Indian Jo urna l of Fibre & Textile Researc h
Vol. 16, Ma rch 199 1, pp. 100-115
Nature of crystals and their influence on fibre properties a
V B Gupta & J Radh akrish nan
Department of Textile Technology, Ind ian Institut e o f Technology.
ew Delhi 11 00 16, India
Received 15 J an ua ry 199 1
Some impo rt a nt structura l and mo rph o logica l as pec ts o f crys ta ls in po lymeric fibres a nd thei r inf"l uence
on the mecha nica l, thennal and optica l properties of the crysta l a nd the fibre are considered. These include the
role o f mo lecu lar arch itect ure, crysta l defects, nat ure of crysta l-a morp ho us coupli ng, crys ta l content , crystal
size and crysta l o rientation o n the stiffness, strength , melting point. therma l expansio n, birefringence. etc. o f
the c rys ta l and the fib re. Some important aspects of crysta l forma ti o n a re also discu ssed . The emphasis
th ro ugho ut is o n poly(c th yle ne te rephth a la te) fib res.
Keywords: Ax ial modulus, Ax ial strength, Birefringence, C rysta l defec ts, Entanglement. Molec ular architecture. r
M o lecu lar co nfo rma ti on. Polymer crys ta l
Introduction
The most successful po lymeri c fibre s, both na tura l
and ma nufactured , are semi-crysta lline in which th e
c rysta ls are rela tively sma ll in size (2-20 nm) a nd a re,
therefo re, often refe rred to as c rys ta llites. Th ey ma y
fo rm a nywhere betwee n 20% a nd 90 % o f the to ta l
mass o f the fibre. Th e crys tals a re known to ma ke a
predom ina nt co ntributi o n to th e macroscopic
stiffness, strength , durabi lity, therma l resistance a nd
sta bility of the fibre. Howeve r, the effec ts of th e fine
structure of the crystals (mol ec ular architecture a nd
o rientati o n in crystals, lattice defects) and their
m o rph o logy (crysta l size, crysta l-a morph o us
co upling) o n the mac rosco pi c prope rti es o f the fibre
are not so well understood . The present contributi on
is a n a ttempt to prese nt a n integra ted , thou gh bri ef,
acco unt of the nature a nd role of crystals in
po lymer-ba sed fi bres with the emph asis being placed
o n po ly(e th ylene terephtha late) (P ET ) fibres. The
fo rma tion of crysta ls is of obv io us impo rta nce in such
a di sc ussion a nd thi s aspect has a lso been brie fl y
considered with particular emph asis o n some
deve lopments in the current thinkin g on thi s
subject.
2 Nature and Role of Molecular Architecture
The nature a nd ro le o f mo lec ul a r architecture will
first be considered with reference ·to axial modulu s
"A condensed ve rsion of this pape r was presen ted a t Polymers ' 9 1,
an in te rn a ti o na l co nference o n Po lymer Science: Co nt empo ra ry
Themes, held at Na tio na l Chemica l Laboratory, Pune.lndia, 1-4
Ja n. 199 1.
100
a nd ax ia l strength of th e crystal a nd th e fibre
respectivel y and la ter a bri ef reference wi ll be made to
thei r effects on some the rmal pro perties, viz. th erm a l
ex pa nsio n a nd meltin g behaviour. These aspects
have acq uired signi fica nce as a result o f the rece nt
deve lo pment s in the fi eld of high-modulus,
hi g h-strength rigid a nd flex ibl e po lyme r fibres whic h
a re visua li zed as agg rega tes of hi ghl y extended
mo lec ul a r chains in a hi ghl y c rysta llin e state and
co nsequentl y have axia l moduli close to th e
th eo re ti ca l crys ta l m oduli in the ch a in d irection a nd
ve ry high a xi a l stre ngth s.
2.1 Axial Modulus
The es timated va lues o f crysta l modu lus in the
chai n direction for three well-known polymers are:
324 GPa (3745 gjde n) fo r hi gh de nsity po lye th yle ne
(Ht>PE), a n a lipha tic fl ex ible c ha in po lymer l ; 11 0
GPa (903 gjden) for P ET , a n a liph atic-aroma tic
semi-ri gid chain polymer 2 ; and 194 GPa ( 15 1~ gjden)
fo r Kevlar, a who ll y aroma ti c ri gid cha in pol ymer 3
Th e highe st meas ured mod uli a re fo r so luti on o r
ge l- spun fibres ma de from these po lymers: 288 G Pa
for HOPEI , 30 GPa fo r PEpa a nd 125 GPa for
Kevla r49 (ref. I ). The moduli of the melt-spun HOPE
a nd PET commercial fibres a re a bo ut 5 GPa a nd 10
GPa respecti vely5. Apparently, the geo metri c
structure of indi vidua l mo lecules has an important
ro le to pl ay: in melt-spun fibres, a significa nt a mo unt
of chain foldin g may occur which co uld a ll ow only
limited chai n co ntinuity in the axia l directi o n.
C onsequently , the numbe r of tie mo lecu les,
pa rticularl y the ta ut tic mo lec ules, which co nnec t the
GUPTA & RADHAKRlSHNAN : NATURE OF CRYSTALS AND THEIR INFLUENCE ON FIBRE PROPERTIES
crystalline blocks in the fibre in the axial direction ,
will be relati vely small. Moreover, there will be
entanglements a nd the overall orientation of the
molecules will not be close to the ideal. As a result, the
axial modulus is low . In gel:spun HOPE and PET
fibre s, the molecules are more likely to be highl y
extended like the molecules in Kevlar fibre and,
therefore, offer much higher resistance to
deforma tion .
It is noteworthy that in the above illustration , the
polymer with the most flexible chain has the highest
estimated and measured modulus. The ph ysical basis
of this was explained by Fra nk 6 as follows . "The
Young's modulus for diamond in the [110] direction is
1160 GPa . In the [110] direction di a mo nd is
composed of full y aligned zig-zag chains of carbon
just like th ose in po lyethylene, utili zin g half the
neighbour-to-neighbour bonds in the crystal, while
the other half of the bonds are at ri ght angles to thi s
direction , con tributing nothing to the Young's
modulu s In this direction , just as the
carbon-to-hydrogen bonds in fully aligned
polyethylene contribute nothing to its longitudina l
Young's modulus. The cross-sectional a rea per chain
in di a mond is 0.0448 nm 2 , four times small er th an
polyethylene, 0. 182 nm 2 • Hence, from thi s analogy ,
we could expect a modulu s of 285 GPa fo r fully
aligned pol ye th ylene, well a bove th a t of steel".
