Molecular Orientation and Fine Structure Development in Polymer Fiber Formation.

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ABSTRACT

ARVIDSON, SARA ASHLEY. Molecular Orientation and Fine Structure Development in

Polymer Fiber Formation. (Under the direction of Drs. Saad A. Khan and Russell E. Gorga).

Understanding structure‐property relationships is key to developing useful

polymeric materials. Developing specific structures in polymer fibers requires precise

control over processing conditions and well understood characterization techniques. We

consider melt spun fibers of polypropylene, which have been in development for decades

but which are still poorly understood in many regards. We use differential scanning

calorimetry and wide angle x‐ray diffraction to study first, what processing conditions result

in formation of the mesomorphic form of isotactic polypropylene, and then, how the

mesomorphic‐to‐α‐crystalline phase transition in polypropylene is affected by processing

conditions, such as extruding the polymer into fibers versus quenching the polymer. We

evaluate the mesomorphic structure for its influence on fiber strength as well as discuss the

duality of mesomorphic phases, which are by definition neither fully crystalline nor fully

amorphous in nature. We also correlate the mechanical properties of polypropylene fibers

with the molecular orientation within fibers to develop a method for estimating molecular

orientation quickly without using methods which are not readily available and/or have

limited applicability.

Developing novel materials and/or properties though alloying polymers provides

additional challenges; incompatible materials suffer from poor adhesion, while complete

mixing can result in loss of the individual, desirable properties of each polymer.

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However, it requires the use of petroleum‐based raw materials and is difficult to

functionalize. Poly(lactic acid), in comparison, requires significantly less energy per pound

to produce, results in the release of less greenhouse gas, and can be composted, but often

costs more, is more brittle, and made degrade prematurely during storage or use. To solve

some of the issues associated with both polymers, we fabricate fibers of polypropylene‐

poly(lactic acid) where the two polymers are contacted in the molten state to form

core/sheath filaments. By contacting rather than blending the polymers, the potential

exists where the properties of each polymer can be combined to produce a fiber with the

benefits of both polymers with the negatives of neither. We find that the stress‐strain

profiles, obtained by applying a constant rate of extension, depend on the polymer that is

the sheath component. The stress‐strain profiles are compared to those obtained from

fibers of either polypropylene or poly(lactic acid) spun individually to evaluate the relative

molecular orientation developed in the sheaths of bicomponent fibers.

As more engineered materials require the properties of multiple components, the

use of “compatibilizers,” or agents which act to reduce the surface tension between

components, continues to become more important. While polypropylene and poly(lactic

acid) spin well together, the mismatch in mechanical properties (such as elongation‐to‐

break or modulus) of the polymers may result in composites that may be prone to

delaminate. We investigate the use of a block copolymer, poly(styrene‐block‐ethylene‐co‐

butylene‐block‐styrene) (SEBS) to compatibilize the polypropylene‐poly(lactic acid) interface

in core/sheath fibers by melt blending it with the poly(lactic acid) in the extruder prior to

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concentrate at the interface and affect the morphology of the polymers at the interface in

such a way that the interfacial layer contains both polymers entangled with each other. We

discuss the morphology of the block copolymer in fibers, in blends, and in solution cast films

and how altering the processing history of the SEBS can results in interesting morphologies.

We also evaluate the ability of the copolymer, added to poly(lactic acid), to influence the

mechanical properties of the polypropylene‐poly(lactic acid) core/sheath fibers.

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Molecular Orientation and Fine Structure Development in Polymer Fiber Formation.

by

Sara Ashley Arvidson

A dissertation submitted to the Graduate Faculty of

North Carolina State University

in partial fulfillment of the

requirements for the degree of

Doctor of Philosophy

Chemical Engineering

Raleigh, North Carolina

2011

APPROVED BY:

_______________________________ ______________________________

Saad A. Khan Russell E. Gorga

Co‐Committee Chair Co‐Committee Chair

________________________________ ________________________________

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ii DEDICATION

This thesis is dedicated to my parents,

in loving gratitude,

for their lifelong support and encouragement.

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iii BIOGRAPHY

Sara Arvidson was born in Wilmington, NC to parents Robert and Brenda. She was

raised in Cheraw, SC with her sister Morgan. After graduating from high school, she

attended University of South Carolina in Columbia, SC. Upon completion of her Bachelor of

Science in Chemical Engineering in 2005, she worked as a Production and Validation

Engineer at 3V, Inc. in Georgetown, SC. Sara joined the Department of Chemical and

Biomolecular Engineering at North Carolina State University in 2006 and the group led by

Prof. Saad Khan in 2007. She was co‐advised by Prof. Russell Gorga in the department of

Textile Engineering Chemistry and Science in the College of Textiles. Following completion

of her graduate studies, she will join the groups of Frank S. Bates and Timothy P. Lodge as a

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ACKNOWLEDGMENTS

Primarily, and with deep gratitude, I must acknowledge my advisors, who have made this

endeavor possible. I am grateful to Dr. Khan for his daily involvement, positive spin, the

atmosphere he provides for his students to work and support each other, and the

camaraderie he provides. To Dr. Gorga, I am thankful for all of the great discussions and

your honest advice, trust in me, and deep caring for students as individuals.

Thank you too to the others who have helped in my academic pursuits: the undergraduates

and high school students who graciously worked with me that the beginnings of their

promising careers, Dr. Richard Spontak, Dr. Kirill Efimenko, Dr. Medhi Afshari, Dr.

Eunkyoung Shim, and especially Dr. Jan Genzer whose generosity of time and resources

helped me through the roughest of patches.

I have enjoyed my time in the department more than I could have imagined. Thank you to

all of the support staff, with my utmost affection, including Mrs. Sandra Bailey, Saundra

Doby, Diane Harper, Clarice Whitmarsh and Sheila Hayes, for all the “little” things you

quietly do for us.

Thank you to my committee members: Dr. Pourdeyhimi, for your efforts to further my

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And, finally, to my friends and colleagues: Certainly, I cannot mention everyone who has

helped me quickly pass my time here, but here are a few who have made the biggest

impact: Michelle, for your eagerness to share your time and because you see the good in

everyone; Kristen, for your deep friendship, exposing me to new people, places and things,

and your professional guidance; Morgan, for understanding what grad school is like;

Christina Tang and Alina, for making the office such a pleasant place to be; Christina Devine

and all of the Joshs for being both fun and sincere; Jeremy, for keeping me in a constant

supply of coffee; Jairus, for being there, always, with a soap box for me to stand on; and to

Wade, for the wonderful person you are and for making that call five years ago today that

has made all the difference in my life.

