morphology and structure of constituent fibres in thermally bonded nonwovens

Morphology and Structure of Constituent Fibres in

Thermally Bonded Nonwovens

Xiao Yan Wang,* Rong Hua Gong

Textiles and Paper, School of Materials, The University of Manchester, P.O. Box 88, Sackville Street,Manchester, M60 1QD, UKE-mail: [email protected]

Received: November 11, 2005; Revised: February 7, 2006; Accepted: February 8, 2006; DOI: 10.1002/mame.200500385

Keywords: crystal structure; fibres; morphology; nonwovens; thermal bonding

Introduction

Nonwoven fabrics are made directly from fibres in a

continuous production line, thus partially or completely

eliminating conventional textile operations, such as car-

ding, roving, spinning, weaving or knitting.[1,2] The

simplicity of fabric formation, coupled with high produc-

tivity, allows nonwovens to compete favourably with

Summary: A new process has been developed to producethree-dimensional nonwovens directly from staple fibres. Inorder to establish suitablewindows of the process parametersto achieve high-quality nonwoven products, the effects ofthermal bonding temperature, dwell time andmould materialon the morphology and structure of the fibre have beeninvestigated using PP/PET bi-component fibres. It wasevident from both scanning electron microscope imagesand Raman spectra that thermal-induced shrinkage of the PPsheath fibre occurred in the thermal bonding process, leadingto deformation and cracking of the PP sheath and exposure ofthe PET core. X-ray diffraction results revealed crystalimperfection and/or less ordered polymer chains, more g-

form and thermal contraction of the crystal lattice for the PPsheath fibre, while birefringence measurements indicatedthat both the birefringence and the orientation factor for thePP fibre decreased after the thermal bonding process. Thedegrees of the thermal-induced shrinkage increased, andthe crystallinity, birefringence and orientation factor of thePP sheath fibre decreased with increasing thermalbonding temperature, dwell time and thermal conductivityof the mould material. All these can be attributed to thedifferent levels of modification of chemical compositioncaused by thermal oxidative degradation and thermal-induced relaxation of the orientation during the thermalbonding process.

Changes of morphology and crystalline features of PP/PET fibre afterthermal bonding process.

Macromol. Mater. Eng. 2006, 291, 499–509 � 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Full Paper DOI: 10.1002/mame.200500385 499

wovens and knits on a performance per cost basis in many

industrial applications, from simple low-cost replacements

for more expensive textiles to high-quality textiles. They

also offer many functions that cannot be filled by traditional

fabrics.[1–3]

Through air thermal bonding offers process versatility,

high product uniformity and loft, and a clean production

process. It is the only thermal bonding process that allows

the entire product to be exposed to a uniform temperature.[4]

Also, through-air bonded nonwovens can be processed

without chemical binders, offering a much safer production

process and working environment for machinery operators

and plant employees. Furthermore, manufacturers are able

to save energy and achieve lower operating costs, as no

binder preparation station is required.[4,5]

Recently, a pilot process to produce three-dimensional

(3D) nonwoven products directly from staple fibres has

been developed in The University of Manchester.[4–8] In

this process, the fibres are formed into a 3D webs using air-

laying and the 3D webs are bonded using through-air

thermal bonding. It has been reported that the most suitable

fibres for this process are poly(propylene)/polyester (PP/

PET) bi-component fibres.[4–8]

During the thermal bonding process, the fibre webs are

subjected to the temperature above or close to the melting

point of the PP sheath of the PP/PET bi-component fibre

in the bonding chamber. Earlier studies[9,10] have shown

that the polymer structure may be changed dramatically

when they are subjected to heat treatment. It is therefore

very likely that the thermal bonding process resulted in

changes in the morphology and structure of the PP/PET bi-

component fibre. The morphology and structure of

polymers are very important for their many properties,

includingmechanical and optical properties,[9,11] and hence

affect the properties of the corresponding fabrics consisting

of these polymer fibres. The knowledge of morphology and

structure of the PP/PET bi-component fibres before and

after the thermal bonding process is important for establish-

ing suitable windows of the process parameters to optimize

the quality of the nonwoven products.

