speed melt spinning with nano silicates - CiteSeerX

Original article

Modification of the mechanical propertiesof polyamide 6 multifilaments in high-speed melt spinning with nano silicates

Wilhelm Steinmann1, Stephan Walter1, Thomas Gries1,Gunnar Seide1 and Georg Roth2

Abstract

In the last few years research activities have been focused on the modification of fiber properties with nano-scaled

additives. One of the most important fields of research is the alteration of mechanical properties such as the tenacity and

the specific breaking load. In this study, we determined the influence of nano-phyllosilicates on the drawability of

polyamide 6 multifilament yarns. It was first demonstrated that the drawability of the fibers drastically increased in an

industrially relevant high-speed melt spinning process. Structural properties of the material are identified by wide-angle

X-ray diffraction (WAXD) and differential scanning calorimetry (DSC). Changes in the crystalline properties as well as in

the alignment of the silicates are compared with the stress–strain curves of the fibers, and a molecular mechanism for the

drawing process is derived from these experiments. In a first step, a significant phase transition in the crystalline structure

unaffected by the silicates occurs for low draw ratios (DRs). Beyond this point, where unmodified fibers start to break,

a gliding between the silicate layers takes place, which is responsible for an extended drawability of the fibers.

This mechanism leads to new possibilities for fiber processing, which can be used to research ultra-fine filaments in

future studies.

Keywords

melt spinning, multifilament yarn, nanocomposites, polyamide 6, structure analysis

Polyamide 6 (PA6), also called Polycaprolactam orPerlon as a trademark, is a polymer of the group ofpolyamides. It was first synthesized by means of ring-opening polymerization reactions from "-aminocapro-lactam in 1938.1 The ‘6’ in the name denotes thenumber of carbon atoms in the repeating base mol-ecule. The resulting sequence in the polymer chain is�½HN� ðCH2Þ5 � CO�n�.

2

The polymer is a polymorphic, semi-crystallinematerial. Three different thermodynamic, metastablephases can coexist in the glassy or rubbery solidstate.3,4 The so-called g phase is crystalline with amonoclinic unit cell. The lattice parameters areknown as a¼ 9.33 A and c¼ 4.78 A perpendicular tothe polymer chains as well as b¼ 16.88 A in the polymerchain direction from wide-angle X-ray diffraction(WAXD) experiments. The value of the angle �between a and c axis is 121�.5 Since g is around 120�,a pseudohexagonal version of this unit cell can be

calculated and is often used in theory for the structuraldescription.4 The second crystalline phase denoted asthe a phase also has a monoclinic cell with lattice par-ameters a¼ 9.56 A, b¼ 17.24 A, and c¼ 8.01 A with anangle �¼ 67.5�.6 As distinguished from the a and gphase, the � phase as the third stable phaseis characterized by a mesomorphic structure. Theshort-range order exceeds the typical values for anamorphous material. Since no long-range order exists

1Institut fur Textiltechnik der RWTH Aachen University (ITA), Aachen,

Germany2Institut fur Kristallographie der RWTH Aachen University (XTAL),

Aachen, Germany

Corresponding author:

Wilhelm Steinmann, Institut fur Textiltechnik der RWTH Aachen

University (ITA), Otto-Blumenthal-Straße 1, 52074 Aachen, Germany.

Email: [email protected]

Textile Research Journal

82(18) 1846–1858

! The Author(s) 2012

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DOI: 10.1177/0040517512456756

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in this phase, no detailed structure can be obtainedfrom diffraction experiments. However, a molecularmodel, based on a nano-crystalline, hexagonal phasewith a¼ 4.8 A and c¼ 8.6 A was proposed. The sidegroups in this structure model are randomly oriented.7

Characterization and phase transitions

Information on the properties of all phases is predomi-nantly gathered from diffraction experiments. Since allreflections from the lattice planes perpendicular to thepolymer chains appear in the same angular range, peakfitting is carried out for an identification of thephases.3,4 Both g and a form crystals appearingtogether with the � phase if a sample is cooled fromthe melt. The a formation is preferred at low coolingrates.3,4,8 Pure � phase material can be obtained byquenching.3,7

