Handbook of Industrial
Polyethylene and Technology
Scrivener Publishing
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Handbook of Industrial
Polyethylene and Technology
Definitive Guide to Manufacturing,
Properties, Processing,
Applications and Markets
Edited by
Mark A. Spalding and
Ananda M. Chatterjee
This edition first published 2017 by John Wiley & Sons, Inc., 111 River Street, Hoboken, NJ 07030, USA and Scrivener
Publishing LLC, 100 Cummings Center, Suite 541J, Beverly, MA 01915, USA
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Library of Congress Cataloging-in-Publication Data
Names: Spalding, Mark A., editor. | Chatterjee, Ananda M., 1946-2016, editor.
Title: Handbook of industrial polyethylene and technology : definitive guide
to manufacturing, properties, processing, applications and markets /
edited by Mark A. Spalding and Ananda M. Chatterjee.
Description: Hoboken, NJ : John Wiley & Sons, 2017. | Includes
bibliographical references and index. |
Identifiers: LCCN 2017030446 (print) | LCCN 2017037911 (ebook) | ISBN
9781119159773 (epub) | ISBN 9781119159780 (pdf) | ISBN 9781119159766
(cloth : alk. paper)
Subjects: LCSH: Polyethylene. | Polymers--Industrial applications.
Classification: LCC TP1180.P65 (ebook) | LCC TP1180.P65 H36 2017 (print) |
DDC 668.4/234--dc23
LC record available at https://lccn.loc.gov/2017030446
Cover image: The Editor
Cover design by Russell Richardson
Set in size of 11pt and Minion Pro by Exeter Premedia Services Private Ltd., Chennai, India
Printed in the USA
10 9 8 7 6 5 4 3 2 1
Dr. Ananda M. Chatterjee started his career at Celanese Plastics (1978) and worked
for Shell Development, Union Carbide, Dow Chemical, Kaneka Texas, and Ingenia
Polymers. He had extensive and in-depth experience in polymer research and
development, product commercialization, polymer film processing and products
technology, polymer additives technology, additive blends technology and processing,
customer technical service, intellectual property (patent) invention and licensing and
technology transfer of UNIPOL polypropylene technology worldwide. He was a Fellow
and Honored Service Member of the Society of Plastics Engineers (SPE). Dr. Chatterjee
held a Ph.D. from the University of Massachusetts Amherst and held over 25 U.S.
patents. Dr. Chatterjee passed away in 2016 during the development of this book.
Contents
Foreword by Barry Morris
xi
Preface
xiii
List of Contributors
xv
Part 1: Principles and Properties of Polyethylene
1 An Industrial Chronology of Polyethylene
Christopher Dobbin
3
2 Catalysts for the Manufacture of Polyethylene
Yury V. Kissin
25
3 Ethylene Polymerization Processes and Manufacture of Polyethylene
Ian D. Burdett and Ronald S. Eisinger
61
4 Types and Basics of Polyethylene
Rajen M. Patel
105
5 Molecular Structural Characterization of Polyethylene
A. Willem deGroot, David Gillespie, Rongjuan Cong, Zhe Zhou and Rajesh
Paradkar
139
6 Thermal Analysis of Polyethylene
Kevin Menard and Noah Menard
217
7 Rheology of Polyethylene
Gregory W. Kamykowski
239
8 Processing-Structure-Property Relationships in Polyethylene
Rajen M. Patel
283
9 Mechanical Properties of Polyethylene: Deformation and Fracture Behavior
Alexander Chudnovsky, and Kalyan Sehanobish
309
Part 2: Processing and Fabrication of Polyethylene
10 Single-Screw Extrusion of Polyethylene Resins
Mark A. Spalding
339
11 Twin-Screw Extrusion of Polyethylene
Yoshitaka Kumura, Amit K. Chaudhary and Mark A. Spalding
357
vii
viii
Contents
12 Blown Film Processing
Thomas I. Butler
381
13 Cast Film Extrusion of Polyethylene
Hyunwoo Kim, Mark A. Spalding, Kurt A. Koppi, Wes Hobson
and Joseph Dooley
411
14 Extrusion Coating and Laminating
Thomas Bezigian
429
15 Injection Molding
Jon Ratzlaff and Thomas Giovannetti
443
16 Blow Molding of Polyethylene
Mohammad Usman and Abdul Sami Siddiqui
475
17 Rotational Molding
Jon Ratzlaff and Glenn E. Larkin, Jr.