Many techniques can be used for the characterization of

the morphology, crystal structure and chain orientation

states of polymers.[9,12–16] The most widely used techni-

ques are infrared dichroism, X-ray diffraction (XRD) and

birefringence. Each of these techniques has its advantages

and disadvantages. The structural features of the molecule

produce characteristic and reproducible absorption bands in

the vibrational spectrum of the molecule. As such, the

Raman spectrum and infrared spectrum can be used as a

fingerprint for identification of structural features by

comparing the new spectrum with previously recorded

reference spectra.[17] The drawback of this technique is that

there are difficulties in knowing whether individual

absorption peaks provide information about crystalline or

highly-ordered chains only, or give some average over both.

X-ray scattering studies of the polymer, which may be

either uniaxially oriented fibres or bulk specimens, can give

wide angle information about the crystal structure,[9] the

crystallinity and themorphology of samples, such as degree

of crystallinity, crystallite size and imperfection and

crystalline modifications. This technique can, in principle,

give the complete distribution of orientations for the

crystallites but suffers from difficulties in correcting the

background scattering due to amorphous materials and

other effects. Molecular orientation in polymers is usually

linked to optical birefringence.[15,18] The observation of

birefringence is therefore a common experimental tool for

the investigation of chain orientation effects in polymers.

Birefringence can be determined in two ways. The first

consists in determining the refractive index in three

directions with the aid of a refractometer, and then

calculating the birefringence as the difference between

different refractive indices. The second method measures

the birefringence directly using an optical method with

polarized light (from the retardation in polarized light on

going through an oriented sample). This method is simple

and quick; moreover, it can get sufficient information of the

chain orientation of the sample.

In this paper scanning electron microscopy (SEM),

Raman spectroscopy, XRD and birefringence measure-

ments with a polarizing microscopewere used to character-

ize the morphologies, crystal structures and chain

orientation states of the as-received fibre and the fibres in

the thermally bonded nonwoven fabrics produced under

different process conditions. Following this paper, the

relationships between the morphology and structure of the

constituent fibre and the properties, including abrasion

resistance and tensile strength, of the nonwoven fabrics will

presented elsewhere.[19] The main objective of the present

study is to understand the process-structure-property

relationships of the thermally bonded nonwoven products

made directly from the commercial PP/PET (sheath/core)

bi-component staple fibres.

Experimental Part

Sample Preparation

The 3D nonwoven process includes two stages: web forming(Figure 1) and thermal bonding, and has been described indetail elsewhere.[4,5]Once the 3Dweb is formed on themoulds,they are moved out of the mould chamber across the machinewidth into a bonding section for consolidation. In the bondingchamber, the hot air is drawn through the fibrous web that issupported on the original mould. The hot air inlet is connectedto a hot air reservoir and duct heater through flexible adiabaticpipes. The air outlet is connected to a suction fan. An air guideis used to improve the flow distribution around the web. Theposition of the air guide can be adjusted along the central axis.

Using various shapes ofmould, a variety of 3D fabrics can beproduced. In the present study, commercial PP/PET (sheath/

500 X. Y. Wang, R. H. Gong

Macromol. Mater. Eng. 2006, 291, 499–509 www.mme-journal.de � 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

core) bi-component staple fibrewith properties listed in Table 1was used to make the 3D nonwoven samples. Different sets ofprocess parameters and mould materials were employed toobtain three series of samples, as listed in Table 2, where thebonding temperature varied from 144 to 159 8C, the dwell timefrom 1 to 5 s, and the mould materials included copper, steeland nylon.

Characterization of Fibre Morphology and Structure

A SM-300 SEM was used to observe the morphologies of theas-received bi-component PP/PET fibre and the fibres in thethermally bonded nonwoven fabrics. Their Raman spectrawere further measured to characterize the relative degrees ofPP sheath shrinkage and PET exposure of the bi-componentfibre after the thermal bonding process. The measurementswere carried out on a Nicolet Raman 950 spectrometer,together with the Raman software, the Nicolet’s OMINICTM,to collect and process theRaman data at room temperature. Thespectra were recorded by plotting intensity against Ramanshift, which is a measure of the difference between theobserved spectra bands and thewavelength of radiation used bythe exciting laser, fromaround400–4 000 cm�1 and an average500 scans at a resolution of 4 cm�1.