There are two noteworthy fea tures of HOPE which
may contribute to its hi gh crystal modulu s in th e
chain direction . First, the cross-sectional a rea per
chain of HOPE is the smalles t (HOPE, 0.182 nm 2 ;
PET, 0.217 nm 2 ; Kevlar, 0.205 nm 2 ), which ensures
that a relatively la rger numbe r of chains per unit
cross-sectional a rea take the load . Second, the HOPE
molecules in the crystal are in the ex tended tra ns
conformat ion, as shown in Fig. I, whereas the PET
molecules ma y no t be exactl y stra ight. The Kcvl a r
molecules, though in the ex tended tra ns
conformation, a re believed to aggregate in the form of
a radial system ofaxially pleated la mellae as shown in
Fig. 2 (ref. 7). Th e angle betwee n the adj acent
components of the plea ts being a bo ut 170°. Thi s
results in a variation from linea rity of 5" (ref. 8). Th e
recently introduced Kevlar 149 fibre is repOrled 9 to
have a modulus of 170 GPa (1343 g/d en) whi ch has
been achieved by producing a more perfect hi ghl y
oriented fibre without pleated structure I o. Tha t chain
conformation is of primary significance ha s been
adequately demonstrated by a study of the crystal
moduli of around 25 pol ymer-ba sed fibres and
film s!l . Fo r crystals with helica l molecules lik e
iso tactic pol ypropylene with 3/1 helix , the
cross-sectional a rea is about twice that for HOPE, the
I
(0)
I
(b)
(c)
Fig. I- Skeletal structure of a mo lecule in the crystal of (a)
polyethylene, (b) po ly(e th ylene terephthalate). a nd (c) Kevla r
Fig. 2- A schematic sketch depicting the radiall y arranged
pleated morph ology of ' Kevlar 49' fibre (ta ken from ref. 7)
force required fo r I % ex te nsion is a round one fifth
a nd the axial modulus 30 GPa. Crystal s with more
loosely packed helices ha ve sli lll owe r ax ial moduli : 7
GPa for 7/2 heli x a nd 4 G Pa fo r 4/ I helix. I n a heli x,
various combinations of tra ns a nd gauche
conformations are possible due to which bond
rota tion beco mes the predv minant deformation
mec hani sm on C1 :\i :l! loadi ng. It needs to be
101
INDIAN J. FIBRE TEXT. RES., MARCH 1991
emphasized that even for extended trans structures,
the presence of bulky side groups can result in a large
cross-sectional area. In this context, an interesting
observation has been reported for polydiacetylene
single crystal fibres for which a plot of the Young's
modulus in the chain direction VS. the reciprocal of
area supported by each chain of the crystal gave a
straight line passing through the originl 2. This
suggests that the product of Young's modulus and
chain cross-sectional area is a constant if the
hack bones of the polymers are assumed to have the
same stiffness. The large side groups on the
polydiacetylene crystal molecules result in a
relatively low modulus of 50 GPa for this polymer
crystal.
The Youn g's moduli of the crystalline regions of a
number of polymer-based fibres and films in the chain
direction, mostly taken fro m reference II, are plotted
as a function of the reciprocal of the cross-sectional
area of the corresponding molecules in Fig. 3. The
data points appear to form three broad gro ups. Th e
polymers with a carbon backbone in the planar
zig-zag o r nearly planar zig-zag conforma tion are
seen to fall on line C. Polymers with a carbon
backbone in a helical conformation tend to fall on line
B. Polymers, which, in addition to carbon, contain
oxygen in their backbone and have a helical or planar
conformation, have lowest moduli for a particular
cross-sectional area and fall on line A. PET (marked
as number 7 in Fig. 3) wh ich also contai ns oxygen in
the backbone, however, does not fall on line C; it has a
relatively higher modulus apparently due to the
stiffeni ng of the structure due to the presence of the
aroma tic ring. The a. form of poly (ethylene
oxybenzoate) (a.-PEO B which is marked as number I
in Fig. 3) has relatively low modulus though it has an
aroma tic rin g in its structure. This is apparently
because the chains are believed to be greatly
contracted in the a. form which results with a
conformation different fro m that of an extended
type.
The development of rigid-rod , aromatic
heterocyclic ordered fibres with extended chain
molecules 10 has made it possi ble to extend the
stiffness range still further . An example of such a
system is the high performance poly (p-phenylene
benzobisoxazole) (PBO) fibre with a density of 1.58
g/cm 3 , a theoretical tensile modulus of 730 GPa,
measured crystal modulus (by X-ray) of 477 GPa and
measured fibre modulus of 318 GPa (ref. 10).
2.2 Axial Strength
Like modulus, the axial strength of the polymer
in the chain direction would also be expected to
depend predominantly on the area supported by
102
l~
"~
>'01
1140
11
~
'"
,80
16
w
,00
:;
~
~
~
160
'"
liO
80
40
n
'jCH AIN CROSS-SECTIONAL
i3
AREA (nm-')
Fig. 3- Plot of measured axial Young's moduli of some polymer
crystals vs the reciprocal of their molecular cross-secti onal area
[I - Poly(ethylene oxybenzoate) (IX-form); 2- Polypivalolactone
IX-form
(heli x
2/1);
3- Pol y(iso butylene
oxide)
(distorted planar zig-zag 2/ 1); 4-Poly(ethylene oxide)
5-Polyoxymeth ylene
(helix
9/5);
(helix
7/2);
6-Polytetrahydrofuran(planar zig-zag); 7- Poly(ethylene
terephthalate) (nearly planar); 8-lsotacti~ polyvinyl tertiary
butyl ether (helix 4/ I); 9- Isotactic poly( 4-methyl pentene-I )
(helix 7/2); IO--Isotactic polystyrene (helix 3/1); II - Isotactic
polybutene-I (helix 3/1); 12- Isotactic polypropylene (helix 3/1);
13- Poly[ I, 6-di(N-carbozolyl)-2, 4-hexadieneJ (planar zig-zag);
14-Polytetrafi uoroethylene (helix 15(7); 15- ~-Polyvinylidene·
fluoride (slightly deflected planar zig-zag); I6-Polyvinyl alcohol
17- Polyethylene
(planar zig-zag);
(planar zig-zag);
18- Diamond (extended)
each chain. Since the limiting value of the load is
defined by chain rupture, Ohta 13 esti mated the
ultimate axial strength of various polymer crystals
a long the chain length assuming that a t the breaking
point, the carbon-carbon bonds a re ruptured . The
estimated ultimate strength values of the crystal
along with the highest reported measured values
and the values for commercial fibres for the three
cases considered earlier are given in Table I . There
are two noteworthy observations. First, the values
of crystal strength in GPa are quite close for the
three polymer crystals, unlike the moduli . This is
obviously due to the very approximate nature of the
model used in calculating the theoretical strength.