Thermal analysis as well as temperature-dependentdiffraction experiments reveal information about phasetransitions. The g crystallites arise from the � phase attemperatures higher than 130 �C,3 whereas g is trans-formed to a phase crystallites in the temperature regimeof the melting point, which is located at 220 �C.9 Incontrast, examples of unstable g crystals, which trans-form far below the melting point,10–12 have beenreported. Also � phase crystals are able to transformto a phase. Therefore, annealing at 190 �C is neces-sary.13 Since polymer chains are more stretched in theg phase, these chains’ formation is also preferred instrain-induced crystallization under mechanical stresswith the result that this phase is preferred after thedrawing processes.8 The g phase is gradually trans-formed to the a phase during annealing at temperatureshigher than 160 �C.3

Melt spinning

Melt spinning is the classic method of fiber productionof thermoplastic materials.2 PA6 is one of the mostcommon fiber polymers, where the processing of mono-filaments and multifilament yarn has a long tradition.The material can be melt-spun as low oriented yarn(LOY) and partially oriented yarn (POY), whereashigh-spinning speeds can be achieved. The materialcan be converted to fully drawn yarn (FDY) in a sepa-rate process as well as to inline drawn yarns (IDY) inone step by melt spinning.2,14

The fiber formation process of PA6 has been widelyexamined with theoretical calculations and online mea-surements of fiber properties. In this case, the fiberdiameter and velocity, mechanical forces, filament tem-perature, and development of molecular orientation

have been investigated.14–16 Other studies deal withthe structure formation in the melt spinning process.The development of crystal phases was examined within-situ measurements of diffraction patterns, whichshowed the qualitative development of the threephases.17–20

Nanocomposites

Nanocomposites of PA6 were characterized in manystudies. Several types of particles at the nanoscale,including carbon nanotubes, carbon nanofibers, silicondioxide, and many silicate-based particles, were dopedin the polymer matrix.21–49 The composites are eitherproduced by in-situ processing during polymerizationor by compounding a nano-powder in the liquid statewith a twin screw extruder.22–29

Much research has focused on the modification ofpolyamide with layered silicates (or phyllosili-cates).22–49 These sheet-like structures are known forthe improvement of fire retardancy, the modificationof mechanical and electrical properties, and influen-cing the sorption.27,30–34 The influence on the polymerstructure and its physical properties were examined byWAXD, differential scanning calorimetry (DSC),transmission electron microscopy (TEM), nuclearmagnetic resonance (NMR), Fourier-transform infra-red spectroscopy (FTIR), and viscosimetry. In theseexperiments, the silicate concentration, the molecularweight of the polymer, and the silicate type werevaried.4,28,32,34–45 The presence of the silicates seemsto affect the crystallization behavior, the occurringcrystal phases, and the crystallite sizes.28,35–42

Further structural changes occur in the annealing pro-cess.46 In addition, the silicates are supposed to sup-port the formation of a rigid amorphous phase(RAF).47

The use of silicates in fibers was mostly examined inelectrospun nanofibers, where fiber diameter is in thesame range as the silicate dimensions. The alignment ofthe particles was observed with TEM.48–50 Only a lim-ited amount of research projects on the effect of nano-silicates on the properties of filament yarns have beencarried out.8,28,51 One group detected changes in thecrystal geometry.8 Another study deals with the fiberproperties of yarns spun at different take-up velocitieswith different silicate concentrations. The presence ofsilicates leads to an increase in the density and the crys-tallinity, where mechanical properties mostly remainunchanged.51 Another group worked on the structuralchanges in the fibers, where nanocomposites with exfo-liated silicate sheets were obtained by in-situ prepara-tion. The monofilaments were spun at very low take-upvelocities and showed a reduced drawability with higherclay content.27

Steinmann et al. 1847



Aim of this study

Current research activities at the Institut furTextiltechnik (ITA) focus on the processing of newpolymer materials to fibers as well as the modificationof standard polymeric materials with dopants fromorganic and inorganic materials. These activities aimat the functionalization of fibers with new propertiesand the enhancement of fiber properties. In thisstudy, we examined the influence of melt-compoundednano-phyllosilicates (or layered silicates) on themechanical properties and drawability of PA6 multifi-laments in a high-speed melt spinning process withindustrially relevant winding speeds. Various experi-ments have been carried out to determine the influenceof these particles and to understand the changes at amolecular scale. According to the crystallographic andthermal properties, we present a model for the extendeddrawability of the multifilaments.