The XRD spectra of both the as-received fibre and those inthe thermally bonded nonwoven samples were measured atroom temperature to identify their phase constituents anddetermine their average crystallite size and lattice strain. TheXRD spectra were collected using a Philips X’pert APD,PW3710 diffractometer (Philips Co., Netherlands) withCu-Karadiation. The experimental conditions included a Cu target,operating at 50 kVand 40 mA, a 0.25 mm automatic divergentslit and a 0.2 mm receiving slit. The collected data angles (2y)ranged from 38 to 758 with a step size of 0.058 and a countingtime of 32.0 s per step.

It is well known that small crystal size, e. g. smaller than100 nm, or crystal imperfection and/or less ordered polymer

chains would result in the broadening of the XRD diffractionpeaks.[20] In the present study, the average grain size and thelattice strain of the nonwovens were estimated based on XRDpeak broadening according to the well-established Scherrermethod:[21]

b ¼ B� b ¼ 0:9ld cos y

ð1Þ

Lattice strain ¼ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiB2 � b2

p

4 tan yð2Þ

where b is the corrected full width at half-maximum of thediffraction peak at the diffraction angle 2y, and B and b are thepeak widths of the sample to be analyzed and the standardsample, respectively. l is thewavelength of the radiation, and dis the average grain size. In the present study, the diffractionpeak of the (117) plane at 2y� 218was used, and the correctionof the full width at half-maximum of the diffraction peak wasachieved by a standard silica specimen to remove theinstrument broadening.

Furthermore, polymers generally exist in partially crystal-line form. In such a case, there is a broad amorphous halo in thecorresponding XRD spectrum, in addition to those diffractionpeaks of the crystalline constituents. According to Machadoand Denardin,[22] the relative crystallinity is related to the areavalues and given by:

xc ¼Ac

Ac þ Aa

ð3Þ

where Ac and Aa are the total areas of the crystalline peaks andthe amorphous halo on the diffraction spectrum, respectively,both without the contribution of the background. The areas ofthe crystalline peaks were obtained by subtracting both thebackground and the amorphous halo from the originaldiffraction spectrum.[22]

The optical birefringencevalues of both the as-receivedfibreand those in the thermally bonded nonwoven sampleswere alsomeasured using a Leitz polarizing microscope at roomtemperature to determine the modification of the molecularorientation of the fibre after the thermal bonding process. Priorto the measurements, a few fibres were laid parallel to one

Figure 1. The 3D web forming system.

Table 1. Properties of the as-received PP/PET bi-componentfibres.

Lineardensity

Fibrelength

Breakingload

Breakingelongation

Breakingtenacity

dtex mm gf % gf � tex�1

1.9 51 9.5 40 0.5

Table 2. Sample codes and the corresponding bondingparameters.

Samplecode

Mouldmaterials

Bondingtemperature

Fabricweight

Dwelltime

Air velocity

8C g �m�2 s m � s�1

BTS1 Steel 144� 1 55� 1 3� 0.5 4� 0.3BTS2 Steel 150� 1 55� 1 3� 0.5 4� 0.3BTS3 Steel 154� 1 55� 1 3� 0.5 4� 0.3BTS4 Steel 159� 1 55� 1 3� 0.5 4� 0.3DTS1 Steel 150� 1 55� 1 1.0� 0.5 4� 0.3DTS2 Steel 150� 1 55� 1 2.0� 0.5 4� 0.3DTS3 Steel 150� 1 55� 1 3.0� 0.5 4� 0.3DTS4 Steel 150� 1 55� 1 4.0� 0.5 4� 0.3DTS5 Steel 150� 1 55� 1 5.0� 0.5 4� 0.3DTN3 Nylon 150� 1 55� 1 3� 0.5 4� 0.3DTC3 Copper 150� 1 55� 1 3� 0.5 4� 0.3

Morphology and Structure of Constituent Fibres in Thermally Bonded Nonwovens 501


another and then cut through using a new razor blade, held at458 to the axes of the fibres. Cutting the fibres in this way gavewedged-shaped ends; then the fibres were mounted in liquidparaffin on a microscope slide and covered with a cover glassfor measurement. The number of whole wavelength bands, asseen in monochromatic light, was counted up the wedge. Thequarter wave or Senamont compensator measured the partwavelength of interference.