Second, the reported axial strength values are much
less than the crystal strength unlike in the case of
modulus where it has been possible to achieve
values which are much closer to the theoretical
maximum. The highest reported 14 strength (80
GUPTA & RADHAKRISHNAN : NATURE OF CRYSTALS AND THE IR INFLUENCE ON FIBRE PROPERTIES
Table I- Ultimate axial strength along the chain direction for
some polymers
Polymer
HOPE
Kevlar
PET
Estimated axia l
Highest
strength
reported axial
GPa (gjden)
strength
GPa (gjden )
31.60
(372)
29.74
(235)
28 . 13
(232)
6.80
(80)
3.50
(27.66)
1.9
(IS .O)(ref. 4b)
A ve rage axial
st rength of
commercial
fibre
GPa (gjden)
0.76
(9)
3.16
(25)
1.15
(9 .5)
gjden) is for a ge l-spun ultra-high mo lecular weight
HOPE; it is less than one fourth the esti mated
theoretica l max imum value for this polymer. This is
primarily because in considering the axia l st rengt h
of fibres , account must also be taken of the flaws
that are present in the fi bre such as microvoids,
particulates, microscopic cracks , chain ends and
othe r sources of stress concentration. Th ese flaws
have little effect on axial modulus, which involves
very low strains, but ha ve significant effect on axia l
strength which is measured at the lim iti ng strain of
the material. The effects offlaws a re best considered
in terms of 'the Griffith theory of fract ure 15
accordi ng to which the breaking stre ngth is
proportional to the square root of the elastic
mOdulus, E. To assess the potential of the various
processing routes for producin g high strength
HOPE, Ohta plotted the measured axia l breaking
strengths o f a number of fibres against the square
root of their axia l moduli . From these linear plots,
he concl uded 13 that the tenacity of the fibre is a
function o f the processing method; the limiting
tenacity values were obtained by ext rapo latin g the
experimental data to the theoretical modulus of
2775 g/den. The fibrillar crysta l growing or ge l fib re
drawing of ultra-high molecular weight HOPE gave
the highest limiting tenacity va lues of 59-93 gjden,
hot drawing or zone drawi ng gave a limiting value
of36 g/den while the solid state extrusion resulted in
a limiting va lue of 7 gjden only.
In a set of elegant experiments, Matsuo and
Ogita 1 6 too k HOPE of molecular weights 1,3 a nd 6
million respectively and studied the concentra tiondependence of draw ratio for the three dried gel
films . A draw ratio of350-400 could be obtained at a
concentration of 1.6 g/ IOO ml in the case of the
lowest mo lecular weight (I million), 0.65 g/ 100 ml in
the case of intermediate molecula r weight (3
million) and 0.4 gj lOO ml in the case of the highest
molecul a r weight (6 million) HOPE. Higher and
lower concentrations than these optimum levels
were found to lead to lower draw ratios and,
therefore, inferior mechanical properties . The
a uthors concluded that most of the chain molecules
in the regime of low concentration are random coils
having coupling entanglements that wi ll be
predominantly intra-molecu la r in nature. On the
other hand , solutions co rresponding to the regime
of high concentration a re thought to consist of
interpenetrating random coils which form a large
number of coupling entanglements that a re both
intra- and inter-molecular. For specimens prepared
from sol ution s with critical concentration, which is
different for different mo lecul ar weights , it may be
expected that there exists a su'itable leve l of
entangled meshes that act as in ter- lamella r
crosslinks a nd effecti ve ly transmit the drawing
force, a nd therefore the possibility of polymer
cha in s slippin g past each othe r without interconn :
ctio n is very low . T he sh ift of the critical
concentrati on to lower vC;!.lues with increasing
molecular weight is probably due to the fact that the
number of the co uplin g entanglement meshes of
HOPE with very high molecula r weight (say 6
million)
increases drastica ll y with
increasing
concentration.
2.3 Thermal Expansion
Polymer crystals like those of HOPE, PET,
polyvinyl alco ho l, Kevlar and nylon 6 with extended
chain structures exhibi t negative thermal expansion
a long the chain direction and positive expansion in
the transverse direction 1 7 . The latter reflects the
weakness of the interchain interactions while the
fo rmer is postulated to arise from the strong elastic
anisotropy in a polymer crysta l due to which
torsiona l and bending motions in the chains are more
highly exci ted than the stretching modes and thi s can
lead to a n effective reduction in the interatomic
distance a long the chai n axis. An a lterna te model 1
postulates that internal stresses are responsible for
this negative thermal expansion .
2.4 Melting Point
At the melt ing point, eq uilibrium ex ists between
the liquid a nd crystal ph ases . Th e equilibrium
meltin g point ofa crystal, Tm o' is given by the !1H m
(enthalpy of melting)/!1Sm (entropy of melting). Th is
definition would only a ppl y to crysta ls of infinite size
(no surface effects) which co ntain only equilibrium
defects, if any. The crystals in the fibres are metasta ble
and contain non-equilibrium defects. The melting
pointof a fibre , T m , may, however, also be considered
in terms of the thermodynamic relationship given
above. The simplest approaches identify large heat of
fusion with strong in termolecular forces l 8 . ]t has
been pointed out by Mande lkern l 9 that attempts to
103
INDIAN J. FIBR E TEXT. RES., MARCH 199 1
correlate the melting 'points of polymers with
intermolecular interacti ons utili zi ng the co hesive
energy density o f the repea ting uni ts as a meas ure of
these interactions has been notabl y un successful
si nce no simpl e correlation is o bserved between ~11
and /j,H m , as is clear from the ex tensive da ta ava ilable
in the litera ture. It was, therefo re, thought th a t it is
more lik ely th a t tlSJ11 may be o f prime importa nce in
esta blishin g the va lu e of TJ11 .
Do le and Wunderlic h 20 emphasised th a t in the
thermodynamic eq ua ti o n th e heats and entropies of
fusio n represent the d iffe rences in e ntha lpy a nd
entropy between the liquid a nd crysta lline sta tes; it is,
therefo re, necessa ry that both these sta tes of ma tter
shou ld be cons id ered in any inte rpreta ti o n of the
meltin g point. They furth er s ugges ted that th e hi gh
melting point of poly ami des is due to the low ent ro py
of the liquid phase (perhaps due to the prese nce of
hydroge n bo nd s) while the low meltin g point of
a liph a tic po lyesters res ults chi efl y from a low hea t of
fusion.
So me interesti ng thermal st udi es 21 have been
recently made on two liquid crysta l po lyes ters
(po lymers A and B) with mo lec ular we ights in the
range 10,000-30,000. Po lyme r B had a stiffer c hain
than polymer A a nd both were, in turn, stiffer tha n the
PET chain. The X-ray c rysta llinities of the polymers
were determined and were co mbined with the a rea of
the DSC melting peak to obtai n the va lues for the heat
of fusion ofa unit mass of three-di mensio na l crysta ls
(/j, H F ). Res ults based on data fo r the sa mples
prepa red by slow coo lin g, ta ke n fro m re ference 2 1,
a re li sted in T ab le 2 a lo ng with the value fo r the
convent io nal P ET. The ta bl e a lso lists va lues for th e
entro py of fusio n /j,SF = /j,HF I Tm . Th e data ha ve
bee n discu ssed with the he lp of a schem a ti c
representation of the mo rph o logies below and above
the crystal me ltin g points (Fig. 4). As show n in T a ble
2, /j, H F for the liqui d crysta l polymers is signi fica ntl y
less than that for P ET. This has been attrib uted to the
imperfections with in th e crysta l la tti ce causing poor
cohesio n of cha ins. Th e rea son for the imperfec ti o n is
stated to be th e non-regular nat ure of the chain in
which th e probabilit y of lo ng runs of reg ul ar
sequences is low. Sin ce th e melting po ints of th e two
rigid cha in pol ymers arc no t ve ry different to th at or
PET. th e low e nth a lpy of fu sio n a lso renects a mllch
lowe r en tro py of fusio n than in PET (Tahle 2). The
reduced /j,S" is a direct conseq ue nce or the extra
stiffness of the chai ns a nd as show n in F ig. 4, th e
melting process is very dilTerent in the liquid
crys ta lline and th e co nven tio na l polymer sys tc ms. It
is conclu ded from this stud y that cha in stiff ness is th e
mam property that will determine whe th er th e
104
Below Tm
a
Above Tm
i \
IiiI \ WI
~
Fig. 4-A schematic diagram depictin g the mOll' ho logies above
and below the crystal melting poi nt for (a) ri gid chain nematic
polymer a nd (b) conventional polymer wi th chain fo lded lamellar
crystals. The thicker poi nt s of the lines represeni regions where the
chai ns fo rm three-di mens ional crystal lattice (ta ken fro m ref. 21)
Table 2- T he hea l a nd entro py of fusio n for some polymers
(ref 2 1)
Polymer
Till
(K)
Po lymer A
Polymer B
PET
5 13
563
530
X- ray
I1H,.
