Indian Journal of Fibre & Textile Research Vol. 16, March 199'1, pp. 65-72

Fibres from polymer blends

B L Deopura, K N Bhaumik & S Mahajan

Department of Textile Technology, Indian Institute of Technology, New Delhi lID 016, India

Received 14 November 1990

Specific improvements in fibre properties can be achieved if they are produced from polymer blends. Fibre properties are highly dependent on phase dispersion of the component polymers. In practice, two polymers are combined to give either biconstituent or bicomponent fibres . Biconstituent fibres are easier to produce and their production and properties are described and discussed in this paper, in particular of those based on the. following combinations: blends of nylon 6 with nylon 66 and other nylons, polypropylene (PP) with polystyrene, fibre-grade PP with plastic-grade PP, poly(ethylene terephthalate) (PET) with block copolyester, and flexible polymers with rigid rod copolyester.

Keywords: Compatibility, Polyamides, Poly(ethylene terephthalate), Polymer blends, Polypropylene

1 Introduction Blending of polymers is a versatile method for

obtaining new materials with improved properties. Polymer blend has been the subject of several recent reviews l - 3 and monographs4 •5 . Polymer blending has also been used to achieve specific improvements in the properties of synthetic fibres6 •7 • Fibres prepared by intimately mixing two polymers are generally called polyblend (or alloy) fibres. An overview of polyblend fibres, including a few typical examples describing different aspects, has been presented in this paper.

2 Polymer Blends 2.1 General Considerations

Polymer blends are simple but intimate physical mixtures of the constituent polymers with negligible covalent bonds occurring between them. Polymers may be miscible or immiscible, depending on the thermodynamics of polymer-polymer interactions and the kinetics of the mixing process. In general, polymer blends are immiscible and form phase-separated systems. These systems are usually opaque and show two glass transition temperatures I - 5. There are some examples of miscible polymer pairs; they show single glass transition temperature and form a single phase. A completely miscible polyblend tends to have properties intermediate between the two polymers comprising the poly blend, whereas a slight incompatibility often provides a superior balance of useful properties.

2.2 Compatibility (Miscibility) of Polymers

Compatible polymer pairs are generally described as polymer mixtures that have desirable properties when blended. Compatibility is the term used to cover a wide range of situations to describe good adhesion, useful and improved properties, and ease of blendingB

•9 for the components. All these are

dependent on the extent of miscibility of polymers of interest. The mixing exhibited at the molecular scale is expressed as 'miscibility'.

The mixing of the two components takes place in accordance with the thermodynamic equation 10:

~GM=~HM - T~SM ... (I)

where ~GM is the change in Gibbs' free energy; ~HM' the enthalpy change; ~SM ' the chaqge in entropy; and T, the absolute temperature.

The thermodynamic requirement for the miscibility of two polymers is that the free energy of mixing ~GM must be negative. For polymeric systems, it is found that ~SM is quite small. This is a consequence of the very small combination entropy change (~SM) which results when two high molecular weight polymers a re blended. The relation of ~SM to the degree of polymerization of the two blended polymers is given by the Flory-H uggins lattice model for polymer solutions II as applied by Scott 12 to polymer mixtures:

~SM =(RV/ ~r)[(<p,ln <p, /x,)+(<p2In <P2/X2)] . .. (2)

where R is the ideal gas constant; V, the total volume of two polymers; Vr, a reference volume which is taken as close to the molar volume; <p, and <P2, the

65

INDIAN J. FIBRE TEXT. RES. , MARCH 199 1

volume fract ions of polymers; XI and Xl, the degree of polymerizat ion of polymers in terms of the reference vol ume Vr . Typica l va lues for the entropic term, TI'1SM , in Eq. (1) for high polymer blends may be less than 0.005 cal/gu . In order that I'1GM be negati ve, the small negative en tropic term cannot be dominated by a large positi ve contribution from the enthalpy of mi xing, I'1 HM . I n other words, polymer mi scibility is limited to those blends with an exothermic or mildl y endot hermic heat of mi xing. This implies that blend mi scibilit y req uires favo urable interactions between the two pol ymers in the blend. Such interact ions, as hyd rogen bonding. di pole-dipole, orcha rge transfer complexation tha t can lead to a favourab le energetic state, are apparen tl y rare among hi gh polymers and only a few cases of ble nd mi sc ibi li ty ha ve been demonstrated.