During the measurements, the fibre was positioned in thediagonal position on the stage and an interference filter insertedinto the light path. A series of dark bands were seen in the fibre.The number of bands was noted, and then the analyser wasrotated until the centre two fringesmet at the centre of the fibre.The angle of rotation was directly proportional to thedifference between the path difference due to the fibre andthe next lower whole number of wavelength. To obtain greateraccuracy, the fibre was rotated through 908 and the measure-ments were repeated. Fresh samples were used for all tests, and20 repetitions were carried out for each condition. Averagevalues from the 20 repetitionswere used for further calculation.Then the birefringence was calculated by:[23,24]

Dn ¼ lt

aþ b

360þ N

� �ð4Þ

whereDn is birefringence, l iswavelength of light used, a andbare the anglesmeasured,N is the number of wholewavelengthsor number of bands and t is fibre diameter.

The orientation factor (f) can be simply defined in terms ofthe angle f between the polymer chains and the elongateddirection of a fibre, as:

f ¼3 cos2 f� �

� 1

2¼ Dn

D0ð5Þ

where the brackets indicate an average of over all angles andD0

is the absolute birefringence:

D0 ¼ XcD0c þ ð1� XcÞD0

a ð6Þ

where Xc is the crystallinity, Dc0 and Da

0 are the intrinsicbirefringence values of the crystalline and amorphous phases,respectively. In the present study, the latter two values used themost widely cited values, Dc

0¼ 0.0291 and Da0¼ 0.060,[9,25]

although values of Dc0¼ 0.0416 and Da

0¼ 0.0379 were alsoreported in the existing literature.[9]

Results and Discussion

Sheath Shrinkage of the Bi-Component Fibre

Figure 2 and 3 present the SEM images showing the

morphologies of the as-received PP/PET bi-components

fibre and those in the thermally bonded nonwoven samples.

The as-received bi-component fibre has relatively uniform

cross section, while deformation and cracks of PP sheath

and exposure of PET core caused by thermal-induced

shrinkage of PP sheath are visible for the fibres after the

thermal bonding process. This suggests that a flow ofmatter

to reduce the surface energy of the fibres probably occurred

during the thermal bonding process. However, as shown in

Figure 3, the deformation and cracks of PP sheath and

exposure of PET core in the nonwoven samples are local

and it is difficult to correlate them directly with the process

conditions through SEM observations. Therefore, the

Raman spectra of the samples were used to compare the

degrees of shrinkage of PP sheath and the exposure of PET

core of the fibres.

As shown in Figure 4–6, the Raman spectra indicate that

there appear two new absorption peaks from the C C

double bonds (1 620 cm�1) andC O carbonyl bonds (1 730

cm�1), for the thermally bonded nonwoven samples

compared to the as-received fibre sample. Moreover, the

heights of the two absorption peaks are clearly associated

with the thermal bonding time, bonding temperature and

mould materials. They becomemore pronounced due to the

increased concentrations of the function groups with

increasing thermal bonding temperature, bonding time

and thermal conductivity of themouldmaterial (in the order

of copper> carbon steel> nylon). The two new peaks of

the nonwoven samples come from the PET core.[23,24,26]

Figure 2. SEM images of the as-received bi-component fibresshowing the uniform cross sections: (a) lower magnification; (b)higher magnification.



These results reveal that the PET core had been exposed to

some extent after the thermal bonding process, and the

degree of exposure increased with increasing thermal

bonding temperature, bonding time and thermal conduc-

tivity of the mould material.