Experimental details

Material

For the melt spinning of the multifilaments, a standardpolymer B24N03 from BASF AG, Germany was used,which is suitable for high-speed spinning. According tothe datasheet of the material, the melting temperature is220 �C and the density at room temperature is about1.14 g/cm3.52

Silicate-modified compounds of the material areproduced with the help of a twin screw extruder.Two different master batches with a total mass fractionof 2% and 1% of phyllosilicates are created.Commercially available phyllosilicates, Nanofil 5 fromRockwood Clay Products, Inc., USA, are used as nano-clays.53 These silicates are based on a natural bentonitethe most important component of which is montmor-illonite. The surface of the silicate is treated by adimethyl, di(hydrogenated tallow)alkyl ammoniumsalt.53 As a consequence, Na+ ions are exchangedwith an organic component, allowing for a betterembedding in the polymer matrix. The typical distancebetween two layers is about 2.8 nm.53

Melt spinning

Melt spinning is carried out on the spinning plant ITAPlus, which allows a broad variation of many spinningparameters at an industrial scale. The polymer in gran-ule form is being transported into a single-screw extru-der from Oerlikon Barmag GmbH & Co. KG,Germany, where the polymer is compressed andmelted through friction and increasing pressureamong the granules. The melt exits the extruder at

pressure levels around 100 bar and is transportedthrough heated pipes to the spin pump. Up to thespin pump the process and throughput is pressure con-trolled depending on the lead pressure at the extruderhead. From the spin pump is a throughput-controlledprocess, where the mass flow rate Q is adjusted to44.0 g/min. The polymer is pumped into the spinpack, which consists of a spinneret and multiple filtra-tion layers. These can be woven metal structures, cera-mic, or metallic nonwovens, and also stainless steelsand is commonly used in melt spinning. After passingthrough all filtration stages, the melt flows into thecapillaries of the spinneret. For the melt spinning ofPA6, a spinning plate with a number of n¼ 24 capil-laries is chosen. The geometry of the capillaries isdefined by their diameter (D¼ 250 mm) and theirlength (L¼ 500mm). Together with the mass flow rateQ and the density of the polymer �, the geometry alsodefines the extrusion velocity ve, where A is the totalarea of capillaries:

ve ¼Q

� � A¼

Q

� � n � �D

2

� �2ð1Þ

At the end of the capillary, the melt exits and forms afiber. The temperature of the extrusion equipment andthus of the melt is the major process parameter deter-mining the viscosity and shearing of the polymer. Thetemperature in the extruder is chosen as 265�C andslightly increased in the heated pipes for lowering theviscosity. The extrusion temperature is chosen as265�C. Over a length of 1500mm the filaments arecooled by a laminar, conditioned air stream at 20�Cwith a velocity of 0.35m/s perpendicular to the fiberaxis and 75% relative humidity. In this step, the solidi-fication and crystallization of the fibers take place. Theas-spun fibers are drawn down vertically usually by atleast one draw down godet rotating at a surface velocityvG. Thus, the draw down ratio (DDR) is fixed asthe ratio between godet velocity vG and extrusionvelocity ve:

DDR ¼vGve

ð2Þ

Depending on the use of the filaments, a wide varietyof drawing states can be realized inline before windingthe filaments on bobbins. The solid-state drawing of thefibers is conducted between two godets, where the con-secutive godet revolves faster than the first. In thisstudy the spinning is realized with inline drawingusing a variation of the drawing ratio between 1.0185and 1.8, whereas the drawing temperature is fixed. Thefilaments are winded on bobbins with a winder from

1848 Textile Research Journal 82(18)



Oerlikon Barmag GmbH & Co. KG, Germany,whereas the winding speed is fixed at 5000m/min. Theparameters used for melt spinning and inline drawingof the multifilaments are summarized in Table 1.

WAXD

Diffraction experiments are carried out on a single-crys-tal diffractometer STOE & Cie IPDS II equipped withan image plate for digital readout. Two independentgoniometer circles allow the exact positioning of thesample for detecting the reflections of the main latticeplanes from the PA6 crystal structures. A special speci-men holder guarantees the parallel alignment of thefibers and a reproducible measurement of the orienta-tion distribution of crystallites in the fiber.

Molybdenum Ka radiation with a wavelength of�¼ 0.71073 A was chosen for the experiments. Theexperimental setup and the specimen holder are dis-played in Figure 1.