The dependence ofthermodynam ie miscibility of a polymer pair on both com posit ion and temperature may be represented by a liquid- liquid phase diagram like the one shown in Fig. 1. At temperature T I , which is be low the upper cri tica l solu tion temperatu re (UeST) for phase separation located at T1, the eq uilibrium mixture wi ll separate into two phases whose compositions lie on opposite sides of the binodal at T I • The binoda l separates the stable (single phase) from the metastable state while the spi nodal

Ts Two phases

QI Tio '-::J

+- T3 d '-

One phose QI a.

Tz E QI l-

T, Two phases

o Fig. I- Idea lized liquid-liquid phase diagram fo r a po lymer blend showing both a n upper and a lower critical solu tion tempera tu re T, and "/ -. resp.:ctivcly. The rmodynamically unstab le regio ns are c~ntaincd within the spi nodal (-----). The bi nod al ( ) dema rks th~ bounda ry between metas table a nd sta ble

(single-phase) regions

66

marks the transition from the unstable to metastable region. At T3 , which is above the UCST but below the lower critical solution temperature (LCST) located at T4 , the blend is a thermodynamically mi scible mixture at all compositions. Above the LCST, e.g. at Ts, two phases aga in co-exist with compositions indica ted by the upper binodal.

Evidence for the existence of a LeST in severa l polymer blends has been cited recen tll 4 - 16 . By compari son, there, is little and usually only indirect evidence for a UCST in polymer blends I6 - 18 .

Inabi lity to detect a UeST for a given polymer blend may resul t from a phase change such as crystallization of one of the components for the melt or vitrifica tion of the amorphous mixture at temperatures above the UeST of the blend 19. Such a change in sta te would provide a diffusional barrier for the attain'ment of thermodynamic eq uilibrium and thereby prevent direct experimenta l observa ti on of any low tempra tu re binoda l. A comprehensive treatment of phase behaviour in polymer blends has been given by Kwei and Wang20.

Fibres are usuall y prepared from crystallizable polymers. Due to crystallization process, phase segregati on takes place. This is further enhanced by the ori entat ion of the molecular chains, leading to significant increase in crysta lliza ti on rates. The two components may, however, be mi sci ble in the amorphous regions.

Partia ll y mi scible systems a llow good di spersion and the formation of an interphase between the polymers with possible improvement in the properties. Immiscible POl Y?lll'r ~y~!CJ1 ~ <; have poorly defin ed interphases which can he rlit (.' )lood use by splitting them for the prod ucti nn of ult rafi ne fibres for synthetic paper, artificia l leather. etc.

2.2.1 Characteriza tion of Miscibilit~

There is a grea t need to quantify miscibility or interaction parameters by experimental means for a blend system so that the structure-property rel a tionships can be developed more completely. There a re various experimental methods for measuri ng miscibility either quantitatively or qua lit atively.

The most common methods a re those relating to the glass transition temperature (Tg) and the depression of melting point (Till)' In the case of incompatible blends, each phase maintains the properties of the pure homopolymer, e.g. Tg of each polymer remai ns the same showing two Tg for the blend system. I I' mi sci bility occurs to some extent, the T~ shifts. For highl y mi scible blend systems, a single T:, intermed iate between those of the pure polymers,

DEOPURA er af.: FIBRES FROM POLYMER BLENDS

is observed . Tg can be predicted by the Fox equation I:

I/ Tgm = (wl / Tg,) + (w2/ Tg,) . . . (3)

where IVI and W 2 are the weight fractions of the constituents: Tg, and Tg" the glass transition temperature, of the constituents; and Tgm' the glass transition temperature of the mixture.