I1S"
crys ta llinit y (k Jkg - ' ) (kJ kg - , K - ' )
(%)
17
21
T ypica ll y 50
40
20
135
0.08
0.04
0.25
po lymer melt will exist as a mesoph ase or as a
conventional iso trop ic melt. Now if the enthalpy of
fusion is no t made low, the po lyme r will have a very
high TJ11. Thu s, to'make processa ble liquid crystal
po lymers, o ne must introduce irregularities in the
cha in to limit the effec tive bonding of the crystals.
3 Crystal Defects
3.1 "Ia ture of Defects
It is ge nera ll y believed th a t crystals in fibres are
hard a nd und eforma ble 22 • Bueche 23 has, howeve r,
emphasized th a t "the crystallites ca nn o t be as perfect
as to impart grea t ri gidit y to the polymer, however, o r
the m a terial will be too brittle to be useful. It is
important neve rt he less that th ec rysta lli tt:s be stabl e
to rather hi gh tem pera tures so th a t they wi ll not melt
GU PTA & RADHA KRI SHN AN : NATU RE OF C RYSTALS A ND THEIR INFLUENCE O N FIBRE PRO PERTI ES
o ut during norma l ha ndling o f the fibre" . C rysta lline
polymers a re much less perfect fro m the po int of view
o f crysta l regularity th a n simple substa nces 2 4, the
possible impe rfecti o ns ra nging fro m poi nt de fects in
the la ttice to a trul y a mo rph o us phase. Amo ngst the
de fects within the la ttice, the cha in end s represent a
di scontinuity tha t ca n lead to d islocatio ns whose
density has been repo rted to be of the o rder of 10 3 1m2
(ref. 12) . The fo ld surface, pa rticul a rl y the gross
crysta llinity defici ency a t the fold surface, represents
a significa nt so urce of imperfecti o n . It is no tewo rthy
th at ri gid cha in po lymers usua ll y crysta lli ze as
ex te nded m o lecules whil e fl ex ible cha in pol ym ers
often fo rm fo lded c ha in crysta ls 25 . Fro m the po int o f
view of defects, po lyeth ylene has bee n studied in
detai l. So me la ttice defects in po lyethylene a re
illustrated in Fig. 5 (ref. 26). Kinks, jogs a nd Reneka r
de fects a re co mforma ti o na l defects. As sho wn in the
fig ure, in the case of kinks a nd jogs, a pa rt o f the cha in
is di splaced perpendi c ula r to the lo ng axis.
A combined study using wide-a ngle X-ray
di ffraction and sma ll-a ngle X-ray di ffraction of a Fig. 5--A schematic rep resen ta tion of some la ttice defects in
pol ye thylene crystals. From left to ri ght: all-trans co nfo rmation
series o f po lyethylenes, mostl y lo w density, with a (defect-free), Renekar defect, kink a nd jog (taken from ref. 26)
wide ra nge o f cha in de fect conce ntratio ns (0 . 1-7 %)
( b)
crysta lli zed fro m the melt has been re po rted 2 7 . The
co nc urrent unit ce ll expa nsio n a nd lon g peri od
decrease wi th increasing cha in de fect co nce ntra ti o n
(a)
lead to a picture o f cha in defects (bra nches,
un sa tura ti o ns) being distributed between the
crysta lline la mell ae a nd the surface layer. Based o n a
7SA.
model, which assumes inclusio n o f defects within the
la ttice by mea ns of genera ti o n of kinks, a n estima tio n
o f the concentra tion ot cha in de fects inco rpo ra ted
into the crystal la ttice « I %) is a ttempted . The
density of defects in no n-crysta lline regio ns turns o ut
to be muc h la rge r tha n their concen tra ti o n in the
crysta lline regio ns a nd s uppo rts the view of a
cl uste ring of defects. F ig. 6 schema ticall y illustra tes
'
the prevail ing exclusion of defects fro m the crysta ls of WllWllUl.U.WllWllWllWUJ.
two samples of co nventi o na ll y bra nched a nd linea r Fig. 6-A schematic rep resen tation of the dist ribution of chai n
po lyethylene with the tota l amount o f defects per unit defects between crystalline lamellae and amo rpho us laye rs fo r (a)
volume being 0.17 a nd 2. 53 % respectively. It is seen low-densi ty polyethylene and (b) high-densi ty polyethylene
(ta ken fro m ref. 27')
that bra nchin g has a most dra m a tic effect o n the
densi ty of defects in th e a m o rph o us phase; the
average cha in sepa ra ti o n in the amo rphou s regio ns co mme rcia l P ET fibre, a cha in end is lik ely to be
registers a n increase as increasin g numbe r o f de fects present in each vo lume eleme nt of 22A as a side- a
very hi gh density o f ch a in ends.
enter thi s layer .
The effect o f crysta l defects on fibre properties will
F o nta ine et al.2 9 have ma de a deta iled stud y o f
no w be briefl y considered .
solid sta te therma l trea tment o f isotropic PET sheets;
they measured the degree of crysta llinity, the di sorder
3.2 Effect on Thermal Properties
pa rameter a nd the average meltin g tempera ture. The
During ra pid c rysta lli za ti o n o f po lymers, de fects results of their stud y a re s umma rized in T a bl e 3. It is
suc h as kin ks a nd chai n ends are inco rpo ra ted in the o bse rved fro m thi s ta ble tha t the crysta ls become
fi brilla r c rysta ls. It has been estima ted 2 8 that in mo re perfect with increasing a nnea lin g tempera ture
i
j
T
105
INDIAN J. FIBRE TEXT. RES., MARCH 1991
Table 3- Morphological parameters for PET subjected to solid state the rmal treatment (ref. 29)
Sample
No.
I
2
3
4
5
Thermal history
C rystalli zed at 200T
As for sample I and
further heat treated
up to 215' C
As for sample 2 and
further hea t treated
up to 230' C
As for sample 3 and
further heat treated
up to 245'C
As for sample 4 and
further heat treated
up to 250'C
Mean long
spacing
(.&.)
X-ray
crystallinity
126
125
Average
melting temp.
(%)
Disorder
parameter
k
53
60
6.7
6.3
240
243
125
59
5.7
248
128
58
5.7
252
129
63
5.4
256
a nd the melting point also increases. However, the
increase in the degree of crystallinity occurs by
heating from 200 to 21YC but not beyond 21YC.
Hence, the increase in melting temperature for
samples treated at the higher temperatures is ascribed
to a n overall increase in crystal perfection.