WAXD is used to detect changes in the lattice para-meters, the crystallite dimensions and the orientationdistribution due to the phyllosilicates. Further analysisof the diffraction patterns is done with the X-rayAnalysis Tool developed at ITA. This tool allows theparameterization of the recorded two-dimensional pixelgraphics in directions 2y (diffraction angle) and j (azi-muthal angle). The diffraction angle is calculated fromthe radial position r’ with the help of the known detec-tor distance d¼ 200mm:

2� ¼ arctanr 0

d

� �ð3Þ

With the help of this data, the independent integra-tion of different regions in the diffraction pattern inboth directions 2y and j is possible. Numerical calcula-tions and peak fits on the resulting intensity diagramsare carried out with the OriginPro software fromOriginLab, USA.

DSC

DSC measurements are carried out on a Mettler ToledoDSC 1 at the ITA polymer laboratory, which isequipped with a FRS5 sensor having 56 thermocouplesfor precise heat flow measurements. For all experi-ments, a heating rate of 10�C/min and a nitrogen atmo-sphere was chosen. The fibers were cut into pieces witha total sample weight of 10mg. The samples wereplaced into 20 ml aluminum crucibles and pressed tothe bottom of the crucible with a lid for good heatcontact. An automatic sample changer was used toplace the crucibles on the DSC sensor for reproducible

Figure 1. Experimental setup (left) and specimen holder (right).

Table 1. Process parameters and variations for melt-spinning of

polyamide 6

Parameter Values

Capillary geometry

(number of filaments

� capillary diameter

� capillary length)

24� 0.25 mm� 0.5 mm

Extrusion temperature 265�C

Mass flow rate 44.0 g/min

Extrusion velocity 25.7 m/min

Draw down ratio 191.0/159.3/134.0/117.7/107.9

Inline draw ratio 1.0185/1.220/1.450/1.650/1.800

Godet temperature

(godet 1/godet 2)

40�C/160�C

Winding speed 5000 m/min

Mass fraction phyllosilicates 0%/1%/2%

Quench air (velocity/

temperature/humidity)

0.35 m/s/20�C/75 %




measurements. Melting enthalpy is calculated by usingthe received thermograms. The melting temperaturebetween fibers in the first and quiescent crystallizationin the second heating process is also determined.Furthermore, the glass transition and correspondingrelaxation processes are examined. The results areused for the identification of changes in the crystalphases a and g as well as for the determination of thetotal crystallinity.

Mechanical properties

A Statimat 4U from Textechno GmbH, Germany wasused to determine the stress–strain curve of all materialswith a clamping length of 100mm and an extensionvelocity of 100mm/min under the standard condition.These curves are used for the calculation of the tensilestrength and the elongation at break. The mechanicalproperties are compared with the structural data deter-mined by the other experimental techniques. This com-bination reveals specific molecular mechanisms for themodification of the mechanical properties due tophyllosilicates.

Results

Crystalline properties

In Figure 2, typical diffraction patterns of the PA6fibers are displayed and are representative for all mea-surements. The small-angle scattering reflections(2y& 2�) appear on the equator of the patterns onlyat high DRs for fibers with phyllosilicates.

For further analysis, intensity profiles are calculatedfor selected regions. Background intensity due to air

scattering and noise from the readout process is sub-tracted. The regions in the diffraction patterns and thecorresponding intensity profiles are displayed in Figure 3.

The 2y profiles can be used for the determination ofthe unit cell and crystallite dimensions meridional andparallel to the fiber axis. Equatorial profiles contain onereflection from the two lattice planes (200) and (001) ofthe g phase (indices for pseudohexagonal description ofthe crystalline form) and (200) as well as (002) of the aphase. The meridional profiles (parallel to the fiber) axiscontain one peak from the (020) lattice plane. For thedescription of the crystalline intensity distributions, aLorentz peak model is chosen, where the center Braggangle 2yc, the half width �2y, and the total intensityI0 are refined:

Ið2�Þ ¼ Io �1

2��

�2�

2� � 2�cð Þ þ�2�

2

� �2ð4Þ

The distance (d) between the crystalline lattice planesis calculated from the center diffraction angle and thewavelength with the help of Bragg’s equation:

d ¼�

2 � sin2�c2

� � ð5Þ

For all fiber samples, the d values for (200), (020),and (001) correspond to well-investigated values of theunit cell of g phase PA6 (4.79 A in the pseudohexagonaldescription)5,54 and do not change with the processparameters. Therefore, they cannot be used to evaluatethe effect of the spinning process.