Several techniques such as differential scanning calorimetry21 (OSC), dynamic mechanical analysis22 - 24 (OM A), dilatomet ry25 and dielectric relaxation26.27 have been used for measurement of Tg .

The melting point depression of crystallizable components can also be correlated with miscibility28 - 31. However, this approach is not free from problems. The melting point measured directly is influenced by the perfection and size of the crystals which are influenced by the condition under which they are grown.

I n the case of immisci ble systems, the size and shape of the inclusions, which are predominantly related to the mechanical properties, can be determined by light and electron microscopy, wide- and small-angle X-ray scattering, etc. Several other techniques to measure miscibility or polymer-polymer inreractions are available and these include heats of dilution 32 , dircet calorimetric studies on low molecular weight analogs33 , neutron scattering34.35 , sorption probes36.37 , FTIR38.J9 , NMR 36.40, etc.

2.3 Phase Morphology and its Control

In the case of polyblend fibres, the phase morphology of a blend is an important factor in the macroscopic behaviour of mixture. Many properties and subsequent use of a poly blend fibre depend critically on the nature of the arrangement of the two phases. The number of potential phase arrangements is almost unlimited . The poss ible morphologies could, howevyr, be considered under two broad divisions:

(I) Structured, i.e. Bicomponent Fibres: Fibres under this group consist of two components divided into two relatively distinct regions within t-he cross-section along the length of the fibre. Typical examples are side-by-side, sheath-core and their variations. Their production requires special fibre formation equipment. A comprehensive account of bicomponent fibres appears in the Iiterature6.7 .

(2) Random, i.e. Biconstituent Fibres: This type of polyblend fibres is composed of a more intimate blend in which one of the components appears in the form of small, discrete phase embedded in a

more or less continuous matrix of the second component. T ypical examples are matrix-fibril (M / F type) and interpenetrating network (IPN type). The dispersed phase is mainly in the form of fibrils. The volume fraction , size and distribution of the di spersed phase and its adhesion to the matrix control the properties of the fibre.

In M j F type fibres, the ma in factors which control the phase morphology are the intensity of mixing, interfacial tension , relative rheological characteristics i.e. mainly elasticity, which, in turn , depends on molecular weight (MW) and molecular weight distribution (MWO), shear stress and phase deformation in processing equipment41 .42 The relevant parameters governing the deformation of fluid droplets of constituent polymers are the viscosity ratio of the two components and the ratio of the interfacial tension and shear stress4J.44. Mixing of the two polymers is better in the case of similar viscosities of the two components under the processing conditions used. Several investigators have examined ma ny other factors influencing the size and distribution of fibrils in M j F type fibre s45 - 48.

Recent studies on partly miscible melt-spun blends have suggested that crystallization of one of the polymers is affected by the presence of the other polymer45.49.5o. Due to change in the crystallization kinetics in the presence of the other polymer, the morphoiogy and the level offinal crystallinity of the blend are affected 5 I.

There is frequently an intermediate zone (Fig. 2) called inte rpenetrating network (lPN), where each phase remains continuously connected throughout the bulk of the blend. Such a structure is formed when the two polymers are in nearly equal proportion in the blend . They may be generated by phase separation in pa rtiall y mi scible systems through the mechanism of spinoda l decomposition 52 or in totally immiscible systems thro ugh the selection of appropriate process ing and rheological conditions53 and by incorporation of certain additives54. This has an important consequence in tha t each component now shares the load beca use of the network structure. As a result , there is improvement in mechanical properties even if the degree of adhesion is not too good 55 . IPN structures offer unique opportunities for blend product development.