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vi

TABLE

OF

LIST

OF

TABLES

CHAPTER 2 ... 16

TABLE 2.1 PHYSICAL PROPERTIES OF IPP RESINS. ... 41

CHAPTER 3 ... 61

TABLE 3.1 PHYSICAL PROPERTIES OF FIBER FORMING POLYMERS. ... 85

TABLE 3.2 POLYMER MORPHOLOGY AND SPINNING VELOCITY WITH CHANGING CORE/SHEATH POLYMER AND ASPIRATOR PRESSURE FOR 50 WT% CORE/50 WT% SHEATH FIBERS. “TRACE” INDICATES THAT THERE ARE ONLY TRACE AMOUNTS OF THE CRYSTALLINE MATERIAL PRESENT BY XRD. THE HIGHEST ASPIRATOR PRESSURE FOR EACH CONFIGURATION INDICATES THE MAXIMUM PRESSURE AT WHICH GOOD QUALITY FIBERS COULD BE OBTAINED. ... 87

CHAPTER 4 ... 104

TABLE 4.1 POLYMER PROPERTIES AND CONSTANTS. ... 130

TABLE 4.2 MEASURED AND PREDICTED VISCOSITY OF MULTILAYERED PP‐PLA MELTS. ... 135

TABLE 4.3 FEATURE SIZE OF SEBS COPOLYMER MORPHOLOGIES DEVELOPED IN MELT SPUN FIBERS. ... 144

CHAPTER 5 ... 149

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ix

LIST

OF

FIGURES

CHAPTER 1 ... 1

FIGURE 1.1 APPROXIMATE SHEAR RATES POLYMERS ARE SUBJECTED TO DURING COMMON COMMERCIAL

PROCESSING (ADAPTED FROM [24]). ... 9

FIGURE 1.2 MOLTEN POLMER IS METERED THROUGH THE SPIN PACK BY AN EXTRUSION PUMP (NOT SHOWN).

AS MOLTEN POLYMER EXITS THE SPIN PACK, IT IS QUENCHED BY COLD AIR, AND THEN DRAWN BY WARM AIR.

SUCTION LOCATED UNDERNEATH THE REVOLVING FORMING BELT AIDS IN LAYDOWN. AFTER LEAVING THE BELT,

FIBERS ARE FIRST COMPACTED AND THEN CALENDARED. WEBS ARE COLLECTED ON THE WINDER. MD DENOTES

MACHINE DIRECTION AND TD THE TRAVERSE DIRECTION. ... 10

FIGURE 1.3 SCHEMATIC SHOWING ESSENTIAL FEATURES THAT DIFFERENTIATE TWO NONWOVENS

MANUFACTURING PROCESSES, SPUN BONDING AND MELT BLOWING. SPUN BONDING INVOLVES QUENCHING THE

MOLTEN POLYMER WITH COOLED AIR FOLLOWED BY ATTENUATION (DRAWING) WITH WARMER AIR, WHEREAS

MELT BLOWING REQUIRES APPLYING HOT ATTENUATION AIR TO THE MOLTEN POLYMER NEAR THE EXIT OF THE

SPIN PACK ORIFICE. ... 11

FIGURE 1.4 COMMON CONSUMER END USES OF SYNTHETIC NONWOVENS [10]. ... 12

CHAPTER 2 ... 16

FIGURE 2.1 WAXD SPECTRA FOR ISOTATIC POLYPROPYLENE (IPP) FIBERS. SUBSRIPTS “S” AND “E” REFER TO

SOURCE OF POLYMER, S CORRESPONDING TO SUNOCO AND E TO EXXON. A. SPUNBOND IPPS, 2000 M∙MIN‐1;

B. MELT SPUN IPPS, 2000 M∙MIN‐1; C. MELT SPUN IPPS, 300 M∙MIN‐1; D. SPUNBOND IPPS, ≈ 20 M∙MIN‐1; E.

MELT SPUN IPPS, ≈ 20 M∙MIN‐1; F. MELT SPUN IPPE, 2000 M∙MIN‐1; G. MELT SPUN IPPE, 1000 M∙MIN‐1; H.

MELT SPUN IPPS, ≈ 20 M∙MIN‐1. SHADED AREAS INDICATE APPROXIMATE SPIN SPEEDS RESULTING IN

MESOMORPHIC FORMATION. ... 42

FIGURE 2.2 WAXD REFLECTIONS OF IPPS FIBER AS‐SPUN AND FOLLOWING ANNEALING. ... 43

FIGURE 2.3 MASS FRACTION OF EACH PHASE FOR IPPS FIBERS DETERMINED BY DSC. SUBSCRIPTS REFER TO

AMORPHOUS (A), MESOMORPHIC (M) AND CRYSTALLINE (C) PHASES. FIBERS WITH FILLED SYMBOLS WERE

PREPARED BY MELT SPINNING AND OPEN SYMBOLS BY SPUNBONDING. ... 44

FIGURE 2.4 BIREFRINGENCE VERSUS TAKE UP VELOCITY FOR IPPS FIBERS. FIBERS WITH FILLED SYMBOLS WERE

PREPARED BY MELT SPINNING AND OPEN SYMBOLS BY SPUNBONDING. ... 45

FIGURE 2.5 (A) TRUE STRESS VERSUS TRUE STRAIN FOR MELT SPUN (SOLID) AND SPUNBOND (DASHED) IPPS

SINGLE FILAMENTS AT AMBIENT CONDITIONS. CURVES ARE AVERAGES OF 8 TO 10 SAMPLES. (B) DATA FROM

ABOVE WITH CURVES SHIFTED HORIZONTALLY TO CREATE A MASTER CURVE. ... 46

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FIGURE 2.6 BIREFRINGENCE (ΔN) VERSUS TRUE STRAIN SHIFT (ΔεT, FROM FIGURE 4B) FOR IPPS FIBERS. ... 47

FIGURE 2.7 ORIENTATION FUNCTION VERSUS TAKE‐UP VELOCITY FOR THE A) CRYSTALLINE, B) AMORPHOUS

AND C) MESOMORPHIC PHASE CALCULATED FOR IPPS FIBERS. FIBERS WITH FILLED SYMBOLS WERE PREPARED BY

MELT SPINNING AND OPEN SYMBOLS BY SPUNBONDING. ... 48

FIGURE 2.8 A) BIREFRINGENCE, B) AMORPHOUS ORIENTATION FUNCTION, AND C) MESOMORPHIC

ORIENTATION FUNCTION VERSUS NETWORK DRAW RATIO FOR PP SPUN BOND (●) AND MELT‐SPUN () FIBERS.

PREDICTIONS BASED ON AFFINE (DASHED LINES, WITH VARYING FITTING PARAMETER N) AND PSEUDO‐AFFINE

MODELS (SOLID LINES). ... 49

FIGURE 2.9 TENACITY (TRIANGLES, LARGE DASH), EXTENSION TO BREAK (SQUARES, SMALL DASH), AND MASS

FRACTION MESOMORPHIC PHASE (CIRCLES, SOLID) VERSUS SPIN SPEED FOR MELT SPUN (FILLED SYMBOLS) AND

SPUNBOND (UNFILLED SYMBOLS) FIBERS. ERROR BARS INDICATE STANDARD ERROR AND TREND LINES ARE TO

GUIDE THE EYE. ... 50

FIGURE 2.10 WAXD SPECTRA FOR A) IPPS AND B) IPPE QUENCHED POLYMER. ... 51

FIGURE 2.11 VISCOSITY VERSUS SHEAR RATE FOR IPPS AND IPPE AT 185OC UNDER 10 L∙MIN‐1 N2. ... 52

FIGURE 2.12 DSC HEATING THERMOGRAMS OF QUENCH COOLED IPPE AND IPPS FIBERS WITH A HEATING RATE

OF 10O

C/MIN. ... 53

FIGURE 2.13 ENTHALPY OF THE MESOMORPHIC‐TO‐α PHASE TRANSITION IN IPP FIBERS AND QUENCHED

POLYMER FOR HEATING RATES OF 1 TO 20OC∙MIN‐1. ... 54

CHAPTER 3 ... 61

FIGURE 3.1 PP/PLA BICOMPONENT FIBERS CROSS SECTIONED AND IMAGED WITH FIB. A) PLACORE/PPSHEATH WITH PLA:PP MASS RATIO OF 50:50. PPCORE/PLASHEATH WITH PP:PLA MASS RATIOS OF B) 15:85, C) 50:50, AND D) 90:10. FIBERS WERE ALL COLLECTED AT AN ASPIRATOR PRESSURE OF 25 PSI. ... 86