535
18 Thermoforming Polyethylene
Roger C. Kipp
573
19 Polyethylene Pipe Extrusion
Pam Maeger, V. Rohatgi, D. Hukill, N. Koganti and B. Martinez
591
20 Polyethylene Foam Extrusion
N. S. Ramesh
603
21 Expanded Polyethylene Bead Foam Technology
Steven R. Sopher
637
22 Polyethylene Fiber Extrusion
Johannes Fink
669
23 Polyethylene Compounding Technologies
Charles D. Park II and Steven Blazey
695
24 Polyethylene Modification by Reactive Extrusion
Adriana I. Moncada, Wenyi Huang and Nicholas Horstman
715
Part 3: Additives for Polyethylene
25 Degradation and Stabilization of Polyethylene
Joseph J. Fay and Roswell E. King, III
753
26 Light Stabilization of Polyethylene
Feng Zuo and Tad Finnegan
771
27 Acid Scavengers for Polyethylene
Robert L. Sherman Jr. and Kimberly E. Kern
793
28 Slip Agents
John Gray and Thomas Breuer
821
Contents ix
29 Antiblocking Additives
Johannes Fink
833
30 Antistatic Additives for Polyethylene
Gina Butuc, Gea Spijkerman, Sue te Hofstee and Ted Kampen
853
31 Antifogging Agents for Polyethylene Films
Michele Potenza and Bjarne Nielsen
865
32 Lubricants for Polyethylene
Johannes Fink
877
33 Fluorinated Polymer Processing Aids for Polyethylene
David A. Seiler, Francois Beaume, Samuel Devisme and Jason A. Pomante
889
34 Chemical Blowing Agents for Polyethylene
Peter Schroeck, Randy Minton, Theresa Healy and Larry Keefe
909
35 Flame Retardants for Polyethylene
Rudolf Pfaendner
921
36 Nucleating Agents for Polyethylene
Darin L. Dotson
935
37 Antimicrobial Agents for Polyethylene
Ivan Ong
967
38 Pigments and Colorants for Polyethylene
Roger Reinicker
985
Part 4: Applications of Polyethylene
39 Fillers and Reinforcing Agents for Polyethylene
János Móczó and Béla Pukánszky
1035
40 Flexible Packaging Applications of Polyethylene
Jeff Wooster and Jill Martin
1071
41 Rigid Packaging Applications
Cliff R. Mure
1091
42 Pipe and Tubing Applications of Polyethylene
Bryan E. Hauger
1109
43 Wire and Cable Applications of Polyethylene
Scott H. Wasserman, Bharat I. Chaudhary, Jeffrey M. Cogen,
Mohamed Esseghir and Timothy J. Person
1125
44 Medical Applications of Polyethylene
Benjamin Poon and Len Czuba
1155
45 Automotive Applications for Polyethylene
Kalyan Sehanobish
1169
x
Contents
46 Textile, Hygiene, Health, and Geosynthetic Applications of Polyethylene
Sanjiv R. Malkan
1179
47 Applications of Polyethylene Elastomers and Plastomers
Kim L. Walton, Tim Clayfield, Jim Hemphill and Lisa Madenjian
1197
Part 5: The Business of Polyethylene
48 Product Regulatory Considerations for Polyethylene
Tor H. Palmgren
1221
49 Sustainability and Recycling of Polyethylene
Thomas Nosker and Jennifer Lynch
1245
50 Bio-Polyethylene and Polyethylene-Biopolymer Blends
Johannes Fink
1253
51 The Business of Polyethylene
Jorge O. Bühler-Vidal
1297
Appendix A1: Polymer Abbreviation Definitions
1331
Index
1333
Foreword
How is it that perhaps the simplest of polymer molecules, containing only carbon and
hydrogen atoms, has such versatility that it accounts for over a third of plastics usage?
Polyethylene is found in such diverse applications as films for food packaging, garbage
bags, industrial sacks, pallet wraps and agricultural use; containers for liquid packaging;
fuel tanks for automobiles; parts for toys; fibers and nonwoven fabrics for healthcare;
and pipes for water transport. Each of the thousands of applications using polyethylene
has its own specific needs that seemingly one molecule could not possibly meet. Yet
in the over 80 years since the high-pressure process to make low density polyethylene
was discovered by ICI researchers in England, the technologies used to manufacture,
modify and process this molecule have evolved to meet these needs. It is a story of
molecular manipulation on an industrial scale and the auxiliary technologies that have
been developed to foster its use.