2.4 Mechanical Properties of Polymer Blends

A compatible polybJend generally exhibits mechanical properties in proportion to the ratio of the componel1ts56 and shows three types of behavi o ur, VIZ. synergistic, additive and

67

INDIAN 1. FIBRE TE XT. RES., MARCH 199 1

VI ::J

::J "'0 o E

E,

o

Para!lel

0·2 0·4 0·6 0·8 1·0 Volume fraction

Fig. 2- Macha nical modulus of va rio us polymer-polymer phase a rra ngements. T he cube shows unit cell of a n idea lized inte rpenetra ting netwo rk (IPN ) structu re. T he shaded a rea is detem1ined by upper and lower bo und estimates fo r this model

>­-L-OJ a. o L-

a..

A (omposition B Fig. 3- Possi ble pat terns of prope rty va ri ation with blend

composit ion. Shaded regiu n deno tes addi ti ve behavio ur

incompatible , as shown in Fig. 3. The modulus genera ll y shows weighted ave rage behaviour. The o ther impo rtant pro perties like strength a nd toughness may not fo ll ow a n additi ve re la tio nship. Thi s usua lly a ri ses due to poor degree o f inte rfac ia l ad hes io n between compo nents tha t provide a multiplicity o f defec ts fo r ea rly fa ilure 55

.

Schre iber a nd R OSS 57 a nd Ka rasz 58 se lec ted the tenacity o f the spun fibre as a n index o f useful pro perty develo pment a nd compa tibility. In principl e, property enha ncement sho uld depend o n

68

the ability o f the constituent ma te ri a ls to remain in mo lecul a r co ntact when the complex structure is subjected to fo rces. In compa tible po lyblends, the mecha nica l behavio ur is a ffected by the properties of the indi vidua l constituents, mode of di spersion, mo rpho logy a nd compatibility in the amo rphous sta te , degree o f crysta llinity a nd sta te of orienta ti o n 59. 60 .

For immi scible blends, the presence of one component a ffects the crys ta lli za tion kinetics and mo rpho logy o f the o ther, which has been proposed as an expl a na ti o n fo r the mecha nica l behavio ur61 .6z •

Provided the res ulta nt phase mo rpho logy can be .defin ed th rough va rio us measurements, o ne migh t expec t tha t pro perties assot ia ted with sma ll defo rma tio ns, such as modulus, could be predicted by composite theory a nd' models5. 56 .6 J.64.

Broadly spea king, the re a re two compet ing mo rph o logical facto rs ex pected to a ffec t the tensile pro perti es: (a) de te riora ti o n of pro perties as a resul t of incompa tibility a nd consequent two-ph ase structure, a nd (b) improvement in properti es as a result o f the reducti on in the ave rage fibri lla r size due to the presence o f the seco nd component, increase in ove ra ll c rysta llini ty, a nd formatio n o f a la rger number of interc rysta lline links65 .

3 Production of Polyblend Fibres The production process used fo r polyblend fi bre

of the bico nstituent type is essentia lly simila r to tha t used fo r producing single polymer fibres.

3.1 Spinning

Spinning invo lves three impo rta nt opera ti o ns, viz. the prepa ra ti o n o f the spinning fluid , its t ra nspo rt to the spinnere tte a nd fi na ll y fi b re fo rma ti o n in the vicinity o f the spinnere tte hole.

For melt spinning, the two polymers a re usua ll y dry blended in gran ula r fo rm , melted , intensive ly mi xed in a n ex truder a nd spu n into fi bre. I n so me cases, the po lymers may be precompo unded in a n ex truder to ensure intimate mi xing a nd pelleti zed . The pe llets a re then used to spin the fi bre. In melt spinning o f condensa ti o n po lymers, ce rta in inte rcha in react io ns such as tra nsamida ti o n, tra nseste ri ficat io n a nd tra nspo lyesteramida tio n , may occur66

.67

. Fo r soluti o n spinning (wet o r dry), the two polymers a re usua ll y dry blended and then di ssolved in the solvent. The two po lymers may be di ssolved sepa ra tely a nd the soluti o ns then mi xed , o r o ne po lymer may be di ssolved a nd mi xed with the mo no mer of the o ther po lymer which is subseq uently po lymerized and spun into the fi bre. The stability of the blend fluid is a n impo rtant facto r. In a solutio n, the components of the blend mi ght be sepa ra ted by coalescence68 . 0 9