FIGURE 3.2 FIBER DIAMETER VERSUS ASPIRATOR PRESSURE AND POLYMER CONFIGURATION FOR 50%

CORE/50% SHEATH FIBERS. ■ PP FILAMENTS, ● PLA FILAMENTS, PLACORE/PPSHEATH, V PPCORE/PLASHEATH. LINES ARE TO GUIDE THE EYE. ... 88

FIGURE 3.3 PERCENT CRYSTALLINITY WITH ANNEALING TIME FOR PP AND PLA ... 89

FIGURE 3.4 REPRESENTATIVE WAXD CURVES FOR PP AND PLA MORPHOLOGIES PRODUCED AT VARYING SPIN

SPEEDS. SPINNING AT 500 M/MIN RESULTS IN MESOMORPHIC PP, WHILE SPINNING AS LOW AS 20 M/MIN

RESULTS IN α‐MONOCLINIC CRYSTALLIZATION. PLA AT 4200 M/MIN RESULTS IN CRYSTALLIZATION, WHILE AT

56 M/MIN, NO SIGNIFICANT CRYSTALLINITY DEVELOPS. ... 90

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xi

FIGURE 3.5 TRUE STRESS VERSUS TRUE STRAIN FOR PP, PLA, AND BICOMPONENT PP/PLA FIBERS COLLECTED

AT 15 PSI ASPIRATOR PRESSURE FOR 50% CORE/50% SHEATH FIBERS. INSET: INITIAL TRUE STRESS – TRUE

STRAIN FOR OBTAINING MODULI FOR PP, PLA, AND BICOMPONENT PP/PLA FIBERS COLLECTED AT 0 PSI

ASPIRATOR PRESSURE. ... 91

FIGURE 3.6 BICOMPONENT FIBERS WITH 50 WT% CORE/50 WT% SHEATH SHOWING FRACTURE AFTER

MECHANICAL DRAWING. A) SEM IMAGE OF PLACORE/PPSHEATH FIBER DRAWN TO FAILURE. DASHED LINE INDICATES LOCATION WHERE A CROSS SECTION WAS TAKEN, WHICH IS SHOWN IN B. B) FIB IMAGE OF CROSS

SECTION NOTED IN A) TAKEN APPROXIMATELY 6 μM FROM FRACTURED FACE. C) SCHEMATIC OF THE STEPS

LEADING TO MECHANICAL FAILURE DURING TENSILE EXTENSION OF PLA CORE/PP SHEATH FIBER IN FOUR STEPS:

1. TENSILE EXTENSION OF BOTH COMPONENTS 2. PLA CORE FAILS WHILE DRAWING CONTINUES 3. PP

SUBJECTED TO FURTHER EXTENSION AND NARROWING IN THE ABSENCE OF PLA 4. FOLLOWING RUPTURE OF THE

PP, THE SHEATH RETRACTS D) FIB IMAGE OF PPCORE/PLASHEATH FIBER PARTIALLY DRAWN. E) CROSS SECTION OF FIBER FROM D). ... 92

FIGURE 3.7 TRUE STRAIN AT WHICH THE PLA SHEATH FAILS AS A FUNCTION OF THE ASPIRATOR PRESSURE USED

DURING FIBER COLLECTION AND FREQUENCY OF STRESS TRANSFER TO PPCORE AFTER PLASHEATH FAILURE FOR 50WT% CORE / 50WT% SHEATH PPCORE/PLASHEATH FIBERS. ... 93

FIGURE 3.8 SEM IMAGE OF PPCORE/PLASHEATH FIBER DRAWN TO FAILURE OF PLA SHEATH. ... 94

FIGURE 3.9 BREAKING TENACITY FOR PP, PLA, AND BICOMPONENT PP/PLA FIBERS WITH 50% CORE/50%

SHEATH COLLECTED AT A RANGE OF ASPIRATOR PRESSURES. ■ PP FILAMENTS, ● PLA FILAMENTS,

PLACORE/PP SHEATH, V PPCORE/PLASHEATH. LINES ARE TO GUIDE THE EYE. ... 95

FIGURE 3.10 A) BIREFRINGENCE FOR PP AND PLA SINGLE COMPONENT FIBERS. B) BIREFRINGENCE

CORRELATED TO TENSILE STRAIN SHIFT FOR PP AND PLA SINGLE COMPONENT FIBERS. □ PP FILAMENTS, ●

PLA FILAMENTS. ... 96

FIGURE 3.11 OVERLAYING BICOMPONENT FIBER TRUE STRESS‐TRUE STRAIN PROFILES WITH A) PLA AND B) PP

MASTER CURVES. (INSETS) STRAIN SHIFT OBTAINED FROM PLOT AS A FUNCTION OF SPINNING SPEED OF FIBERS. ... 97

FIGURE 3.12 THE AMORPHOUS MOLECULAR ORIENTATION FOR PLA FA. LINES ARE TO GUIDE THE EYE. ... 98

CHAPTER 4 ... 104

FIGURE 4.1 POLYMER MORPHOLOGY AND DIAMETER AS A FUNCTION OF ASPIRATOR PRESSURE AND SPINNING

CONFIGURATION. NOTE, ALL FIBERS CONTAINED SOME AMORPHOUS MATERIAL. ONLY CRYSTALLINE OR

PSEUDOCRYSTALLINE MORPHOLOGIES ARE NOTED IF PRESENT AT THE GIVEN SET OF SPINNING CONDITIONS.

MORPHOLOGIES ARE MARKED ACCORDING TO: = α‐CRYSTALLINE PP, = MESOMORPHIC PP, AND z =

CRYSTALLINE PLA. FIBERS SPUN AT THE SAME ASPIRATOR PRESSURE WITH DIFFERENT CONFIGURATION ARE

CONNECTED WITH TIE LINES. ... 129

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xii

FIGURE 4.2 DIAMETER OF FIBERS AS FUNCTION OF SPINNING PRESSURE, WITH AND WITHOUT COPOLYMER.