Trznadel[27] and Pakula[28] reported that when PP

polymer with fixed molecular orientation is subjected to

increased temperature, without being subjected to external

mechanical forces, it shrinks due to the relaxation of

orientation of the extended polymer chains. They also

Figure 3. SEM images of the fibres in the thermally bonded nonwovens: (a) BTS1, whole view; (b) BTS1, local view showing acrack on the surface of PP sheath; (c) BTS4, whole view; (d) BTS4, local view showing deformation and shrinkage of PP sheath andexposure of PET core; (e) BTS4, local view showing severe cracks on the surface of PP sheath and exposure of PET core; (f) BTS4,another local view showing relatively weak cracks on the surface of PP sheath.



pointed out that the maximum shrinkage of highly oriented

polymer fibres occurs at temperatures very close to the

melting point or slightly above it. During the present

thermal bonding process, the fibre was heated above the

melting point of the PP sheath, so the shrinkage of the PP

sheath is significant. As for the PET core, itsmelting point is

significantly higher than that of the PP sheath, so the

molecules in the PET corewere lessmobile than those in the

PP sheath. Accordingly, the amount of shrinkage for PET

corewould be far lower and it is negligible compared to that

for the PP sheath, thus easily leading to the exposure of the

polyester core. Therefore, the cracks at the surface of PP

sheath and two new peaks in the Raman spectra of the fibres

in the thermally bonded nonwoven samples, due to the

exposure of the PET core, can be attributed to the far lower

shrinkage of the PET core than that of the PP sheath.

The thermally induced shrinkage of highly oriented

polymers is dependent on temperature and time.[27,28] The

chain mobility of a fibre increases on increasing the

temperature. Thus, the shrinkage of the PP sheath increases

with increase in the bonding temperature, leading to the

increased exposure of the polyester core, as shown in

Figure 3 and 4. Although the chain mobility has a limit at a

certain temperature and hence the shrinkage of a specific

fibre has a maximum at each temperature, the shrinkage of

the fibre PP sheath increased with increase in dwell time, as

shown in Figure 5. This indicates that the dwell times used

during the thermal bonding process were insufficient for the

chain mobility of the PP sheath fibre to reach its limit.

In the thermal bonding process, when the fibrewebs were

moved into the bonding chamber, a mould with higher

thermal conductivity would be rapidly heated up to higher

temperature than a mould with lower thermal conductivity,

given an approximately identical dwell time. Consequently,

when the samples left the bonding chamber, the more

thermally conductive mould would continue to heat

these samples for a longer time than the less thermally

conductive mould would do. Therefore, the amount of fibre

shrinkagewas in the order of copper> carbon steel> nylon

mould (Figure 6). This is in a good agreement with the

dependence of the thermal degradation of the nonwovens

on the mould materials.[29]

As shown in Figure 3, the exposure of the PET core

mainly resulted from the shrinkage and cracking of the PP

sheath. As aforementioned, there is a considerable differ-

ence in the amount of shrinkage exhibited by the fibre

sheath and core. Wagner[30] reported that the different

shrinkage between the core and the sheath of the thermal-

plastic-based bi-component fibres can lead to axial thermal

Figure 4. Raman spectra of the as-received fibre and those in thethermally bonded nonwoven samples produced at differentthermal bonding temperatures.

Figure 5. Raman spectra of the as-received fibre and those in thethermally bonded nonwoven samples produced using differentdwell times.

Figure 6. Raman spectra of the as-received fibre and those in thethermally bonded nonwoven samples produced using differentmould materials.



stresses in the order of 10 GPa, due to the interfacial

adhesion between the fibre PP sheath and the PET core. In

the present work, DSC curve reveals that the melting

temperature of the PP sheath of the as-received fibre is

around 129.5 8C.Upon start of the thermal bonding process,

fibre oxidative degradation gradually developed.[29] Oxy-

gen grafting to PP macromolecules led to an increased

density of the fibre surface layer, while the elimination of

the volatile compounds resulted in the decreased weight.

These may cause further shrinkage of the PP sheath and

hence increase the axial stresses. Also, the oxidation was

highly heterogeneous, thereby creatingmechanically active

defects: oxidized domains.[31] Once the final sheath length

could not accommodate the core length, cracks would thus

initiate and develop around these oxidized domains on the

PP sheath, as shown in Figure 3.