On the basis of the above results a nd low angle
X-ray intensity, Fontaine e t at.29 proposed a
mechanism for the reorganization at the
crys tal-a morphou s bo undary laye r as·a result of the
an nealing treatment (Fig. 7). The crystal
imperfections inside the crystal (not shown in the
fig ure) are considerably reduced on annealing as they
ap parently migrate out of the crystal.
The crystals with high defect density have been
shown to have very low melting points while drawn
fibres annealed for long durations have relatively
higher melting points 3 0 . 31 ; the difference being of the
o rder of IOoC (247'C to 257'C). When the fibre is
constrained in the DSC cell so that it cannot shrink,
the melting point of lhe heat-se t fibre can go up to
264°C (ref. 32). The crystalline density of PET
originally calculated at 1.455 glcc (ref. 33) is now
believed to be closer to 1.5 15 glcc (ref. 34); it appears
to be morphology-dependent.
3.3 Effect on Optical Properties
The intrinsic crystalline birefringence of PET
which represents the limiting birefringence of its
perfectly oriented crystalline unit, assuming it to be
transversely isotropic, has been reported to be
anywhere between 0.212 and 0.31 (ref. 35). It was
observed that the lower values arose from studies on
cold-drawn fibres while the higher values were based
on studies on heat-set fibres. It was suggested 36 that if
IOn
-~
CC)
o
(T e =200C)
----ri~
----tl~
---~
----+-I':Y
I
I
-----------+i~
----'-i~A
_ _ _~i-::-'" \'
==-----+~? \)
-----------fl~
Fig. 7- A schematic sketch depicting the ordering and smoothing
effect at the crystalline amorphous interface of the lamellae
(taken from ref. 29)
a perfect crystal had an intrinsic birefringence of ~co,
then a crystal with perfection index Pc will have a
birefringence Pc !:lnco' assuming that an ideally perfect
crystal will have Pc = I . The measured birefringence
!:In of a fibre of crystallini ty P and crystallite
orientation fc can then be written as
!:In = Pc !:lncop!c + !:lnamo (I - P}famPam
where Pam is the perfection index of the amo rphous
phase, .1nam " its intrinsic birefringence and Jam its
orientation. Since cold-drawn fibres have crystals
with high defect density compared to the heat-set
fibres, the dependence of measured intrinsic
birefringence on sample morphology reported in the
literature is not surprising.
GUPTA & RADHAKRISHNAN : NATURE OF CRYSTALS AND THEIR INFLUENCE ON FIBRE PROPERTIES
The morphology-dependence of the measured
values of intrinsic birefringence of the crystalline and
amorphous phases of PET can be very approximately
represented by the schematic diagram shown in Fig. 8
which is based on the data obtained on fibres with
different structures and morphologies and on the
assumption that the defects which get incorporated in
the crystal can migrate out of it during solid state
thermal treatment. Such an assumption is implicit in
the model of Balta Calleja et al. 2 7 The wide range of
intrinsic birefringence values reported in the
literature can thus be traced to the differences in
morphologies of the PET samples investigated. This
model receives added support from the production of
PET fibre with /).n = 0.26 by suitable combination of
drawing and heat-setting 3 7 .
Sheldon 40 had shown that PET film extruded at
low rates showed higher crystallization rate than film
extruded at a higher rate. He attributed this to the
disruption of crystalline remnants in the melt at high
speed of extrusion which can then no longer act as
o-boiling waler (100°C)
60
/
I
I
40
.
I
I
I
20 I
I
I
Q
In some recent work 3 8 on drawn PET films and
fibres drawn to different extents under identical
conditions, the films showed relatively lower
shrinkage at 220°C compared to the fibres at draw
ratios of 4 and above, as shown in Fig. 9. X-ray
diffraction studies on a drawn film and a drawn fibre ,
both of draw ratio aro und 4.16 and prepared under
identical conditions, had shown that the crystal
disorder parameter, k;,is 10.31 for the film and 5.5 for
the fibre. As shown in Fig. 10, the crystal disorder
reduced on heat-setting of both sets of samples but the
free-annealed samples showed marginally higher
perfection than the taut-annealed samples. Since
defects within the lattice reduce the energy for
molecular motion 39 , it was proposed that the
presence of defects can result in faster crystallization
which would stabilize the structure, thus reducing
shrinkage.
(200"( )
• -silicone oil
w
3.4 Effect on Crystallization
FILM
(0 )
80
<t
~
0
Z
i'i
I
If)
FIBRE
80 -
( b)
60
,
I
i
40
20
0
2
1
3
5
I,
6
DRAW RATIO
Fig. 9-Shrinkage as a function of draw ratio for PET film and
fibre (taken from ref. 38)
12 .0 , - - - - - - - - - - - - - - - - -
-PERFECTION INDEX FOR AIoIORPHQUS lHY
HIGH - - - - MEOIUM - - - - - LOW
10 ·0
...u
0 · 32
o TA
Cold Drawn-Hot Drnwn -Heat-set
o FA
8 .0
z
'"~ 0.26
a:
....
.......
~
I.L
~
'"a:
iii
6·0
~QI:!trol Fi b rez
«
a:
~
ffi
a
Film
8
g
100
140
~
4·0
~Fibrez
a:
a
III
o
0
LOW
MEDIUM
PERFECTION INDEX OF
2·0
HIGH
CRVSTAL---
Fig. ~A.schematic representation of the dependence of intrinsic
crystalline birefringence (I1n"o) and intrinsic amorphous
birefringence (6.n amo) on the perfection of the two phases
0
0
HEAT - SETTING
180
220
TEMPERATURE
260
(C)
Fig. IO--Dependence of disorder parameters of PET film and
fibres on heat-setting temperature
107
INDIAN J. FIBRE TEXT. RES., MARCH 1991
nuclei. Since the film used in the present investigation
was cast at 10m/min while the fibre was produced at
spinning speeds of 1000 m/ min, the presence of
incipient nuclei in the film was possible. PET fibres
spun at low speed (10m/ min) were, therefore ,
produced and drawn under identical conditions to a
draw ratio of 4.16. They were found to crystallize (as
measured by increase in density when subjected to
heat-setting at 8SOC) to a greater extent and
approached the crystallization characteristics ofthe
films, as shown in Fig. II. The shrinkage
characteristics of this ma terial were also closer to
those of the film . The higher shrinkage of the fibre
spun at 1000 m/min and then drawn, in relation to that
of the film cast at 10m/min and then drawn, may thus
be related to the lower crystallization potential of the
former.
4 Coupling Effects
4.1 Nature of Coupling
In a fibre, a single molecule can form part of several
ordered regions (crystals) as can be concluded from
the following ohservation . The usual length of a
molecular chain is, in general, far greater than
the size of the crystallites. For example, considering
only the extended molecule J the length of the
zig-zag polyethylene mol ec ule of molecular
weight 50,000 is about 4500 ~ (ref. 41). Acrystalline
region, on the other hand , may be only 100-500 Along
and hence one molecula r chain is considered to pass
through many crystalline and non-crystalline regions
successively. The ordered and disordered regions are
thus coupled. The coupling may be in series or in
parallel and the nature of coupling can have
important consequences on fibre properties.
