid fibers, higher modulus tire cords (via increased hot stretching tensions) have poorer fatigue resistance. In tire cords having a high enough twist multiplier such that abrasion is the dominant factor, the rate of strength loss increases directly with increasing cord stress and decreases rapidly with increasing twist multiplier. 209 ,224,228 For example, in the 'disk fatigue' test with 7 TM (25 helix angle) Kevlar® tire cords, strength loss at a fixed compression setting (3.5 0/0) increased linearly with cord stress from about 20/0 at 0.25 dN tex - 1 to 50% at 0.9 dN tex - 1. This degree of compression is exaggerated vs. compressions wl)ich occur in radial tire belts. Similarly, with ropes from resinimpregnated Kevlar® strands containing a wax lubricant, abrasion cycles to failure in a reverse bending over sheave test increased exponentially with increasing safety factor (e.g. from 2000 cycles at a SF of two to 500000 cycles at SF of four). For most fibers, abrasion resistance decreases with increasing temperature, at least in part a result of the increasing coefficient of friction. Abrasion resistance is frequently poorer for wet yarns or strands than dry,230 cycles to failure decreasing by as much as an order of magnitude. Of course, abrasion can be modified with the use of proper lubrication, if feasible, as has been mentioned above. Thus for non-impregnated poly(p-phenyleneterephthalamide) yarns 230 the use of various paraffin waxes (e.g. 3-60/0) effected major improvements in the strength loss of a 1667 dtex yarn abraded wet for 10000 cycles under a load of 3.5 dN tex -1. The bare yarn failed in 350 cycles, while the lubricated yarns showed 5-25% strength loss after 10000 cycles, depending on the level and type of lubricant. 0

8.6.4 Degradation Environmental factors which cause a reduction in molecular weight have a sensitive effect on tensile properties and durability. It has been mentioned previously that tensile strength increases as about the 0.4-0.5 power of molecular weight (M n) for moderate to high molecular weights. Thus a

280

Generic Polymer Systems and Applications

200/0 change in molecular weight should cause a ca. 100/0 change in tensile strength, in the absence of complicating uniformity factors. However, when the change in M n occurs as a result of degradation of an oriented fiber, the effects are much more sensitive 21 0,218 so that a 20% loss in AI n often causes a 50% loss in strength. Most fibers lose strength when exposed to air at high temperatures below their melting point. However, they differ greatly in their sensitivity to oxidation. Thus unprotected aliphatic polyamides such as nylon 6,6 lose strength rather rapidly when heated in air at 180°C (20% strength retention after 24 h). However, when the polymer contains Cu 2+ -based antioxidants (e.g. 60 p.p.m. Cu 2+), the degradation rate is reduced by a factor of 40 or more 218 so that strength retention is about 85%. Degradation rate is judged by the linear plot of reciprocal of strength retained vs. (time)0.5 for this reaction, which is oxygen diffusion controlled. PET fibers have similarly good stability without antioxidants. The sensitivity of the unprotected nylons stems from the relatively easy abstractability by peroxy and other radicals of the hydrogen atom adjacent to the amide nitrogen. This helps propagate a chain degradation reaction. Aramid fibers also have high oxidation stability since this type of readily abstractable hydrogen atom is not present. Thus Kevlar® aramid tire yarns retain > 85% of their strength after 24 h at 180°C, while Nomex® aramid yarns (m-aramid) lose essentially no strength in this time. 199 Photo-oxidation (UV radiation) is also a potential source of fiber strength loss. In this case, the rates are normally much slower than high temperature thermal oxidation so that the rate is normally not controlled by oxygen diffusion and the reciprocal of strength retained is linear with time. The labile hydrogen atoms of the aliphatic polyamides are still an important factor in photo-oxidation and major improvements are possible by incorporation of Cu 2+ -based antioxidants. However, absence of these sensitive points does not insure good photostability since other pathways for degradation exist. Thus the p-aramids, which are yellow and have an absorption maximum at about 330 nm tailing out.to about 410 nm 231 form free radicals when UV light is absorbed. In the absence of oxygen, these rearrange to form a carbonyl group without breaking the chain,232-234 the 'photoFries' reaction, and degradation rate is slow. However, when oxygen is present, it reacts with the free radicals to break the chain. Since the absorption coefficient is so high, the degradation rate for Kevlar® filaments or yarns is about the same as for unprotected nylon. In spite of this, when larger assemblies are used, the high degree of self-screening can result in much higher stability. For example, a fabric of Kevlar® 29 lost 50% of its strength after five weeks exposure to sunlight in Florida but a half inch thick rope only lost 10% in six months. 180 The relative stability of various fibers can depend strongly on the wavelength distribution of the incident light. This is particularly true when the absorption bands are further out in the UV than they are for the p-aramids. Thus, in outdoor Florida sunlight, bright nylon 6,6 (i.e. containing no Ti0 2, a photoinitiator) is superior to PET. However, when the samples are exposed under glass, PET is superior because of screening out of wavelengths to which the polyester is sensitive. 235 Polypropylene fibers are very unstable to UV exposure so that they lose a great deal of strength in times as short as 25 h. However, hindered amine stabilizers are available which can decrease degradation rate by an order of magnitude by formation of nitroxyl radicals, which can scavenge the peroxy radicals responsible for the chain degradation reaction.236-239 The degradation rates for fibers exposed to large doses of ionizing radiation do not necessarily parallel those for photodegradation. Thus, for irradiation in air, nylon 6,6 yarns lost 80% of their strength when subjected to 200 Mrad of2 MeV electron irradiation, while the aramid fibers lost only 10-25% of their strength for 600 Mrad irradiation. 16, 210 The yield of free radicals and other products (e.g. H 2) is very low for the fully aromatic polymers, about 1% of those for aliphatic polymers 16 because of de-excitation pathways provided by the aromatic rings. Cross linking can also occur with radiation but the extent is normally not sufficient to cause significant modulus or recovery changes in typical semicrystalline polymers unless some type of monomer is also included to magnify the effects of the radiation. Hydrolysis can be a source of strength loss in certain end uses. For example, polyester tire cords can lose strength as a result of aminolysis and hydrolysis in tires which run too hot. However, use of low amine rubber stocks and fibers which have reduced levels of -C0 2H end groups (catalyst for both reactions), along with cooler running rubber stocks, have controlled this problem. 228,240 The aliphatic polyamides generally have good resistance to hydrolytic degradation in the pH range of 5-13.,218 but are subject to degradation, e.g. in boiling water, at lower pH values because of the affinity of acids which hydrogen bond to the amide groups. On the other hand, inorganic bases are not readily solvated so that degradation is not a problem at higher pH. For the aramid fibers, hydrolytic stability is good over the pH range of 4_8 231 with an estimated time to 500/0 strength loss of about 180 days at 100°C. PET fibers are more sensitive to degradation by inorganic bases in spite

Polymers for Fibers

281

of a low affinity for them. In this case, degradation proceeds by progressive hydrolysis and polymer removal from the surface with little loss of molecular weight in the interior of the fibers. In the case of cationically dyeable PET fibers, which contain acidic functional groups (e.g. S03X), a low pH environment (e.g. during dyeing) converts these to free sulfonic acid groups, which, at elevated temperatures (e.g. 100°C), promote hydrolysis to reduce both molecular weight and strength. 218 This problem can be alleviated by addition of Na 2 S0 4 (e.g. 0.2%) to the dye bath, which shifts the critical pH for good strength retention by almost two units and makes it possible to dye without significant strength loss. 8.7

REFERENCES

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