Crystallinity and Crystal Structure of the PP Sheath

Figure 7 presents the XRD spectrum of the as-received bi-

component PP/PET fibres, with the appropriate separations

between the crystalline, amorphous and real background

sections. Figure 8–11 present the XRD spectra of the fibres

in the thermally bonded nonwoven samples produced using

the different process conditions, in the range of 2y¼ 0–358.The data of the crystallite size, lattice strain and crystal-

linity, calculated using Equation (1), (2) and (3), for these

samples are listed in Table 3.

FromFigure 7–11 andTable 3, it can be seen that both the

crystallinity and the crystal structure of the PP sheath of the

bi-component fibre had changed after thermal bonding

process. Some peaks in the spectrumof the as-received fibre

sample disappear, and a few new peaks, specially the (113)

peak of the g-form, appear for the fibres after the thermal

bonding process. The crystallinity values of the fibres in the

nonwoven samples are far lower than that of the as-received

fibre sample, and they decrease with increase in the thermal

bonding temperature, dwell time and thermal conductivity

of the mould material. Due to thermal contraction of the

crystal lattice resulting from thermal-induced shrinkage,

the corresponding peaks for the nonwoven samples are

shifting to larger angles compared to the as-received fibre

sample. The angle shift is also influenced by the thermal

bonding process and mould materials. It increases, and

accordingly the lattice strain decreases, with increasing

bonding temperature, bonding time and thermal conduc-

tivity of the mould material. These results are in good

agreement with the results from the Raman spectra, as

shown in Figure 4–6.

The crystallization behaviour of PPfibre is very complex.

So far, four different crystal forms (a, b, g, and smectic) for

PP have been identified mainly using XRD.[32–37] The four

crystal forms have different stability under a certain

condition. In most cases, the g-form is slightly more stable

than the a-form.[38–40] Its content generally increases with

decrease in the molar mass of the PP fibre molecule, in the

Figure 7. XRD profile of as-received fibre, with the appropriateseparations between the crystalline, amorphous and real back-ground sections.

Figure 8. XRD profiles of the as-received fibre sample and of athermally bonded nonwoven sample.

Figure 9. XRD profiles of the fibres in the thermally bondednonwoven samples produced using two different bonding temper-atures.



range of 1 000–3 000 g �mol�1.[33,41] As mentioned in a

previous paper,[29] the thermal oxidative degradation of

fibre PP sheath during thermal bonding process followed b-scission mechanism, which resulted in chain breakage and

led to decreased molar mass. Earlier studies[34–37,42] have

indicated that the g-form crystals develop from the short

isotactic propylene sequences under similar conditions.

From the XRD spectra of both the as-received fibre and

those in the nonwoven samples, only a and g forms are

detected (Figure 8–11). Therefore, it is reasonable to

conclude that more g-form was probably formed during the

thermal bonding process. The g-form (113) peak newly

appearing on the XRD spectra for the fibres in the thermally

bonded nonwoven fabric samples can be ascribed to the

increased content of this form of crystals.

The g-form crystals formed by short crystalline sequence

usually have lower melting temperature than that of a-formcrystals formed by long crystalline sequence.[34] Thus, the

increased content of g-form crystals is one of the reasons

why the melting point of the fibre PP sheath had decreased

after the thermal bonding process, as determined by the

DSC curves.[29]

That the crystallinity decreases with increasing thermal


of the mould material can be explained as follows. The

crystallinity drop was probably associated with the melting

states of the PP sheath subjected to the hot air flows at the

thermal bonding temperatures that are higher than its

melting temperature of around 129.5 8C as revealed by

DSC. During the thermal bonding process, when a lower

bonding temperatures or a shorter dwell time was used,

some molecular order may be present in the melts which

had not been heated much above the crystalline melting

point. Such regions of order would provide sites for crystal

nucleation and hence crystallization would be more rapid

when cooling was carried out, resulting in a higher

crystallinity. On the other hand, the thermal oxidation of

the PP sheath during the thermal bonding process resulted

in broken chains and led to the reduction in chain length.[29]