The nature of coupling in a textile fibre can be
illustrated with the help of Prevorsek's model 42 for
1.384
Film
1.380
C"')
-
E 1.376
u
( ]I
>
1·372
t-
iii
~
1.368
o
1.364
1.360
0
0 .4
0.8
1.2
1.6
2.0
2.4
'2.6
) .2
. T I M~ (n )
,
.'
Fig. ll'-Density data for PE! ,nl~and fibres heat treated at 85"C
for different time .periods
10 ~
PET fibre (Fig. 12), which is extensively quoted in the
literature. Two noteworthy features of this model are:
(i) the intercrystalline links in fibrils, which assist in
load transfer, are less in number than the number of
molecules in the crystal; and (ii) the interfibrillar
extended molecules provide relatively more
continuity. The former, i.e. intercrystalline tie
molecules, provide series coupling in the fibril
between the crystallites and the latter, i.e. the
interfibrillar molecules, provide the parallel coupling
between the fibrils .
There has been quite a detailed investigation of the
structural and morphological changes that take place
in a conventional highly oriented, crystalline
high-density polyethylene fibre or film as a result of
annealing. The model of a conventional drawn
HOPE filament 43 is shown schematically in F ig.
13(a); on heat-setting, considerable reorganization
occurs and the poor phase separation of the
cold-drawn filament may be replaced by a more
distinct phase separation [Fig. J3(b)]. Fischer and
Fakiroy44 also studied the reorganization that occurs
in oriented crystalline HOPE on annealing under
different tensions and temperatures. At low
temperature and high tension, the phase separation
was represented by the model shown in Fig. 14(a)
while at higher temperatures and lower tensions,
there was a very distinct phase separation, as
observed by low-angle X-ray diffraction which is
schematically shown by the model in Fig. 14(b).
Ward 4s has recently considered the mechanical
anisotropy at low strains in polymers and has pointed
out that the various models fo r anisotropic polymers
fall in two distinct categories, depending on whether
molecular orientation or the composite nature of a
crystalline polymer is the starting point. The two
model hierarchies are then applicable to different
polymer systems. On the basis of molecular
orientation, the single-phase aggregate model 4 6 may
be successfully applied to amorphous polymers,
low-crystallinity PET, drawn low-density PE, Kevla!f
a nd carbon fibre; the single-phase sonic modulus
modd 47 is applicable to all oriented polymers and the
two-phase sonic modulus model 4 8 is applica ble to
polypropylene fibres and to a limited ex tent to PET
fibres. If the composite model approach is used , the
series-parallel unit cube model 4 9 can be applied to
annealed drawn linear PE a nd PP and the lamellar
9rientation modep o to oriented sheets of PE,
annealed PP and PET. If tie molecules are added to
composite models s 1, they can be applied to all drawn
polymers . For high modulus polyethylene, the
addition of crys'talline bridges to composite model s I
allows their mechanical anisotropy to be understood .
GUPTA & RADHAKRlSHNAN : NATURE OF CRYSTALS AND THEIR INFLUENCE ON FIBRE PROPERTIES
ttti+H#Hl--.. CRYS TA LLiT ES
EXTENDED
NON-CRYSTALLINE
MOLECULES
DISORDERED
DOMAINS ~/
Fig. 12- Prevorsek's model for PET fibre (taken from ref. 42)
(a)
( a)
( b)
( b)
\J ~
I UU U
Fig. 13-Model for conventional drawn HDPE: (a) cold drawn
with poor phase separation and (b) annealed ~ith distinct phase
separation
The short fibre reinforced polymer composite
modeJ52 has also been applied to high modulus
polyethylene.
The above considerations suggest that there can be
no unique model for a fibre since the structure and
morphology of the fibre are so sensitive to thermomechanical treatments. The model 42 ofPrevorsek for
PET is thus representative of a very specific
Fig. I4-A model depicting (a) limited phase separation for
low-temperature, high-tension annealed PE and (b) more distinct
phase separation for high-temperature low-tension annealed PE
(taken from ref. 44)
morphology out of a wide spectrum of possible
morphologies of PET fibre .
4.2 Effect on Mechanical Properties of Fibres
Extensive studies s3 .s4 on heat-set PET have shown
that high temperature (200°C and above) annealing
of PET fibres and films results in predominantly series
type of coupling with clear-cut phase separation
109
INDIAN J. FIBRE TEXT. RES. , MARCH 1991
occurring. The amorphous orientation factor drops
fro m 0.6 to 0.2 a nd the elastic modulus a lso decreases
by a factor of 4 or so . Annea lin g at low temperature,
o n the other hand, reta in s significant a mount of
parallel coupling and the samples have a relatively
hi gher amorphous orientation and elastic modulu s.
The samples heat-set with their ends ciamped have
greater degree of parallel coup ling. It is interesting to
recall that in industrial processes, the fibres a nd
fabrics a re subjected to heat-setting treatment under
considerab le constraint.
'eo
.,.c:
"
1/1
It has recently been reported 56 that the dynamic
mechanical relaxations with peaks at around 50°C
and 90°C in HOPE seem to a rise due to defect
diffusion in the crysta lli tes with some influence of the
amorpho us matter in the interfacial region . This type
of interaction between the two phases would be
expected if they are in tandem or are coupled in series.
Series coupling is significant in fibrillar fibrous
struct ures, as is evident from the following two
examples. In six HOPE fibre and fi lm samples with
crystallinity rangin g from 50 to 85 % the macroscopic
moduli ranged from 0.7 to 15 GPa (ref. II), though the
crystal modulus was constant (around 250 GPa). The
resistance to initial deformation (stiffness or
modulus) is thus apparent ly dominated by the more
complia nt amo rph ous phase and , therefore, homogeneity of stress or series coupling may be assumed to
be predominant. In another study 5 7 though the
crystal modulu s of PET remained constan t a t 110
GPa from 25 to 21YC the axial modulus of the
filament decreased from 9 GPa at 2YC to I GPa
a t 200°C. This decrease is apparently due to a
decrease in a morphous modul us; a series co upling is
again indicated.
110
0
Series (AI
0
Poroll.l
(~l
6
II(
...
1&1
1&1
1.
Z
"'~
II(
.:I
z
:l
0-6
Q.
::;)
0
U
..J
T o illustrate the a bove observations, the experimental data on series and parallel co upling
parameters for free-ann ea led and taut-annealed P ET
fibres and their Instron and sonic moduli a re
presented in Figs 15 and 16 respectively. The
correlation between modulus a nd couplin g parameter is quite apparent. In taut -annealed samples, the
parallel coupling is hi gher and so is the modulus. In
free annea led samples, the series co upling is relatively
higher a nd they have low modulus.
The recovery behaviour of these two sets of PET
fibres from tensile strains up to 0.15 was also
in vestigated 55 . The taut-annealed samples showed
superio r recovery behaviour which was dominated
by both the crysta lline a nd amorphous phases; the
recovery behaviour of the free-annealed samples, on
the other hand , was controlled principally by the
amorpho us phase.
Fno- onn4lol4ld
1&1
......J
"'
Toul-onn.ol.d·
II(
~
0
Z
4
1/1
1&1
II(
0·5
1&1
1/1
0-2
°O~'~I~O~0--~~--~18~O~--2~2~0--~2~50~
HEAT-SETTING TEMPEAATUAE(OC)
Fig. l5--Coupling parameters as functions of heat-setting
tempe ratu re for PET fibres
1B O r - - - - - - - - ---, . - - - - -- - -----...