DEOPURA et af.: FIBRES FROM POLYMER BLENDS

After the mixing step, the blend spinning fluid is delivered by laminar flow under relatively low stress to the spinnerette. During this transport , small particles of either polymer may coalesce into larger ones due to the relatively low stress conditions. This is not desirable . One 'way to avoid this is to add some form of in-line mixing system before the spinnerette; static mixers (e.g. Kenics mixer) have been used for this purpose by several researchers7o.71. At the hole entrance, a rapid acceleration begins, which produces an elongational flow that deforms dispersed particles into fibrils . The size (L/ D ratio) and number offibrils in the matrix are related to the relative rheological characteristics of the two phases44, the interfacial tension 72 and the hole geomet ry 73 . Small relative interfacial tension and viscosity or elasticity of the dispersed phase to the matrix produces fibrils with large L/D ratios. There is an optimum die entry angle depending upon rheological behaviour for the smallest diameter and greatest number of fibrils in the matrix.

After emerging from the spinnerette hole other factors such as the draw down ratio 74- 76 and the rate of solidification (crystallization, coagulation, etc)? 7 become important for controlling and preserving the fibril s th at have been formed. In melt spinning, the amount of draw down ratios is high in certain cases and may have a large effect of fibril geometry75. In wet spinning, the draw down ratio is often less and consequently elongational flow in the spinnerelte hole must be the dominant mechanism for fibril formation .

3.2 Drawing

In blend fibre drawing under given conditions, each polymer phase will have a maximum draw ratio that can be achieved without breaks and is predominantly affected by its Tg,molecular weight and specific interactions between the molecules. Generally, the diameter of the fibril s in M/ F fibres decreases with increase in draw ratio 78 .79. The length of fibrils increases provided there is no breakage of fibrils into shorter fragments . In the absence of interaction between the two polymers, early damage may occur due to phase separation. I n the presence of interaction between the two phases, the i11aximum achievable draw ratio may exceed that of either pure polymer50. The opposite situation exists as well 78 .

This generally occurs in the case of rigid polymers in the form of inclusions. It is frequently observed in M /F fibres that drawing induces voids in the blend fibre because one component is drawn more than the other and there IS inadequate interfacial adhesion45 .50.80.

4 Examples of Polyblend Fibres 4.1 Fibres from Blends of Nylon 6 and Nylon 66

A detailed study on fibres from nylon 6/nylon 66 blends 50.81 - 83 in a wide range of compositions has been reported. Thermal studies show that % crystallinity and heat of fusion for blends are lower than those for corresponding homopoiymers. The major component at low fractions does not crystallize at all.

Fibres from blends with nylon 6 as major component ( > 90%) crystallize in the y-form as compared to predominantly ex-form for un blended nylon 6 fibre 82 . When the major component is nylon 66 (> 90%), the blend has a crystalline fraction lower by 10% of that for unblended nylon 66. It is observed that nylon 66 acts as a nucleating agent for nylon 6. The overall crystallization rate of either nylon decreases with blend composition and is minimum in the range of 30-50% of nylon 66 in nylon 6 (ref. 83).

During the melt spinning of blends, sufficient uniformity is obtained, indicating good mixing of components. The lower crystalline content, increased crystal size distribution and/or the presence of the metastable y-crystalline form in nylon 6 spun samples, and the associated interfacial sli ppage at the boundary between components during drawing enhance the drawability of the spun samples and result in an improvement in properties. The polyblend fibre has teqacity of 1.0 GPa and modulus of 7.0 GPa83 . These improvements in properties for blend filaments are attributed to increased birefringence, greater amorphous orientation and smaller fibrillar size. The formulation of an interconnected intercrystalline fibrillar network with fewer defects in terms of fibrillar ends may also make a contribution . The blend fibres show higher storage modulus.TMA shrinkage decreases with blending of nylon 66 in nylon 6. The blends of nylon 6/nylon 66 show good potential for producing technical-grade filaments, e.g. tyre and V-belt cords, yarns for geotextiles, etc.