The short chains of PP can decrease the nucleation and

crystallization kinetics of PP from melt.[42] Also, the

content of the oxygen-containing groups became high

enough to render more difficulties to the inter-chain

association, thereby leading to crystallinity decreased for

the fibres in the thermally bonded nonwoven samples. The

level of thermal oxidation increases with increase in the


of the mould material.[29] Thus, the nucleation and

crystallization kinetics of PP sheath from the melt would

decrease with increase in the bonding temperature, dwell

time and thermal conductivity of the mould material. This

agrees very well with the observed result that the

crystallinity decreases on increasing the bonding temper-

ature, dwell time and thermal conductivity of the mould

material.

The broadening of the XRD peaks for the fibres in the

thermally bonded nonwoven samples can be more

explained as an indication of crystal imperfection and

disorientation of the polymer molecular, rather than small

crystal size. This is due to the fact that the degrees of

broadening of the different peaks are inconsistent with each

other. Most changes of crystal perfection involve reorgan-

ization of the surface of macromolecular crystals, which

contain chain folds, loops, ends and tie molecules.

Chemical alteration of the PP macromolecular crystals

surface resulting from the thermal oxidative degradation of

the PP during the thermal bonding process may cause

crystal imperfection. The crystal imperfection can also

reduce the melting temperature of the PP sheath, and is

another factor that caused the melting peaks of the PP in

DSC curves of nonwoven samples to shift to lower

temperatures 126.5–128.8 8C from 129.5 8C for the as-

received fibre.[29]

Figure 10. XRD profiles of the fibres in the thermally bondednonwoven samples produced using two different dwelling times.

Figure 11. XRD profiles of the fibres in the thermally bondednonwoven samples produced using two different mould materials.



Fibre Birefringence

Figure 12 presents the optical microscope images of the

fibres under polarized light. A series of dark bands are

observed and the number of bands is three for both the as-

received fibre and the fibres coming from the thermally

bonded nonwoven samples. The number of the whole

wavelength bands does not change; only the distance

between the two centre fringes decreases after the thermal

bonding process. A previous study[43] indicated that the

birefringence of PP is far lower compared with that of PET.

As such, it can be assumed that thewholewavelength bands

result from the PET core, and the PP sheath only causes part

wavelength interference. The number of whole wavelength

keeping as a constant can be ascribed to the fact that the

thermal bonding process lasted for a very short time and the

thermal conductivity of polymer is very low. Consequently,

no appreciable change of crystalline structure occurred and

thus the birefringence value of the PET core did not change

after thermal bonding process. As aforementioned, the

crystallinity and crystalline structure of the PP sheath

significantly changed after the thermal bonding process.

The decreased distance between the two centre fringes can

be easily attributed to the decrease in the birefringence of

the PP sheath.

As described in the Experimental Part, the birefringence

and orientation factor of the fibres can be calculated using

Equation (4) and (5) respectively, based on the parameters

obtained from their optical microscope images under

polarized light, the same as those in Figure 12. The

calculated values of the birefringence and orientation factor

for both the as-received fibre and the fibres in the thermally

bonded nonwoven samples are summarized in Table 3 and

Figure 13–15, where the error bars stand for the 95%

confidence limits of the average birefringence values. Here

Table 3. The measured parameters of crystal structure, birefringence and chain orientation factor for the fibres in the different samples,calculated using Equation (1)–(6).

Sample Crystallinity Crystallitesize

Latticestrain

Birefringence Orientationfactor

% A %

As-received fibre 63.4 135 1.67 0.024 0.596DTC3 33.2 230 1.09 0.009 0.224DTS3 36.8 205 1.11 0.016 0.330BTS4 32.7 232 1.08 0.009 0.013DTS5 35.3 223 1.11 0.011 0.291DTN3 38.8 184 1.21 0.018 0.377

Figure 12. Optical microscopy images of the fibres underpolarized light: (a) the as-received fibre and (b) a representativeof the fibres in the thermally bonded nonwoven samples.