ERROR BARS INDICATE STANDARD ERROR FOR APPROXIMATELY 10 MEASUREMENTS. ... 131

FIGURE 4.3 SEM IMAGES OF MELT BLENDS AFTER ETCHING IN DCM. A) 50 WT% PP, 50 WT% PLA. B) 50

WT% PP, 47.5 WT% PLA, 2.5 WT% SEBS COPOLYMER. ... 132

FIGURE 4.4 SCHEMATIC OF LOADING STRATEGY FOR MULTILAYERED PP/PLA MELTS. A.) PP, B.) PLA ON

BOTTOM, PP ON TOP, C.) PP ON BOTTOM, PLA ON TOP, D.) 8 ALTERNATING LAYERS STARTING WITH PLA AND

ENDING WITH PP, E.) PLA. ... 133

FIGURE 4.5 ZERO‐SHEAR VISCOSITY FOR LAYERED PP‐PLA COMPOSITES WITH STANDARD ERRORS INDICATED.

DASHED LINE – S‐ROM PREDICTION. ERROR BARS ON S‐ROM INDICATE UNCERTAINTY OF PREDICTION BASED

ON STANDARD DEVIATION OF INDIVIDUAL PP AND PLA VISCOSITIES USED TO CALCULATE THE LAYERED

VISCOSITY. ... 134

FIGURE 4.6 BREAKING TENACITY OF PP/PLA FIBERS WITH AND WITHOUT COPOLYMER. A) PLACORE/PPSHEATH B) PPCORE/PLASHEATH. ERROR BARS INDICATE STANDARD ERROR OF APPROXIMATELY 10 TRIALS. ... 136

FIGURE 4.7 A) TRUE STRAIN AT WHICH THE SHEATH OF PPCORE/PLASHEATH FIBER RUPTURES. HOLLOW SYMBOLS INDICATE THAT THE PLA SHEATH RUPTURED AT APPROXIMATELY THE SAME STRAIN AS THE PP CORE. B)

FREQUENCY OF SHEATH RUPTURE PRIOR TO CORE RUPTURE DURING MECHANICAL STRAIN FOR PPCORE/PLASHEATH FIBERS. ... 137

FIGURE 4.8 DSC HEAT TRACES OF PPCORE/PLASHEATH FIBERS, WITH AND WITHOUT COPOLYMER ADDED TO THE PLA PHASE, IN THE VICINITY OF THE GLASS TRANSITION TEMPERATURE OF PLA AT TWO ASPIRATOR PRESSURES. . 138

FIGURE 4.9 DYNAMIC RHEOLOGY OF PLA AND PLA/COPOLYMER BLEND. ... 139

FIGURE 4.10 OPTICAL MICROGRAPHS OF FIBERS UNDER POLARIZED LIGHT FOR MEASURING OPTICAL

BIREFRINGENCE ACCORDING TO THE METHODS OF.40 THE REFRACTIVE INDEX (RI) NOTED ON EACH IMAGE

INDICATES THE REFRACTIVE INDEX OF THE LIQUID SURROUNDING THE FIBER. ... 140

FIGURE 4.11 STRAIN SHIFT OF PP/PLA FIBERS AS A FUNCTION OF SPINNING VELOCITY. ... 141

FIGURE 4.12 TEM IMAGES OF FIBER CROSS‐SECTIONS. A) (PLA+COPOLYMER) SPUN AS A SINGLE FIBER (NOT

CORE/SHEATH), B) PPCORE/(PLA+COPOLYMER)SHEATH, C) (PLA+COPOLYMER)CORE/PPSHEATH, D) PPCORE/PLASHEATH (NO COPOLYMER). ... 142

FIGURE 4.13 TEM IMAGES SHOWING MORPHOLOGIES OF SEBS COPOLYMER. COMPOUNDED WITH PLA AND

SPUN INTO FIBERS: CYLINDERS (A), CONCENTRIC CYLINDERS (B), SPHEROIDS CONCENTRIC TO CYLINDERS (C).

SEBS COPOLYMER AS RECEIVED (D). SEBS MELT BLENDED WITH PLA (E). SEBS CAST FROM DCM (F). ... 143

FIGURE 4.14 SCHEMATIC OF SEBS MORPHOLOGIES FORMED BY MELT COMPOUNDING WITH PLA AND

EXTRUDING INTO FIBERS. (A) BILAYERED TUBULES; (B) SWELLED, BILAYERED TUBULES; (C) CONCENTRIC,

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ENDBLOCKS, CONCENTRIC TO BILAYERED TUBULES; (F) SPHERES CONTAINING PE/B MIDBLOCKS WITH A

BILAYERED PS SHELL, CONCENTRIC TO BILAYERED TUBULES. RED = PS, BLUE = PE/B, GREEN = MIDBLOCK

SELECTIVE SWELLING POLYMER SUCH AS PLA OR COMPOUNDING MATERIAL. ... 145

CHAPTER 5 ... 149

FIGURE 5.1 MECHANISMS OF MESOMORPHIC TO α‐MONOCLINIC PHASE TRANSITION IN ISOTACTIC

POLYPROPYLENE. ADAPTED FROM [2]. BLACK TRIANGLES REPRESENT RIGHT HANDED HELICES AND WHITE

TRIANGLES REPRESENT LEFT HANDED HELICES. RED ARROWS INDICATE THE CHAINS INVOLVED WITH A

PARTICULAR MOLECULAR MOTION. ... 157

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1

CHAPTER

1 I

NTRODUCTION AND

O

VERVIEW

1.1 INTRODUCTION

Predicting properties of polymer fibers, such as mechanical, thermal, and others,

necessitates a solid understanding of processes associated with fiber formation from the

melt phase and the ability to accurately quantify characteristics, including the molecular

orientation of macromolecules therein, of the fibers. Spunbond fiber production has been

in existence nearly as long as melt spinning, yet the depth with which spunbond fiber

characterization has been studied pales in comparison to melt spinning. While this may in

part be due to the end uses of spunbond webs, it is more likely due to the many challenges

associated with characterizing fiber properties. Therefore, the first objective of this study is

to develop a method for quantifying molecular orientation in fibers using the mechanical

properties. During our investigation, we find it necessary to explore the way in which the

processing of a polymer affects the thermodynamic phase transitions in that polymer. A

comparison of the molecular orientation development in spunbonding relative to melt

spinning and the effect of spinning conditions on orientation is conducted as a means to

understanding the differences, if any, between these important fiber forming processes.

Another goal of this research was to develop a means of measuring the molecular

orientation of bicomponent fibers that did not necessitate regular cross sections, opaque

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two polymers on the orientation and molecular structure on the polymers could be

elucidated. Finally, we aim to determine the ability of a triblock copolymer to compatibilize

co‐spun polymers, where the block copolymer is an A‐B‐A type and the polymers to be

compatibilized are not A or B type.

1.2 BACKGROUND AND MOTIVATION

1.2.1 SPUNBOND SPINNING PROCESS

Melt spinning of polymer fibers started in the late 1930s.1‐3 Melt spinning involves melting

a polymer and extruding it through a die called a spinneret. As the polymer exits the die, it

is drawn through a quenching fluid (typically cooled air) under tension into a thin cylindrical

fiber. As the polymer solidifies, the roller (godet) that maintains the tension rotates at up to

several thousand m/min. Shear rates encountered during fiber extrusion can be

approximately 105 s‐1 (Figure 1.1). Fibrous filaments are then transferred onto additional

rollers that rotate at higher speeds to further draw the fiber in a continuous process.