Figure 2. Typical diffraction patterns of PA6 fibers (left: 0% of phyllosilicates with a DR¼ 1.22, right: 2% of phyllosilicates with a

DR of 1.8).




A qualitative value for the crystallite dimensions (D)is derived from the peak position and half width withthe help of the Scherrer’s equation:

D ¼�

�2� � cos2�c2

� � ð6Þ

Further evaluation is carried out for diffraction pat-terns with additional intensity distributions at smalldiffraction angles, where two different interpretationsare possible. In the first case, the intensity is consideredsmall-angle scattering, where the slope d of the intensitydistribution provides information about typical sizes ofregions having other densities than the polymer in thefiber structure and follows Guinier’s law:55

IðQÞ ¼ ��2N2pV

2p exp �

Q2�2

3

� �ð7Þ

The variable q is the norm of the scattering vector,which can be determined from the diffraction angle:

Q ¼4�

�� sin

2�

2

� �ð8Þ

The other parameters are the density differencebetween void and polymer �r as well as the numberNp, and the volume of the density fluctuations Vp. Sincethe reflections appear on the equatorial direction of thepatterns, these voids are perpendicular to the fiber axis.The resulting distances are interpreted as spacesbetween polymer chains due to the embedding of sili-cate layers. The corresponding parameter d (Guinierslope) is determined by a linear regression, where thelogarithmic intensity is plotted against the square of thescattering parameter Q2.

The second interpretation is based on the shapetransformation of the silicate crystal morphology,where the intensity distribution is a partially visibleBragg reflection of the silicate crystal lattice. A smalleramount of stapled silicate layers lead to an increase ofthe full width at half maximum (FWHM). The size Ds

and number of stapled silicates is then calculated withthe help of the Scherrer equation. Both interpretationslead to similar values for the number of embedded sili-cate layers between the polymer chains and the sameadjacent molecular mechanism.

The azimuthal scans provide information on theorientation distribution of the crystallites (fc) as wellas the orientation of the phyllosilicates (fs). The meandeviation from the peak center is used for the calcula-tion of the orientation factor f ranging from –1/2 for

Figure 3. Regions in diffraction patterns and corresponding intensity profiles: (a) scans in radial direction, (b) scans in azimuthal

direction (intensity distribution in region 3 corresponds to region 1 in (a) and vice versa).




samples oriented in the other direction (not possible forfibers) to 0 for unoriented samples up to 1 for perfectorientation:56

f ¼ 1�3

25 sin2ð’Þ4 ð9Þ

In Table 2, an overview of the crystalline propertiesof all samples is given.

Thermal properties

Typical DSC thermograms with and without phyllosili-cates as well as from quiescent crystallization (secondheating process) are displayed in Figure 4. Calculatedvalues for the glass transition temperature Tg, the corre-sponding change in heat capacity �cp, the melting tem-perature Tm and the melting enthalpy �H are indicated.

With increasing DR from 1.0185 to 1.22, a signifi-cant change in the heat flow curves takes place for bothtypes of fiber. First, a relaxation process can beobserved at the glass transition for the less drawnfibers, which also appears in the second heating process.The second change is a modification of the peak formfor the higher ratio together with the appearance of ashoulder at lower temperatures. For the fibers withphyllosilicates allowing higher DRs (1.8), no significantchange in the thermograms takes place. For all fibers,as well as for quiescent crystallization, the changes inheat capacity �cp¼ (0.20� 0.02) J/gK at the glass tran-sition Tg¼ (49.0� 0.4)�C as well as the meltingenthalpy �H¼ (72.7� 0.4) J/g at the melting tempera-ture Tm¼ (220.0� 0.6)�C remain constant. The degreeof crystallinity is calculated from the fraction of thetheoretical value for pure crystalline material of PA6:57

Xc ¼�H

�H0¼

�H J=g½ �

190:6� 100 % ð10Þ

The thermal properties of all samples are summar-ized in Table 3.

Mechanical properties

Stress strain curves of fibers with DR 1.22 with andwithout phyllosilicates are displayed in Figure 5. It isobvious that the silicates affect the mechanical proper-ties of the fibers, whereas the elongation at break isincreased and the tensile strength is reduceddramatically.

The values for the tensile strength Fmax and the elon-gation at break Emax are summarized in Table 4. Inaddition, the values for the fineness T are displayed.Since extrusion velocity and winding speed were con-stant for all materials, the fineness is nearly the same forall fibers.