(a)
Sonic
(b)
Inst ron
~Qntrol
·....el--..._.I--_e
TA
"0
g; 12 0
U1
3
90,-----
::J
o
o
~
Fig. 16-Dependence of (a) Instron modulus and (b) sonic
modulus of PET fibres on heat-setting temperature
Wool has a ve ry compliant matrix reinforced by
ex-helices, wh ich form the crystalline phase 5 8 . It ha s
been quite convincingly shown that the helices
dominate the axial defonnation of WOOP 8 and thi s is
attributed to the parallel coup lin g betwee n the
globular protein matrix and the helix; the stress is in
fact transferred from the spring-like helix, as it
deforms, to the m a trix.
Takayanagi49 realized quite early that in fibre s
both series and parallel coupling are prese nt and hi s
unit c ube model has been very successfu l.
GUPTA & RADHAKRlSHNAN : NATURE OF CRYSTALS AND THEIR INFLUENCE ON FIBRE PROPERTIES
4.3 Effect on Deformation Mechanisms in Oriented Polymers
Fourier-transfonn infrared (FTIR) studies made
on oriented crystalline PET films prepared by
heat-setting at elevated temperatures were made as
the films were axially deformed s9 . It was observed
that in the sample heat-set under constraint, crystal
deformation commenced at a very low strain while
the deformation of the slack-set sample was
dominated by chain uncoiling in the amorphous
phase. Only beyond 20% strain, chain unfolding
could be detected in the slack-set sample. These
effects are a direct consequence of the nature of
coupling in the two samples; a strong parallel
coupling in the taut-annealed samples ensures the
participation by the crystallites in the defOlmation of
the sample right from the commencement of the test
while the predominant series coupling in the
slack-annealed sample results in the dominance of the
coiled amorphous regions in controlling the
structural changes which occur in the first 20%
extension.
4.4 Effect on Thermal Behaviour
Heat-set PET filaments subjected to DSC tests for
studying the melting characteristics, with the sample
in the unconstrained and constrained states, showed
some interesting features 32 ; the relevant data are
summarized in Table 4. It may be noted that the
free-annealed fibres , which have been annealed at the
higher temperatures, have lower crystallite and
amorphous orientation relative to the corresponding
taut-annealed samples, as expected. The crystal
disorder parameter, crystallinity and crystalli te size,
however, have been shown to be quite close in these
two sets of samples 3 2.5 3 with the free-annealed
samples having slightly higher crystallinity,
crystallite size and crystal perfection.
It is interesting to note that in the constrained state,
when reorganization of the structure is inhibited and
the orientation is expected to be retained right up to
melting, the melting point is enhanced by 6° to 9T
over the melting point observed in the unconstrained
state. Thi s enhancement was attributed to the
en tropic retrictions on chain conformation in the
amorphous regions. The coupling effects are quite
obvious when the data are examined in detail. The
crystal perfection, size and crystallinity are all slightly
higher in the free-annealed fibres (annealing
temperature, ·J 220-250·C) but their ' ·' melting
temperatures are lower. This could be attdbuted to
the greater degree of parallel coupling in these
samples and is a reflection of the higher orientation in
the taut-annealed samples .
Table 4-Structural and thermal characteristics of heat-set
PET filaments (ref. 32)
Sample
Crystallite Amorphous
orientation orientation
( P2(9)c ( P2(9)am
Control
Melting temp. ('C)
Unconstrained
technique
Constrained
technique
0.940
0.582
254.7
263 .5
0.952
0.947
0.947
0.934
0.937
0.936
0.944
0.519
0.494
0.470
0.433
0.416
0.369
0.318
254.6
254.9
255 . 1
255.5
256.6
256.0
256.2
263.8
263.8
263 .8
263.6
263.0
262.4
262:0
0.944
0.952
0.922
0.912
0.87,7
0. 859
0.860
0.522
0.505
0.430
0.258
0.200
0.156
0.039
255.6
255.0
255.2
254.8
254.8
255 .0
255.0
263 .7
262.0
263 .0
262.2
261.9
261 .8
261.4
Taut-annealed
100·C
140"C
180· C
220"C
230·C
240·C
250·C
Free-annealed
IOO·c
140·C
180·C
220·C
230·C
240·C
250·C
5 Crystal Content and Size
The most dramatic effect of crystallinity is
observed for isotropic polymers in the rubber-like
state . For example, the Young's modulus of raw
rubber (smoked sheet) increases by a factor of the
order of 100 by the presence of about 24 %
crystallinity 60. The rubbery modulus of amorphous
unoriented PET at 90·C also increases by about 100
times on crystallizationS . In both cases, the Tg also
increases on crystallization . The sit uation for
oriented, semi-crystalline fibres is more complex. For
example, when commercial PET fibre is heat-set
isothennally at temperatures between 100°C and
255T for 5 min , the Tg first increases with increase in
heat-setting tempera ture (or with crystallinity),
reaches a peak value at about 180°C and then shows a
decrease 3l , as shown in Fig. 17. The dye uptake or
diffusion coefficient of PET, say at 130°C, on the other
hand, first decrea ses with increasing temperature of
heat-setting and then registers an increase as shown in
Fig. 18 (ref. 62); the minimum in dye uptakecoinciding with the maximum in Tg (ref. 61) as
may be seen by comparison of Figs 17 · and 18.
These observations have been expla'inedas follows :
Up to heat-setting temperature of 180°C, new crystals
ate fonned by coming together of. ' well-para\leliied
chains in the amorphous regions which increase the
number of small crystals in the fibre. As a result, at a
111
INDIAN 1. FIBRE TEXT. RES., MARCH 1991
160
1400~~----------------------~
M
e-FA
o-'TA
"
~
.e 120
- 1000
01
><
~
0
UJ
~
800
::l
SO
0 100
140
ISO
220
260
-J
0
>
III
::l
HEAT-SETTING TEMP. (oel
Fig. 17- Variation of Tg of PET fibres with heat-setting
temperature
..
1200
70
80r---------------------------~
0
400
~
'200
a
~
-<
~
600
0
0
70
120 140
160 180
200 220
21,0
ANNEALING TEMP., ·C
Fig. 19--DependenIX of amorphous volume per crystal of PET on
heat-setting temperature (taken from ref. 63)
UJ
i5
It must be emphasized that the mechanical
properties of extended chain fibres like gel-spun
____ ____ __ ____ __
HDPE, Kevlar and carbon fibres are very sensitive to
120
140
160
180
200
220 240
crystal orientation: In the standard melt-spun fibres ,
SETTING TEMPERATURE;C
the mechanical properties are not equally sensitive to
crystal
orientation. This is because they are
Fig. 1~Effect of setting temperature on dye uptake [:LuI. OISp." ""
predominantly related to the amorphous orienta ..