Drawing reduces the fiber diameter and concurrently increases crystallinity and/or

molecular orientation by reorganizing the secondary molecular structures. After the fibers

are spun, a separate process, such as weaving, is necessary to create a fabric from the

polymer fibers. Several variations on melt spinning have been conceived and developed to

manufacture nonwoven fabrics from polymer fibers in one streamlined process. The first of

such processes was introduced in 1956 to produce nonwoven synthetic microfibers for use

in a device to collect and to monitor radioactive particles in the upper atmosphere.4,5

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efficiently produce nonwoven polymer fabrics.* Like traditional melt spinning, spunbonding

involves extruding a molten polymer through a die with micron‐sized holes (Figure 1.2). As

the molten polymer exits the die it too passes through a cooling medium. In spunbonding,

rather than drawing the fibers with a godet roll, warm jets of air are injected by aspirators

into the spunbond chamber to lengthen, narrow, and tangle the fibers as they are forced

onto a forming belt. The fibers are collectively sent through calendaring rolls, which bond

the fibers together and create a cohesive web. Spunbonding is considered the fastest

production method of manufacturing nonwovens, at a rate of about 100 m/min of fabric.

For comparison, knitting machines or looms can produce approximately 2‐8 m/min of

woven product.7 Figure 1.3 compares the spunbonding process with the similar melt

blowing process. Spunbonding generally produces fibers of 12 to 50 μm, while melt

blowing can produce fibers <1 μm.7 Spun bonding consists of five main areas of control:

filament extrusion, quenching, drawing/attenuation, laydown, and bonding. This study

focuses primarily on the first three of these steps, extrusion, quench, and draw.

1.2.2 NONWOVENS APPLICATIONS

Nonwovens products have the advantage of being simple and inexpensive to manufacture.

However, because nonwovens lack a post‐spinning draw step as in melt spinning, they lack

the high degree of order, and hence strength, of traditional melt‐spun fibers. The first

commercial uses of modern nonwovens were predominantly in the automotive industry,8

*

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where they replaced more expensive, woven filler materials.† Though nonwovens were

used extensively in the decades following their inception, early reports of melt blown or

spun bond nonwovens in the literature were few compared with traditional melt spinning.9

Because nonwoven fibers were generally used as low cost materials, research concentrated

on either the mere ability to produce a nonwoven material from a given thermoplastic or

the mechanical properties of the resulting nonwoven webs, not the individual filaments.

While the bulk of nonwovens material is still manufactured for disposable or low

value‐added products, nonwovens have found use in a number of high value‐added

applications (cf. Figure 1.4).10 The majority of fabrics used in filtration filter media, for both

liquid and dry applications, are already based on nonwovens, which represent about $2

billion annually.11 Nonwovens’ application in filter devices is based on both size exclusion

(pore size of web) and interaction of the filtered particles with the surface of the fibers. For

example, melt blown polypropylene (PP) nonwovens have been used to adsorb oils and fats

from industrial waste water emulsions12 and to size‐exclude suspended solids while

adsorbing organics13 from reclaimed waste water. Separation of white blood cells and/or

leukocytes from blood has been accomplished by surface‐modified polyester

nonwovens.14,15 A significant fraction of the commercial nonwovens market consists of

nonwovens used for their electrical properties.5 Further, carbonized nonwovens have been

investigated for use as the electrode material in electrochemical capacitors because they

†

Nonwoven fabrics may be the oldest type of fabric, dating back thousands of years to felted material produced

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5

have good electrical and thermal conductivity and are able to accumulate electrical charge,

compared to other active carbon forms.16 Fibrous active carbons are also partially ordered,

which facilitates electron conductance along the fiber; they are easy to fix spatially in the

product, providing good contact compared to grainy forms of active carbon.16 Another

promising area involving fibrous materials involves tissue scaffolding: nonwovens have

been studied extensively as matrices for tissue scaffolding because they have the potential

to provide excellent control over fiber diameter and modulus, which have been shown to

greatly affect cell proliferation and differentiation.17,18

Nonwovens possess extraordinary potential for many new commercial markets

owing to their ease of production and low cost.9 Current research in nonwoven

technologies requires focusing on the fundamental mechanisms by which polymer fibers

are formed in order to comprehend and control the relationships between processing

conditions and fiber morphology and to aid in the prediction of macroscopic mechanical or

other properties without requiring the production of numerous samples. Relating individual

fiber properties to macroscopic properties of nonwovens webs presents another set of

challenges. Though the formation of synthetic fibers has been studied for more than 70

years, progress in understanding the fundamental mechanisms of fiber formation,

especially in non‐traditional melt spinning processes, that are necessary to effectively

control fiber structure and properties has been slow and techniques to analyze and

(22)

6

more detailed molecular insight that links the macroscopic properties of fibers with

microscopic organization of polymers inside fibers.

1.2.3 SELECTED FIBER STRUCTURE‐PROPERTY RELATIONSHIPS

Fiber strength and technical properties are related to the fiber microstructure. The

relationship between the optical properties of a polymer and the orientation of the

molecules therein is a well observed phenomenon. Macroscopic stretching of rubber was

shown to result in a measurable change in the birefringence of the material; by reorienting

the molecules, crystallization took place, and both the orientation and crystals contributed

to the increase in observed birefringence by stretching.19 Ward (1962) [20] later derived

expressions to relate the optical birefringence to the molecular orientation of polyethylene

and introduced the idea that the relationship between mechanical anisotropy in polymer

materials and molecular orientation was more complex than that of birefringence to

molecular orientation. Subsequent work in these areas has focused on describing the

theoretical bounds of polymer strength and methods of characterizing the molecular

orientation, as reviewed by Young and Eichhorn (2007).21

Extensive work has been carried out towards understanding structure‐property

relationships for commodity polymers, such as PP and poly(ethylene terephthalate) (PET).

Mechanical testing of semi‐crystalline fibers and films reveals that that high non‐crystalline

(amorphous) orientation reflects the amount of taut tie molecules present.22 Therefore the

non‐crystalline orientation in fibers required for high tenacity23 is proportional to fiber

(23)