4.2 Fibres from Blends of Nylon 6 or Nylon 66 with Other

Polyamides

Fibres produced from blends of nylon 6 with nylon 610, nylon II , and nylon 12 have been investigated by Kiato et al. 50. A high tenacity (1.2 GPa) filament is obtained from 50/50 nylon 6/nylon 610 blend . The marked increase in tensile strength and modulus of this sample is believed to be related to the highly interconnected interfibrillar network system.

A blend fibre from 80/20 mixture of nylon 66 and poly(hexamethylene isophthalamide) or poly(m-phenylene adipamide) shows improved Tg

69

INDIAN 1. FIBRE TEXT. RES., MARCH 1991

and modulus66 .B4

. Both Tg and modulus are higher than that expected from additivity based on a simple rule of mixtures. Since the fibre consists of two polyamides, a homogeneous blend is formed. It is also likely that some transamidation might have occurred, forming a block copolymer in the blend66 .

This polyblend fibre has improved nonftatspotting properties.

4.3 Fibres from PET and a Block Copoly (ester-etherT(PEE)

Fibres produced from blends of PET and a block copoly(ester-ether) (PEE) have been studied4 7 .

These types of blend form partial compatible systems. Fibres are produced by melt extrusion and then hot-drawn several times in water. It is observed that PET fibrils are dispersed in the matrix of PEE wi th considera ble en tanglemen t. The microstruct ure of the annealed blended fibre has a close resemblance to that of natural wool except that the sizes of the fibrils are large. The tenacity, tensile modulus and elongation-to-break of these fibres are 0.39-0.49 GPa, 2.9-4.9 GPa and 35-50% respectively. These values are significantly higher than those of the conventional wool fibres having corresponding values 0.12-0.2 GPa, 1.3-2.9 GPa and 25-35% respectively. In addition, the elastic recovery of these blended fibres is almost equal to or greater than that of na tura l wool.

4.4 Fibres from Blends of PP with PS, PET, Nylon 66 and

Copolymer

Fibres produced from PP with a small percentage of atactic PS show promise in terms of improved texturizability and dyeabilityB5. As-spun fibres are found to have substantially low crystallinity (34%) for 5% PS blend samples; this is a prerequisite for feeder yarn to get satisfactory crimp rigidity . TtXtured yarns have low crystallinity, leading to easy dyeability with di sperse dyes. It is likely that the reduction in crystallinity of PP is partly due to the partial compatibility ofPS with PP in the melt and in the fibre.

Shimizu ef a1. 86 .S7 have shown that it is possible to produce ultrafine fibres by melt spinning technique. They selected a PP IPS system where the disperse phase can be manipulated by varying the shear viSCOSllies of the components through the appropriate choice of spinning temperature . When the volume ofPS exceeds 60%, PS forms the matrix phase above 230°C, as the shear viscosity of PP is higher than that of PS. The dispersed phase (PP) is deformed into ultrafine fibrils by attenuating the matrix phase (PS) during spinning. The dispersed fibrils are made much thinner by increasing the take-up speed. After spinning the fine PP,· fibriis

70

distributed in the matrix are separated by dissolving the matrix polymer. For 20/80 (PP/ PS) blend , the average diameter of the ultrafine fibrils obtained at ~ake-up speeds of 500 and 3500 m/min are 550 and 400 A respectively and the average > fibril lengths are estimated to be 2 a nd 4 cm respect ively. Thus, the ultrafine fibrils obtained in this manner have aspect ratios reaching 105_106 .

The blend fibres from PP/PET are used for easy therma l-bondable nonwovens comprising melt blown fibres averaging less than 10 ~m in diameterss.s9. At the spinning conditions, PET remains as a core while PP forms a sheath due to the low viscosity and solidification temperature of PP. When the melt blown web is them1al bonded, the PET in the core crystallizes while the PP component bonds at the point of intersection . The PET serves as a supporting network so th at the web retains its porosity and fibrous nature. The fibres become bonded together to give a web. The polymer components are generally used in amounts between 40 and 60% volume, but can vary outside this range.