Fast Scarlet B 150 at lOOT for 90 min] for ' PET fibre
(taken from ref. 62)
tion. For PET fibres, the strain in the crystalline
regions was shown 64 to be a factor of about ten less
heat-setting temperature of 180·C, the fibre contains than the macroscopic strain for macroscopic strains
a large number of small crystals so the amorphous up to l.8%.
regions are highly constrained. Heat-setting at higher
The orientation of the crystal or lamellar planes
temperatures can result in the formation of crystals of can also have a very large effect on the Young's
larger size. The fibre consequently contains a small modulus offibres. The earliest studies were made on
number oflarge crystals and the amorphous volume low-density polyethylene films 50 which showed that
per crystal increases; as a result, molecular mobility is inter-lamellar shear could make a significant
enhanced. Fig. 19 (ref. 63) shows this effect very contribution to the compliance of the film,
clearly.
particularly in samples showing four point low angle
Crystal size can have a significant effect on the . patterns with lamellar planes inclined to the fibre
melting point of fibres ; larger crystals showing higher aXlS.
melting temperatures than smaller crystals.
7 Crystal Formation
The fibre forming polymers generally crystallize
6 Crystal Orientation
The orientation of the molecules In the crystals when their melts or solutions are cooled, the
with respect to the fibre axis can have important crystallization being induced by nuclei. For spherical
consequences on mechanical and thermal properties nuclei of radius r, the Gibbs surface energy is
of the fibre. The measured crystal modulus of PET proportional to r2 while the crystallization energy,
reported in the literature ranges between 75 GPa and which is a volume effect, is proportional to r3 (ref. 26).
137 GPa. This wide range ofvalues has been shown 2 These energy terms are of opposite signs and so the
to arise because the angle between the crystalline energy of nucleus formation becomes negative only
c-axis and the fibre direction can vary in different above a critical nucleus size. The nucleus can then
samples which, if not accounted for, can lead to an grow into, for example, a spherulite.
underestimate of crystal modulus. When suitable
When polymeric materials are heat-set,
correction was applied, a value of 110 GPa was crystallization can occur either by the melting of
smaller imperfect crystals at the crystallization
obtained 2 .
40
JO~~~
112
~
~
~
~
~
GUPTA & RADHAKRISHNAN: NATURE OF CRYSTALS AND THEIR INFLUENCE ON FIBRE PROPERTIES
temperatures followed by recrystallization into a new
form or by solid state transfonnation in which
molecules in the amorphous regions, which are not in
exact register, can, with the aid of thennal energy,
gain enough inobility to reduce their free energy by
fonning small crystallites. In some interesting work
on PET, Wu et al. 65 monitored the low-angle X-ray
scatterIng patterns in situ as the films defo'rmed and
concluded that the latter process predominates.
Fakirov et al. 30 have also arrived at the same
conclusion from DSe studies. However, Groeninckx
and coworkers 66 •67 state that the samples heat-set at
low temperatures have small imperfect crystallites,
which reorganize by partial melting and
recrystallization. In samples heat-set at relatively
higher temperatures , no large-scale melting occurs;
the crystal thickening which occurs without any
change in long period is presumed to arise from
crystal perfectio n at the boundary laye rs of the
crystalline and amorphous phases, a process which
has been briefly referred to earlier in this paper.
Another process which has received considerable
attention is the high-speed melt spinning process
which involves uniaxial stress fields in the spin-line;
the orienting influence of the stress field overcomes
the disorienting influence of thennal relaxation
process 68 .69 . During melt-spinning, the temperature
at which crystal formation can commence is
considerably enhanced at higher spinning speeds 70 ;
however, it is not raised to a high enough level to
prevent molecular relaxation. This has been a cause
of disappointment to the fibre producers who had
dreams of manufacturing oriented crystalline fibres
in one step.
A novel discovery reported recently71 holds a
glimmer of hope for the revival of interest in one-step
standard fibre production process·for flexible chain
polymers. The studies reported are on high molecular
weight HDPE which is unprocessable at l600C but
was found to have minimum viscosity between 150
and l52°e in which range it fonned a highly mobile
hexagonal crystalline mesophase making it possible
to use the liquid crystal spinning with this flexible
system which so far has been used only with rigid
systems like Kevlar, aromatic copolyesters and
cellulosc and perhaps which the silk worm has been
using for ages to spin silk, the queen of fibre.S.
The spinline crystallization of PET at high speeds
involves significantly large undercooling; the 'melt
temperature may decrease from 280 e to 180~C iJi
0.005 s. The nucleation rate is consequently very high
and the cancent'r ation of nuclei is very dense. Tht
critical size of the nucleus for crystallization to
proceed is considerably reduced and app roaches the
unit cell dimension of PET . This occurs because in an
oriented melt, the entropy change on crystallization is
small. At a given supercooling the free energy change
is correspondingly increased. This has been termed as
nucleative collapse 72 or as spinodal defect
decomposition 73 and is also referred to as regime III
of crystallization. The large number of small nuclei
come together to fonn a crystallite and the
crystallization is distinctly different compared to the
usual nucleation and growth process that dominates
the crystallization of a melt in the stress-free state or at
low stresses .
There is another type of crystallization that ha s
received
considerable
attenti0I;l,
viz.
the
crystallization of ultra-high molecular weight HDPE
to give highl y crystalline, extended chain
morphology (Fig. 20a) (ref. 74). A similar model has
been proposed for Kevlar (Fig. 20b) (ret. 8). As
discussed in an ea rlier section, these high
performance fibres are produced from a gel or a
solution. Mackleyand Sapsford 75 have made the
interesting observation that high performance
H D PE fibres had been earlier prod ueed by Ward and
coworkers 1 and Wu and Black 76 also from the melt; in
( b)
( a)
"I
0
•• '
'.: I "
I.
::!:,,::
t..c
•
. ": ~.,
'::
..
I '.' . .
, ;
Fig. 20-(a) Extended chain morphology of a ..gel.-~P\ln ni gh
, mq l.~~phH ,w,t;jiW-\.,liDP.I;: PR~~~ .Cry~t~llini~, >.j.O.%
.
(taken fr~ ~e~, ,(4) ~,.: '1.' I . ' . .
, (b) Scnt;matic represeo.tation of llU"uctun: 'of Kevlar fibre
'. ,(lake n from,ref. 8)
113
INDIAN J. FIBRE TEXT. RES., MARCH 1991
both the cases, high-strength, high-modulus fibres
could be obtained by processing at the elevated
temperatures of the order of 250°C and then rapidly
cooling the filament to give an essentially isotropic
fibre . The fibre is then subsequently drawn in the
temperature range 80-12SOC and draw ratios of the
order of30 can be achieved. Mackley and Sapsford 75
offer a physical exp!anation as to why these
processing conditions should give such a high draw
ratio. They believe that the processing conditions
used resulted in a reduction in the entanglement
concentration in the precursor melt and filament. As
in gel drawing, this meant that the fibre could then be
drawn to a high draw ratio. It is concluded that the
polymer must be processed in a manner to provide an
entanglement concentration high enough to ensure a
connected network, yet sufficiently low to enabfe high
draw ratios; whilst increased molecular weight
provides enhanced strength.
A structural model for high modulus polyethylene
has been derived from entanglement concepts by
Grubb 77. Starting from the concept that the
entanglement network is a controlling factor in
polymer deformation, a fibrillar morphology, rather
than a crystalline-bridge type of morphology, similar
to that derived from solution grown Shish Kebab
material, is proposed.
I7
18
19
20
21
22
23
24
25
26
27
28
29
30
31
32
33
References
34
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