There are very few molecular characteristics that can be modified to affect properties. These include the molecular weight, the molecular weight distribution (MWD),
long-chain branching, short-chain branching and short-chain branching distribution. Each of these, however, can have a substantial influence on properties. Consider
short-chain branching as an example. Increasing short-chain branching through introduction of alpha olefin comonomers (typically butene, hexane or octane) reduces the
ability of the chains to form organized structures or crystals. This produces the wide
range of density available in linear PE, and is responsible for producing tough, flexible
polymers on one end of the scale and stiff, temperature resistant polymers on the other.
The technology used to manufacture PE influences these five molecular characteristics.
The high pressure free radical polymerization process produces a molecular structure
with long-chain branches, which enhances melt strength and allows LDPE to excel at
being blown into film and extrusion coated. The drive for lower pressure processes
led to the development of catalysts that produce linear HDPE and LLDPE resins with
substantially no long-chain branching. Scaling up the catalyst polymerization process
led to specific manufacturing technologies, such as solution, slurry, and gas phase reactors, each with their own characteristic molecular fingerprint that influences properties. Recent developments in catalyst technology have allowed more precise control of
the molecular architecture, creating new product performance features. For example,
narrower MWD and comonomer distribution enabled by single-site catalyst technology produces lower seal initiation temperature and fewer extractables for flexible packaging. Further engineering of the molecular architecture has been achieved through
reactor design, such as combining reactors in series to create bimodal MWD polymers
for enhanced processability.
xi
xii Foreword
Manipulating the molecular architecture is not enough. Additives have been developed to quench the activity of catalysts, improve thermal oxidative stability, reduce flow
instabilities, prevent film blocking, reduce film friction, increase the rate of crystallinity, provide reduced resistivity, allow for foaming, retard flames, and provide color.
Additives are often the unsung heroes of polymers and play a critical role in fine-tuning
the polymer for specific applications.
Creating the polymer is only the first step; it must be transformed into a useful film,
sheet, container, part, profile, or fiber. The converting process plays an important role.
Understanding how the stretching, stress development, and cooling rate involved in the
converting process affects final part properties is critical for meeting application needs.
Many of the current processes did not exist or were in their infancy when polyethylene
first became available. These processes had to be developed alongside the resin to meet
the application requirements.
This book brings together a host of industry experts on polymer design, polymer
manufacturing processes, converting processes, additive technology, major applications, and regulatory, environmental sustainability, and supply chain issues to tell the
story of polyethylene. Each aspect of the technology goes together with the others. All
must be considered to understand the technology of polyethylene. The authors truly
are experts in their subject matter and have presented their material in a way that is
approachable to the industry practitioner. Academic and original source materials are
cited for the reader to look up if desired, but this book is designed for everyday use. No
other book on the market today brings all these elements together into a single resource
in such a useful way. Resin producers, additive suppliers, equipment manufacturers,
fabricators, and end-users alike will find this book invaluable for understanding this
polymer technology.
Recounting the story of polyethylene is important because of the role it plays in society. Its versatility and value-in-use is often overlooked by members of the community
who only see it as a “cheap” plastic for carry-out bags and the like. The low cost of polyethylene, driven by its technology, certainly has enabled it to be used in so many diverse
applications. The industry has challenges, however, to better communicate the benefits
of plastics and build upon the early efforts to improve sustainability. This book is an
outstanding resource for understanding the current state of polyethylene technology,
which has evolved considerably in its 80-year history. The book provides current and
future engineers with the knowledge to build upon this history with new technologies
and applications that will further enhance its growth and relevance in society.
Barry Morris
Technical Fellow
DuPont
August 2017
Preface
Polyethylene is the largest volume of polymer produced globally with about 84 million
metric tons produced. The success of polyethylene is due to a combination of its relatively
low cost to produce, the large scale of the production plants, numerous end-use applications, and the variability of the physical properties possible. The physical properties of
polyethylene are controlled by a nearly limitless combination of solid densities, structures, molecular weights, molecular weight distributions, and different comonomers. It
is the properties generated by these control factors that allow the resin to be processed
using a wide range of techniques and applications. Resins can have very low, solid densities and be elastic or they can have a very high, solid density and be very rigid. Resins
are produced that are viscoelastic and high molecular weight designed for blow molding or pipe applications to low molecular weight resins that are used for adhesives. The
innovations around polyethylene production, processing, and applications started in
the 1930s and continue to this day. No other polymer class exists that can compete with
the properties and applications of polyethylene.