Figure 13. Birefringence of the fibres in the different thermallybonded nonwoven samples as a function of the thermal bondingtemperature.



the effects of the thermal bonding temperature, dwell time

and mould material have been taken into account. It can be

observed that both the birefringence and the orientation

factor have decreased after the thermal bonding process.

Similar to the dependence of crystallinity of the PP sheath

on the process conditions, both the birefringence and the

orientation factor decreasewith increasing thermal bonding

temperature, dwell time and thermal conductivity of the

mould material.

The birefringence has been proved to be a sensitive and

reliable fingerprint of polymer structure.[11,44] Birefrin-

gence of a polymer fibre stems from two basic features in its

construction: (i) the chemical composition of the constit-

uent molecules and (ii) the structural arrangement of the

molecules within the fibre.[45] The chemical composition of

fibre is partly responsible for the birefringence level of the

fibre. The molecular structure, side groups and their

relationship to one another are important and any

alternation to the molecule or side groups can cause the

change of the birefringence to some extent. As for the

structural arrangement, the following features, including

the orientation of the polymer chain, crystallinity and

density, phase constituent and intrinsic molecular config-

uration, all affect the birefringence.

It was found that the chemical composition of fibres was

changed after the thermal bonding process due to the

oxygen-containing groups produced by the thermal oxida-

tive degradation.[29] As mentioned before, the crystallinity,

phases and perfection of crystals were also changed after

the thermal bonding process. The significant shrinkage of

the fibre PP sheath occurred due to the thermally induced

relaxation of the orientation after the thermal bonding

process. All these could attribute to the observed fall in the

birefringence of fibres after the thermal bonding process.

The thermally induced shrinkage and the level of the

oxidative degradation progressively increased and the

crystallinity decreasedwith increase in the thermal bonding


mould material. These were accompanied by a correspond-

ing progressive decrease in overall polymer chain orienta-

tion, which is consistent with that the birefringence of the

fibre decreased with increasing the thermal bonding


mould material.

Conclusion

The effects of the thermal bonding temperature, dwell time

and mould material on the morphology and structure of the

PP/PET bi-component fibre in the thermally bonded

nonwoven samples have been studied. Based on the results

obtained, the following conclusions can be drawn:

(1) Thermal-induced shrinkage of PP sheath caused

deformation and cracking of the PP sheath and exposure of

PET core after the thermal bonding process. The cracks

initiated and developed around the oxidized domains

formed due to the thermal oxidation of the PP sheath. The

degree of the thermal-induced shrinkage of the PP sheath

and the exposure of PET core increased with increasing

thermal bonding temperature, dwell time and thermal

conductivity of the mould material.

(2) Only a and g forms for PP fibre are detected from both

the as-received fibre and those in the nonwoven samples,

and more g-form was formed after the thermal bonding

process. The crystallinity of the PP fibre decreased with

increasing thermal bonding temperature, dwell time and the

thermal conductivity of themouldmaterial. The broadening

of XRD peaks of the fibres in the thermally bonded

nonwoven samples revealed crystal imperfection and/or

Figure 14. Birefringence of the fibres in the different thermallybonded nonwoven samples as a function of the dwell time.

Figure 15. Birefringence of both the as-received fibre and thefibres in the different thermally bonded nonwoven samplesprepared using different mould materials.



less ordered polymer chains induced by the thermal

bonding process. Due to the thermal-induced shrinkage,

thermal contraction of the crystal lattice was evident for the

PP sheath fibres after the thermal bonding process.

(3) Both the birefringence and the orientation factor

decreased after the thermal bonding process. Similar to the

dependence of crystallinity of the PP sheath on the process

conditions, both the birefringence and the orientation factor

decrease with increasing thermal bonding temperature,

dwell time and thermal conductivity of the mould material.

(4) The dependence of thermal-induced shrinkage,

crystallinity and birefringence of the PP sheath on the

process conditions can be attributed to the different levels of

modification of chemical composition caused by thermal

oxidative degradation and thermal-induced relaxation of

the orientation after the thermal bonding process.

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morphology and structure of constituent fibres in thermally bonded nonwovens

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