7

depends on the interplay between crystallinity, and the orientation of crystalline and

amorphous regions.23 Spunbond PP fiber‐to‐fiber bonding, tensile strength, and breaking

elongation can be improved by altering the molecular structures.24,25 Spunbond PET fibers

and polyethylene (PE) films exhibited increasing tension at break and/or tensile strength

and with increasing crystallinity and crystalline orientation.22,26 PP webs were shown to

have better tensile properties with increasing overall orientation and decreasing fiber

diameter.27

1.3 RESEARCH GOALS AND ORGANIZATION OF DISSERTATION

The theme of this work is exploring the function of a high‐shear flow field (i.e., fiber

extrusion) in developing morphologies that are unique relative to those formed in the

absence of such a flow field. One of our objectives is to understand how the presence of

the high‐shear flow field of fiber extrusion results in the formation of certain polymer

morphologies as well as how the thermal stability of these phases are affected by the

processing history of the polymer. Another primary objective is to understand and develop

methods to quantify the molecular orientation in fibers that is a result of the high‐shear

flow fields encountered with fiber extrusion. Finally, we wish to investigate the

compatibility between polymers in core/sheath fibers and the effect a copolymer additive

has on improving the mechanical properties of core/sheath fibers. To these ends, Chapter 2

investigates the mesomorphic phase of PP created during spunbond fiber spinning and

demonstrates that this phase is thermally more stable than the mesomorphic phase

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8

molecular orientation of single component fibers from the stress‐strain curves of fibers is

established. Chapter 3 contains a study of the effect of co‐spinning PP with a second

polymer, poly(lactic acid) (PLA), in core/sheath configuration. Co‐spinning has the potential

to develop greater molecular orientation in one of the polymers than can be achieved

individually due to the presence of the polymer‐polymer interface, yet the methods of

measuring this molecular orientation can prove difficult or impossible for some polymer

combinations. To answer the question of how molecular orientation is developed in

bicomponent fibers, the mechanical property – molecular orientation correlation from

Chapter 2 is extended to bicomponent fibers. Finally, in Chapter 4, we address the concept

of “compatibilization,” or improving the adhesion between polymers, in core/sheath fibers

by adding a triblock copolymer, poly(styrene‐b‐ethylene‐co‐butylene‐b‐styrene). This work

reveals the effect of the high‐shear extrusion process on the block copolymer morphology

resulting from melt compounding an “A‐B‐A” type of block copolymer with a “C” type

(25)

9

Figure 1.1 Approximate shear rates polymers are subjected to during common commercial

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10

Figure 1.2 Molten polmer is metered through the spin pack by an extrusion pump (not

shown). As molten polymer exits the spin pack, it is quenched by cold air, and then drawn

by warm air. Suction located underneath the revolving forming belt aids in laydown. After

leaving the belt, fibers are first compacted and then calendared. Webs are collected on the

winder. MD denotes machine direction and TD the traverse direction.

molten polymer

spin pack

forming belt

quench (cold air)

aspirator jets (warm air)

TD

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11

Figure 1.3 Schematic showing essential features that differentiate two nonwovens

manufacturing processes, spun bonding and melt blowing. Spun bonding involves

quenching the molten polymer with cooled air followed by attenuation (drawing) with

warmer air, whereas melt blowing requires applying hot attenuation air to the molten

polymer near the exit of the spin pack orifice.

forming belt cooled

quench air

warm

attenuation air

hot

attenuation air

spin pack

molten polymer

(28)

12

(29)

13 REFERENCES

[1] Dreyfus, H. Patent GB346689 (1931), “Improvements in the production of artificial

filaments, threads, ribbons and the like and in products resulting therefrom.”

[2] Triggs, W. Patent GB495790 (1938), “Improvements in the manufacture of

polyamides.”

[3] Greenewalt, C. US patent no. 2217743 (1940) “Melt spinning apparatus suitable for

the manufacture of vinyl polymer fibers, etc.”

[4] Wente, V. Industrial and Engineering Chemistry 1956, 48, 1342.

[5] McCullough, J. International Journal of Nonwovens 1999, 8, 139.

[6] Li, D.; Xia, Y. Advanced Materials 2004, 16(14), 1151‐1170.

[7] Genis, A.; Usov, V.; Razboeva, V. Fibre Chemistry 2007, 39 (1), 1‐6.

[8] Shailer L. Industrial and Engineering Chemistry 1959, 51, 901.

[9] Choi, K.; Spruiell, J.; Fellers, J.; Wadesworth, L. Polymer Engineering and Science

1988, 28(2), 81‐89.

[10] World Synthetic Fibers 1995 – 2005, Vol 2, Sect. VI, pp. 532‐561.

[11] Palexpo, G. Filtration & Separation 2005, 42(2), 26‐29.

[12] Kalużka, J.; Lebiedowski, M. Fibres and Textiles in Eastern Europe 2003, 11 (1), 64‐

(30)

14

[13] Lui, L.; Xu, Z.; Song, C.; Gu, Q.; Sang, Y.; Lu, G.; Hu, H.; Li, F. Desalination 2006,

201(1‐3), 198‐206.

[14] Callaerts, A.; Gielis, M.; Sprengers, E.; Muylle, L. Transfusion 1993, 33(2), 134‐138.

[15] Natori, S.; Gomei, Y.; Higuchi, A. Journal of biomedical materials research. Part B,

Applied biomaterials 2006, 78B(2), 318‐326.

[16] Cisło, R.; Krucińska, I.; Babeł, K.; Koszewska, M. Fiber & Textiles in Eastern Europe

2004, 12(3), 70‐74.

[17] Yasuda, K.; Inoue, S.; Tabata, Y. Tissue Engineering 2004, 10(9‐10), 1587‐1596.

[18] Engler, A.; Sen, S.; Sweeney, L.; Disher, D. Cell 2006, 126(4), 677‐689.

[19] Treloar, L. Transactions of the Faraday Society 1941, 37, 84‐97.

[20] Ward, I. Proceedings of the Physical Society 1962, 80(5), 1176.

[21] Young, R.; Eichhorn, S. Polymer 2007, 48, 2‐18.

[19] Branciforti, M.; Pimentel, R.; Bernardi, A.; Bretas, R. Journal of Applied Polymer

Science 2006, 101(5), 3161‐3167.

[20] Samuels, R. Structured Polymer Properties. New York: John Wiley & Sons, 1974.

[21] Bhat, G.; Nanjundappa, R.; Kotra, R. Thermochemica Acta 2002, 392, 323‐328.

[22] Nanjundappa, R.; Bhat, G. Journal of Applied Polymer Science 2005, 98(6), 2355‐

(31)

15

[22] Wang, H.; Jin, X.; Wu, H.; Yin, B. Journal of Material Science 2007, 42(19), 8006‐

8009.

[23] Zang, D.; Bhat, G.; Sanjiv, M.; Wadsworth, L. Textile Research Journal 1998, 68(), 27‐

35.

[24] Painter, P.; Coleman, M. Fundamentals of Polymer Science. Washington, D.C.: CRC

press, 1997.