Discussion

Crystallinity, phase contents and morphology

The DSC analysis of the multifilaments showed anearly constant melting enthalpy in all samples, which

Table 2. Crystalline properties of the PA6 fibers with different mass fractions of silicates and draw ratios revealed from WAXD

Sample 1 2 3 4 5 6 7 8

Silicates 0 % 0 % 1 % 2 % 2 % 2 % 2 % 2 %

Draw ratio (DR) 1.0185 1.22 1.22 1.0185 1.22 1.45 1.65 1.8

Lattice spacing perpendicular dII [A] 8.47 8.45 8.52 8.52 8.46 8.49 8.48 8.45

Lattice spacing parallel DII [nm] 20.2 17.9 18.5 19.8 18.8 17.6 17.7 17.5

Orientation factor of crystallites (perpendicular) fc,r 0.85 0.91 0.89 0.86 0.92 0.93 0.91 0.92

Orientation factor of crystallites (parallel) fc,II 0.86 0.90 0.88 0.87 0.91 0.90 0.92 0.93

Orientation factor of silicates fs – – – – 0.88 0.89 0.92 0.93

Spacing due to silicates from Guinier slope d [nm] – – – – 2.35 2.15 1.85 1.49

Spacing due to silicates from FWHM Ds [nm] – – – – 2.45 2.11 1.67 1.35

Figure 4. DSC thermograms (with and without phyllosilicates,

exothermal heat flow on positive axis).




also occurs in the second heating process. This indicatesthat the total crystallinity is not affected by the processconditions chosen for the production of the PA6 fibers.Furthermore, the phyllosilicates do not seem to affectthe crystallization behavior of the material. Comparedwith other studies showing a nucleation of g phase onthe silicates,8 this emphasizes the importance of stress-induced crystallization in the high-speed spinning pro-cess. The comparison of the fiber samples with thesecond heating leads to the conclusion that the fractionof crystalline material is completely unaffected by theextreme cooling conditions in the melt spinning pro-cess. Although the total crystallinity Xc& 38.1% isconstant all over the processing conditions, the formof the DSC thermograms changes significantly. Theappearance of the shoulder before the melting peakfor higher DRs indicates a transition phase to a greateramount of the crystalline a phase during the heatingprocess. This conversion to a phase during heat

treatment is special for polymeric fibers since thereverse transition can be observed in samples with noorientation (e.g. in compression molding).57 The tem-peratures for the shoulder observed in the experimentscorrespond to a phase transition to a high-temperatureform similar to a phase known from another study.8

This transition occurs by increasing the DR from1.0185 to 1.22, whereas the peak form remains constantfor higher DRs in the silicate doped samples. Thisobservation indicates that the g to a transformationoccurs in the early stages of drawing. In this step,only the polymer chains are affected by the mechanicalstress of the inline drawing process. At higher DRs, thethermal response of the polymer is constant, so thatfurther drawing is assumed to affect the silicates.However, the corresponding changes in the silicatesystem are not detectable due to the small mass fractionin the polymer matrix.

The structural phase transition is accompanied bytwo other effects. First, the relaxation process duringthe glass transition of the amorphous fraction is sup-pressed when the material has been transformed to thea phase. Since glass transition occurs in the noncrys-talline regions, this different nature of the processcannot be generated directly by the crystal phases.However, it is assumed to be caused by a change inthe morphology, where amorphous and crystallineregions are arranged differently. Like the transforma-tion in the crystalline regions, this effect is unaffectedby the presence of phyllosilicates in the polymermatrix. The second effect is a slightly smaller valueof the crystallite size parallel to the fiber axis. Thisoccurs in combination with the increasing g to a trans-formation. Therefore, this observation also providesan indication of changes in the crystal morphology,whereas in this case the crystallite geometry (detect-able in the crystallite size) is changed. Like the otherobservations affecting the crystal phases, this change isalso independent of the presence and amount of sili-cate dopants.