Recent work 90 on matrix/fibril type fibres from PP (30-70%) and nylon 66 (70-30%) being done at lCl (UK) involves melt spinning of blends. The fibres so produced have fibrils in the structure, arranged along the axis and randomly joined with interconnections that penetrate through the fibrils of the other component so that the resultant structure of both the components exists in the fibre as interpenetrating network (IPN). Thi s requires a high degree of shear during spinning and is mainly controlled by the passage through the spinnerette of appropriate hole size and shape.

An aminoalkyl acrylate-ethylene copolymer (8%) has been blended with PP and melt spun to produce a staple fibre· with improved dyea bility9 1.

4.5 Fibres from Blends of Different Molecular Weights of the Same

Polymer

There are a few reports on studies of blends using different molecular weights of the same polymer92

Hinrichsen and Green93 studied the rheological behaviour of blends of two diffe rent molecular weights of nylon 6 and observed that it is possible to melt spin high molecular weight nylons. Oeopura ef al. 94 .95 have studied the crystallization behaviour of PP and high molecular weight PP (HMPP) blend system of various compositions. In the case of 3% HMPPcomposi tion , increased crystallization rate is observed as a result of enhanced nuclea tion and growth rate than in PP sample. However , at higher HM PP compositions, the growth rate decreases. The breaking stress and initial modulus of PP/ HMPP (6%) blend fibre are 0.75 GPa and 7.34 GPa as

DEOPURA et al.: FIBRES FROM POLYMER BLENDS

compared to 0.65 GPa and 5.2 GPa respectively for PP fibre. The improved mechanical properties are attributed to increased amorphous orientation and in terconnected-i n tercrystalli ne tic molecu les. Heat-setting of these samples shows decrease in amorphous orientation; however , the rate of decrease is found to be less in the case of6% HMPP blend samples compared to PP samples.

4.6 Fibres from Blends of Flexible and Rigid Rod Polymers

The blends of liquid crystalline polymers (LCP) with conventional thermoplastic polymers have been studied in recent years96 - 102 and it is observed that the blends of' thermotropic copolyesters (e.g. vectran, Xydar , X7G, etc.) with thermoplastics (e.g. PET, polyimides, polycarbonate, polysulphone, etc .) have ma ny advantages, such as the ' direct use of conventional melt processing and ease of processability due to reduction in blend viscosity. Most LCPs have an intrinsically low melt viscosity in the mesophase and a tendency towards easy orientation in the flow direction99 .

Melt-spun fibres from blends ofPHB-PET (LCP) and polycarbonate (PC) show improvements in the mechanical properties 1 00. Both the tensile strength and modulus increase linearly whereas the elongation a t break decreases sharply. Fibres produced from a blend of5% HNA-HBA (LCP) and PC show a modulus improvement over that of melt-spun PC fibres 10 1. Blend fibres with 10% 60-PHB-PET and 90% PET exhibit significant increase in modulus from 1.9 GPa to 2.5 GPa and tensile strength from 60.5 MPa to 67.8 MPa 102.

Oeopura el al. 1 03 reported a significant improvement in the modulus and strength of90/l0 PET/PHB-PET blended melt drawn fibre; the blended fibre had a modulus of 17.72 GPa a nd a strength of 1.01 GPa as compared to 1l.75 GPa and 0.78 GPa respectively for PET fibre. Simjlar improvements have been observed for fibres from blends of PBT and PHB-PET. The mechanica l properties of nylon 6 fibres are improved by the addition of a small amount of a LC copolyestcramide 104

. Thus, it is possible to reduce the high cost of'LC P by blending with low cost polymers and impa rt unique properties such as dyeability a nd hyd rolytic res istance to the resulting fibre.

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