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16

CHAPTER

2 M

ESOMORPHIC

‐

α

‐

M

ONOCLINIC

P

HASE

T

RANSITION IN

I

SOTACTIC

P

OLYPROPYLENE

:

A

S

TUDY OF

P

ROCESSING

E

FFECTS ON

S

TRUCTURE AND

M

ECHANICAL

P

ROPERTIES

Chapter 2 is essentially a manuscript by Sara Arvidson, Saad Khan, and Russell Gorga published in

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17

M

ESOMORPHIC

‐

α

‐

M

ONOCLINIC

P

HASE

T

RANSITION IN

I

SOTACTIC

P

OLYPROPYLENE

:

A

S

TUDY OF

P

ROCESSING

E

FFECTS ON

S

TRUCTURE AND

M

ECHANICAL

P

ROPERTIES

Sara Arvidson, Saad Khan, and Russell Gorga

ABSTRACT

We report the enthalpy for the mesomorphic to α‐monoclinic phase transition in

polypropylene under varying thermal treatments. The mesomorphic phase is created by

fiber spinning and rapid quenching methods and identified using wide‐angle X‐ray

diffraction and differential scanning calorimetry. Fiber mesomorphs are found to have a 3‐

fold increase in enthalpy of transition per gram of mesophase compared with our

measurements of quenched polypropylene and previous reports of quenched

polypropylene. In addition, systematic tensile testing over a range of spin speeds and

polymer morphologies reveals that the presence of mesomorphic regions does not

correlate with reduced fiber strength as has been previously suggested. Fiber true stress‐

true strain curves obtained at varying take‐up velocities are compared to determine the

(34)

18

We find that the tensile strain shift correlates with birefringence, thereby providing an

alternative method to assess molecular orientation in fibers, an important factor for fiber

strength. This approach can prove useful for fibers in which measuring the molecular

(35)

19 2.1 INTRODUCTION

Much work has been done to produce and model high‐speed melt spinning for the

production of high tenacity filaments for a number of polymers.1‐5 Such processes generally

include post‐spinning drawing steps that result in greater molecular orientation in the fiber.

Though similar in many ways to melt spinning, nonwovens processes that aim to create high

strength webs from fibers in a single continuous process such as spunbonding or melt

blowing may present extra challenges due to their coupled spin‐draw step. While both melt

spinning and nonwovens processes were commercialized by the 1950s,6 a much greater

emphasis on modeling in the literature has been applied to generalized melt‐spun fiber

formation and use.

Due to relative ease of spinning, low cost, and rapid crystallization, isotactic

polypropylene (iPP) is one of the most commonly spun polymers in nonwovens processes.

iPP crystallization is a complex process that involves several crystalline morphologies and

competing crystallization mechanisms that depend greatly on temperature and stress. iPP

has been shown to crystallize predominately in the α‐monoclinic form under isothermal

crystallization,7 slow cooling,8 and during melt spinning of filaments.9 The β–hexagonal or

γ–triclinic polymorphs may be obtained while crystallizing from the melt at high

undercoolings or pressures, or with the addition of nucleating agents.10‐14 The “meso”

polymorph, often described as a smectic or paracrystalline phase, has been identified in iPP

(36)

20

and 40oC;17 and in fiber spinning at moderate take‐up velocities, high extrusion

temperatures, low molecular weight distributions, and low draw‐down ratios.9,10,18‐20 The

mesomorph is not an imperfect crystal, but rather has molecular ordering between that of

the amorphous phase and a true crystalline phase; 21 mesomorphic iPP has a high degree of

order in the direction of the chain axis but little in its lateral packing.22 Mesomorphs are not

unique to iPP, having been shown for a wide variety of polymers including polyesters and

polysiloxanes23 and are important for improving polymer clarity24 and possibly polymer

processability.22

Both the α and meso iPP polymorphs are composed of 31 helices. While the

arrangement of the left‐ and right‐handed helices of the mesomorphic form are disordered,

left and right helices follow a well‐defined sequence of handedness in the α form.25,26 Upon

heating, the meso phase transforms into the α phase, possibly by thickening of existing α

crystals and/or by structural rearrangements in the mesomorphic phase.27 The mechanism

of this structural rearrangement has yet to be determined in any conclusive way.16,26,28‐31

Androsch26 observed the meso‐to‐α phase transition through atomic force microscopy of

nano‐scale domains and determined that initial mesomorphic domains were not destroyed

during the phase transformation but could not rule out local melting within domains.

Consequently, it is unclear if the meso‐to‐α phase transition occurs via melting into the

(37)

21

meso‐to‐α transition have been discussed, one involving chains unwinding to reverse

handedness and the other requiring chain translocation without changing handedness.26,31

Extensive work has been published describing the effect of spinning conditions on the

development of orientation in the α, β, and amorphous phases in iPP fibers, yet orientation

development of the mesophase (and the mesophase in general) has largely been ignored in

fibers.19 While orientation of the amorphous phase has been linked to fiber strength,32 with

crystallites acting as network junctions,33 the role of mesophases on mechanical properties

is ambiguous. It has been suggested that mesophase content poses significant detriment to

the strength of iPP fibers34 and increases ductility in quenched films,35 yet polymer

mesophases are credited with improving mechanical properties and processibility over

crystalline phases.22 Further gains in defining structure‐property relationships for

semicrystalline polymers have not been made due to an incomplete understanding of

viscoelastic and viscoplastic responses of the various morphologies present in the polymer,

during the melt solidification stage or during mechanicals tests at ambient temperature.36

An overwhelming majority of studies on the thermal behavior of iPP surrounding the meso‐

to‐α transition focuses on mesomorphic iPP formed by rapid quenching rather than fiber

spinning. Estimated enthalpies of the meso‐to‐α phase transition range from 8.8 to 16.7 J/g

for quench‐formed meso iPP due to varying polymer thermal histories, experimental

methodologies and polymer grades,28,37‐39 which makes accurately determining the meso

(38)

22

Additionally, it has been shown that the enthalpy of the meso‐to‐α phase transition may

differ for fibers and quench‐formed meso iPP.34

In this study we compare the enthalpy of the iPP meso‐to‐α phase transition of

fibers to that of quenched polymer and discuss the implications of the enthalpy in terms of

supporting existing theories of the mechanism of this phase transition. Further, we report

the mechanical properties and orientation functions of meso‐containing iPP fibers in the

context of melt spinning and spunbonding routes of iPP fiber formation and assess the

dependence of the mesomorphic phase orientation on take‐up velocity with a three‐phase

model. We begin with a discussion of fibers containing mesophase iPP and follow with that

of quenched‐formed iPP, ultimately comparing the thermal properties of each.

2.2 EXPERIMENTAL

2.2.1 MATERIALS AND SAMPLE PROCESSING.

Commercially available iPP formulations with similar molecular weight profiles were

selected from two manufactures: iPPS (Sunoco Chemicals Polymers Division, product

CP360H) and iPPE (ExxonMobil Chemical, product PP3155) (Table 2.1). Melt‐spun iPPS and

iPPE fibers were produced at the Hills spinning line in the College of Textiles Fiber Science

Lab at NC State University at take‐up velocities up to 2000 m∙min‐1. No secondary drawing

was applied to melt‐spun fibers. Spunbond iPPS fibers were produced at the Nonwovens

Cooperative Research Center (NCRC) Partners’ Pilot facilities located at NC State over a

range of aspirator pressures. Fibers were collected following extrusion but before the

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23

mass throughputs of 0.4 grams∙hole‐1min‐1. Fibers were also collected for each process

without the use of the godet or aspirator pressure with the intent of creating unoriented,

isotropic fibers at a low take‐up velocity. This velocity (≈20 m∙min‐1) is referred to as the

“freefall” velocity. The densities of semicrystalline fibers were estimated from the densities

for 100% crystalline, amorphous, and mesomorphic iPP: ρC=0.936, ρA=0.858, and ρM=0.916

g∙cm‐3.40,41 Due to the nature of spunbonding, individual fiber take‐up velocities are

variable and unknown at the time of formation. The equivalent take‐up velocities for

spunbond fibers are calculated from the continuity equation (Equation 1) to facilitate

comparison between spunbond fibers and melt‐spun fibers, for which take‐up velocity V

(m/min) is directly controlled.

c A Q V

⋅ ρ

= (1)

In Equation 1, Q is the throughput (g∙hole‐1∙min‐1), ρ is the polymer density (g∙m‐3) and Ac is

the average cross‐sectional area of an individual fiber (m2).