Table 3. Thermal properties of the PA6 fibers with different mass fractions of silicates and draw ratios revealed from DSC

(Q: quiescent crystallization, second heating)

Sample 1 2 4 5 6 7 8 9 Q

Silicates 0 % 0 % 1 % 2 % 2 % 2 % 2 % 2 % 0 %

Draw ratio (DR) 1.0185 1.22 1.22 1.0185 1.22 1.45 1.65 1.8 -

Glass transition temperature Tg [�C] 48.4 49.7 49.3 48.9 48.5 49.1 48.7 49.3 48.8

Change in heat capacity �cp [J/g�K] 0.19 0.19 0.21 0.22 0.18 0.2 0.18 0.23 0.21

Melting temperature Tm [�C] 219.3 220.5 221.1 219.5 220.3 219.9 219.2 220.8 219.5

Melting enthalpy �H [J/g] 73.3 73.0 72.1 72.3 73.0 72.8 73.1 72.5 72.3

Crystallinity Xc [%] 38.5 38.3 37.8 37.9 38.3 38.2 38.4 38.0 37.9

Figure 5. Stress strain curves of fibers with DR 1.22 (with and

without phyllosilicates).




Orientation distribution and dispersion of silicates

Information about the orientation distribution isobtained from the evaluation of the intensity profiles.Since small angle scattering in the patterns due to voidsis assumed to be an effect of the silicates embedded inthe polymer, the corresponding intensity distribution isused for the determination of the placement of silicatesin the fiber. The orientation factor f of the crystallites aswell as the silicates is displayed in Figure 6a. The crys-talline regions are well oriented in all fibers, but thisorientation increases even more together with the g toa transformation. For higher DRs, the factor remainsalmost constant. Since values for f with different massfractions of silicates are very similar, the dopant doesnot seem to affect the reorientation. Instead of that, theeffect is also assumed to take place due to a significantchange in the morphology of the fiber structure, wherethe fibrils of the a phase have a stronger orientationalong the fiber axis.

For the lowest DR, no small-angle scattering signalcan be detected, so that no orientation information canbe obtained. In this case, the silicates are assumed to bemore or less disoriented respective to the fiber axis so

that no orientation occurs by the drawing of the melt.For higher DRs in the solid state, the silicate orienta-tion increases continuously and reaches values as highas the crystalline ones. Since small-angle scattering isobserved in the equatorial plane of the diffraction pat-terns, the normal to the silicate layers defined by dis-tance to the next silicate sheet is perpendicular to thefiber axis.

Further information on the embedding of the silicatelayers in the polyamide matrix can be obtained fromthe slope d determined in the Guinier plot described inthe experimental section, or alternatively from theFWHM of the reflection and the Scherrer formula.Both values represent a medium distance of polymericmaterial, whereas this void is filled with silicates.Consequently, the parameter can be used for the calcu-lation for an average thickness of silicates in the matrix.With the help of the typical distance between two sili-cate layers, the corresponding number of adjacentlayers can be evaluated, which is a benchmark for thedispersion of silicates in polyamide. As seen in Figure6b, the Guinier slope d decreases with increasing DR.The final value for the highest DR 1.8 (1.49 nm) corre-sponds to the dimension of one silicate layer (with

Figure 6. (a) Orientation factors of crystallites and silicates dependent to DR, (b) Guinier slope d as a function of DR.

Table 4. Mechanical properties of the PA6 fibers with different mass fractions of silicates and draw ratios revealed from stress strain

curves (fineness T from reeling)

Sample 1 2 3 4 5 6 7 8

Silicates 0 % 0 % 1 % 2 % 2 % 2 % 2 % 2 %

Draw ratio (DR) 1.0185 1.22 1.22 1.0185 1.22 1.45 1.65 1.8

Tensile strength Fmax [cN/dtex] 5.68 5.75 4.06 3.08 3.19 3.72 4.05 4.31

Elongation at break Emax [%] 49.2 46.6 65.4 76.9 73.6 63.8 55.8 51.8

Fineness T [dtex] 91.4 91.4 91.2 91.2 91.4 91.3 91.5 91.5




regard to the organic component, which is compatiblewith the polymer and does not cause voids). Therefore,a defoliation of silicates takes place in the drawing pro-cess and results in single silicate layers dispersed in thepolymer matrix. The same effect of silicate defoliation,remarkable in smaller voids in the polymeric material,was detected by Ibanes et al.8,58 Compared with thisstudy with winding speeds less than 1000m/min, theeffect was demonstrated here for high-speed melt spin-ning (5000m/min). Furthermore, the material wasdrawn at lower temperatures (40�C compared with140�C). Both variations cause a slower phase transfor-mation since kinetics is slower at lower temperaturesand the transition has less time to take place at higherspeeds. Therefore, the effect of a phase formation is lessdetectable here. It is remarkable that a complete defo-liation of the silicates could be achieved in an indust-rially relevant spinning process.

Drawability

As seen from the mechanical characterization, the ten-sile strength decreases with an increasing amount ofphyllosilicates, whereas the elongation and thereforethe drawability increase drastically. These results arein contrast to prior work, where nanocomposites werecreated by in-situ polymerization and silicate layerswere already dispersed. In this case, the maximum elon-gation was reduced.27 Thus, this seems to be an effect ofmelt compounding, where silicates are first agglomer-ated in the granule. Since dispersion takes place in thedrawing process, drawability can only be increasedwhen the nanocomposites are prepared by melt com-pounding. With increasing DR, the tensile strengthincreases together with a decreasing value for the elon-gation but never reaches the higher values for pure

PA6 filaments. The dependency of both mechanicalproperties from the mass fraction of phyllosilicates andthe DR is almost linear. Compared with the study ofIbanes et al.,8,58 tensile strength in particular is muchhigher here. This is the effect of the high-speed spinningprocess, showing the possibility of an industrial applica-tion of silicates for achieving higher drawability.The corresponding diagrams are shown in Figure 7.

Molecular interpretation

With the help of information discussed in the previousparagraphs, a molecular interpretation of the drawingprocess is possible. For both pure and nano modifiedPA6 fibers, the mechanical stress in the process firstaffects the polymeric matrix. In this step, a significantchange in the crystal morphology takes place, causing astructural phase transition from the crystal phase g toa, a better molecular orientation, and smaller crystal-lites in the lamellar fiber structure. For the PA6 nano-composites, a first orientation of the silicates takesplace. In this step, the silicate layers are more or lessagglomerated as revealed from the higher values for theGuinier slope. Beyond the point of break for pure PA6multifilaments, extended drawability is caused by thepresence of silicates in the case of agglomeratedsheets. Mechanical stress acts on the silicate layers,where weak binding forces between two layers allow agliding and a defoliation of stapled silicate sheets. Thesebinding forces are stronger than the conformationaland structural energy saved in the polymer matrix,since the process takes place at higher DRs withincreased mechanical stress. However, they areweaker than the forces acting between a polymerchain and an adjacent silicate layer. The other wayaround, no gliding would be possible since no

Figure 7. Tenacity in dependence of mass fraction of phyllosilicates and DR.




mechanical forces would be transmitted from the poly-mer to the silicate system. Weak binding forces can alsoexplain the decrease of tensile strength. For the highestDR, the defoliation is fulfilled and single-layer silicatesare well dispersed and oriented in the polyamidematrix. Since no further orientation and gliding pro-cesses can take place, the fiber breaks after this point.The complete transformations at the molecular scaleoccurring in the drawing process are exemplifiedin Figure 8.

Conclusion

The presence of phyllosilicates in PA6 multifilamentsleads to an increase of the drawability in the inlinedrawing process. For achieving these properties, thepreparation should be done by melt compoundingsince agglomerated silicate sheets are necessary. Inthis study, it was first demonstrated that this effectcan be used in an industrially relevant high-speed spin-ning process at 5000m/min. For lower DRs, only thePA6 polymer chains are affected by the mechanicalstress. This effect is independent of the silicates andleads to a higher orientation of the polymer chains aswell as a transformation in the crystalline regions. Thistransformation is assumed to be the conversion fromthe g to the a phase. For higher DRs, beyond the elon-gation of break for pure PA6, a defoliation of the sili-cate layers takes place. Drawing of the fibers is possibleuntil single silicate layers are dispersed in the polymermatrix. Compared with other studies with lower wind-ing speeds, the structural phase transition from g to a isless dominant, but the effect of the silicates is larger. Inthis case, the doping of PA6 with nano phyllosilicatescould be used to reduce the fiber fineness and to pro-duce ultrafine filaments. Even though mechanical prop-erties are lower than for unfilled PA6, they are still

good enough for an industrial application of these fila-ments. For this purpose, further experiments with areduced mass flow rate have to be carried out.Furthermore, the reduction of tensile strength has tobe compensated with an appropriate procedure of pro-cessing the fibers.

Funding

Contract grant sponsor: Deutsche Forschungsgemeinschaft(German Research Foundation, DFG); research project‘‘FiberSage’’, contract grant number: GR 1311/34-1.

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