Quenched iPPS and iPPE polymers were prepared by heating the respective resins in

a furnace to 225oC until molten (~5 min), pressing the molten polymer between sheets of

aluminum, and holding in a furnace at 225oC for 15 minutes. Samples were removed from

the furnace directly into one of several quenched baths used: ice‐water (0oC), acetone‐dry

ice (‐78oC), pentane‐liquid N2 (‐131oC), and liquid N2 (‐196oC).39,42 Quenched polymers were

(40)

24 2.2.2 DSC MEASUREMENTS

The crystalline (xC), mesomorphic (xM), and amorphous (xA) fractional content was

determined using a TA Instruments Q2000 model differential scanning calorimeter. Scans

were carried out on 6 to 12 mg samples in standard aluminum pans calibrated to an indium

standard. Heating rates, unless otherwise specified, were 10oC∙min‐1 under 50mL∙min‐1 N2

purge. To identify the fraction of polymer in the crystalline phase, the heat of fusion for a

sample ΔHm was compared to the heat of fusion of a gram of 100% crystalline

polypropylene ΔHmo according to Equation 2.

_o

m m C

∆H

x = (2)

ΔHmo = 207 J∙g‐1 was used for 100% crystalline iPP.43

2.2.3 WIDE ANGLE X‐RAY DIFFRACTION (WAXD) MEASUREMENTS

X‐ray diffraction studies were conducted with a CuKα radiation source (λ=1.542 Å) at 40kV x 30 mA for 1800s on a Bruker D‐5000 diffractometer equipped with a Highstar area detector.

Diffraction patterns were analyzed with Bruker General Area Detector Diffraction System

(GADDS) software. Fibers were aligned and placed in the apparatus oriented vertically, with

the plane of the fibers perpendicular to the x‐ray beam. In transmission mode, the intensity

was recorded for 2θ in the range of 10o to 32o. Scans taken with an empty sample holder

served as a background for subtraction. Orientation of the crystalline regions with respect

(41)

25

rotational symmetry about the fiber axis, as shown in Equation 3 and by employing the

Hermans‐Stein orientation factor, Equation 4.44‐46

∫

π π φ φ φ φ φ φ φ = φ 2 2

0 hkl ,jz j,z j,z

0 j,z j,z j,z

2 z ,j hkl z ,j 2 d )sin ( I d )sin ( )cos ( I

cos (3)

2 1 -) ( cos 3

f ,jz

2

j

φ

= (4)

Here, Ihkl is the intensity defracted from the (hkl) planes which are normal to the j‐

crystallographic axis. In α‐monoclinic iPP, the helices lay along the c‐crystallographic axis,

which lacks a convenient reflection, so the intensities of the <110> and <040> bands were

integrated as a function of the azimuthal angle. The average cosine squared of the angle

between the fiber axis and the c‐axis can be calculated according to Equation 5.

cosφ2c,z =1−1.1099cos2(φ_110,_z)-0.901cosφ2040,z (5)

Mesomorphic iPP was modeled as a hexagonal crystal.47 Due to symmetries within a

hexagonal crystal,48 only one reflection is required to characterize the mesomorphic

orientation, which was estimated according to Equation 6.

cosφ2c,z =1−2cos2(φ_0002,_z) (6)

2.2.4 FIBER MECHANICAL TESTING

Mechanical testing of fibers was conducted on an Instron model 5544 at ambient conditions

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26

5 N load cell. Single filaments with a gage length of 28.6 mm were drawn at a crosshead

speed of 25.4 mm/min until breakage occurred.

2.2.5 MICROSCOPY

Fiber diameters measured by optical microscopy were used to calculate fiber cross‐sectional

areas. The refractive indices of polypropylene fibers were measured by a Mach‐Zehnder

type interference microscope by Aus Jena with polarized green light (λ=546 nm). Fiber

birefringence (Δn) was calculated from refractive indices (n) in the directions parallel and

normal to the fiber axis.

B D λ Z n

n _or _o

⋅ ⋅ + =

⊥ (7)

∆n=n -n⊥ (8)

Here, Z is the width of fringes, and B is the fringe shift, and no is the refractive index of the

immersion oil.

Birefringence relates to the molecular orientations of each phase according to

M M oM

o A A A o C C

Cf x f x f

x

n= Δ + Δ + Δ

Δ (9)

where we have included the mesomorphic phase contributions to the total birefringence.

Intrinsic birefringences (Δo) used were 0.331, 0.040, and 0.0468 for the crystalline,

mesomorphic, and amorphous iPP phases, respectively.32,41,49,50 Calculation of the

orientation functions of the crystalline and mesomorphic phases (fC and fM) are described by

(43)

27

Equation 9 ignores the form birefringence which takes into account the crystallite shape

and generally only contributes 5‐10% of the overall birefringence for most polymers.

However, polypropylene has been reported to have negligible form birefringence.32

2.3 RESULTS AND DISCUSSION

It is well known that resin material properties (isotacticity, molecular weight, viscosity, etc.)

and polymer processing (time duration at a given temperature and/or rate of temperature

change) determine the crystallinity, crystal form, molecular orientation, and ultimately fiber

strength. In this study, we investigate the effect of spinning conditions by comparing

spunbonding to melt spinning for iPPS resin and we investigate material properties by

comparing melt spinning and quenching of iPPS and iPPE resins.

2.3.1 FIBER EXTRUSION: MORPHOLOGY AND MOLECULAR ORIENTATION

Figure 2.1 shows WAXD spectra for iPPS spunbond, iPPS melt‐spun, and iPPE melt‐spun fibers

for a range of take‐up velocities. WAXD patterns exhibiting strong peaks located at 2θ=

14.2, 17.1, and 18.7 degrees correspond to the (110), (040), and (130) α‐monoclinic

reflections, respectively,51,52 while broad reflections at 2θ= 15 and 21 correspond to

mesomorphic iPP.53 Melt‐spun and spunbond fibers exhibit α reflections at both low and

high take‐up velocities. Due to the ease of controlling the take‐up velocity on a melt

spinning line, fibers were obtained at very low to moderate velocities (20 < V < 1600

m∙min‐1) that were not always possible with the spunbond process. However, where similar

take‐up velocities were obtained by both processes, melt‐spun and spunbond fibers

Molecular Orientation and Fine Structure Development in Polymer Fiber Formation.

ACKNOWLEDGMENTS

TABLE

OF

CONTENTS

LIST

OF

TABLES

LIST

OF

FIGURES

CHAPTER

1

I

O

CHAPTER

2

M

‐

α

‐

M

P

T

I

P

:

A

S

P

E

S

M

P

M

‐

α

‐

M

P

T

I

P

:

A

S

P

E

S

M

P

∫

∫