This handbook provides a nearly complete description of polyethylene. The chapters
were written by some of the most experienced and prominent authors in the field, providing a truly unique view of polyethylene. The book starts with a historical discussion on
how low density polyethylene was discovered and how it provided unique opportunities
in the early days. New catalysts are presented and show how they created an expansion
in available products, including linear low density polyethylene, high density polyethylene, copolymers, and polyethylene produced from metallocene catalysts. With different
catalyst systems a broad range of structures are possible with an equally broad range of
physical properties. Numerous types of additives are presented that include those that
protect the resin from the environment and processing, fillers, processing aids, anti-fogging agents, pigments, blowing agents, and flame retardants. Common processing methods, including extrusion, blown film, cast film, injection molding, and thermoforming,
are presented along with some of the more specialized processing techniques such as rotational molding, fiber processing, pipe extrusion, reactive extrusion, wire and cable, and
foaming processes. The business of polyethylene, including markets, world capacity, and
future prospects, are detailed. This handbook provides the most current and complete
technology assessments and business practices for polyethylene resins with nearly 3000
references.
The concept and development of this handbook was initiated by Dr. Ananda M.
Chatterjee in April 2013. Without Andy’s attention to detail and his knowledge of the
PE industry this handbook would have never been started. On April 5, 2016, Andy
passed away at the age of 69 before this handbook was completed. Andy was a pioneer
in the field of PE technology, a friend, a great mentor, and he had a vision for the future.
xiii
xiv Preface
Andy started his career at Celanese Plastics (1978) and worked for Shell Development,
Union Carbide, The Dow Chemical Company, Kaneka Texas Corporation, and Ingenia
Polymers. He had extensive and in-depth experience in polymer research and development, product commercialization, polymer film processing and products technology,
polymer additives technology, additive blends technology and processing, customer
technical service and intellectual property invention. He was also an expert in licensing
and the technology transfer of UNIPOL Polypropylene Technology. Andy was a Fellow
and Honored Service Member (HSM) of the Society of Plastics Engineers (SPE). This
handbook is dedicated to Dr. Ananda M. Chatterjee and his vision for the future. In
late April 2016, I took over as the editor so that the project that Andy started could be
finished.
Kim Walton (Chapter 47) from The Dow Chemical Company passed away unexpectedly during the development of this handbook. He spent his career advancing
polyethylene technology and was an expert in polyethylene elastomers and plastomers.
He had broad impact on the commercialization of polymers and formulations to meet
customers’ needs. He was a forward-thinking scientist, a friend, and a passionate mentor. He led the technology development of polypropylene impact modification which
revolutionized the material design of bumper and fascia for the automotive industry.
This handbook would not have been possible without the contributions and worldclass expertise of the contributors; there were 91 authors from 38 different affiliations.
It was a great pleasure, a learning experience, and an honor to work with these folks.
Their knowledge and willingness to share are the keys to the success and acceptance of
this handbook.
Mark A. Spalding
Fellow
The Dow Chemical Company
August, 2017
List of Contributors
Francois Beaume is a senior development engineer at Arkema S.A. (France). He graduated in Chemistry
from Chimie ParisTech and received a PhD in Polymer Science from ESPCI ParisTech in 1996. He has
been with Arkema for over 20 years and is actively involved in Fluoropolymers R&D since 2002.
Tom Bezigian is a graduate of the UMASS-Lowell Plastics Engineering program, where he is now an
adjunct professor specializing in film and sheet extrusion. He is the author/editor of TAPPI’s Extrusion
Coating Manual. He has worked in various capacities in the industry for companies such as Cryovac, Mobil
Chemical, James River Corporation, and Felix Schoeller GmbH after which he founded and formed his
own specialty film and extrusion laminating business in Syracuse, New York. He now consults and teaches
on many aspects of the film and laminating industry, both here in the USA and abroad.
Steven Blazey received his PhD in organic chemistry from Case Western Reserve University in 1978 and
currently works in A. Schulman’s Technology Center in Akron, Ohio, as a Research and Development
Engineer developing new products for targeted applications. His career spans 30+ years formulating and
compounding polymers. He has a broad knowledge of plastic materials and additives and colorants as well
as polymer processing.
Thomas E. Breuer is retired from Witco Chemical/Chemtura, now PMC Biogenix. He was Director of
their Research and Applications Laboratory when he retired after 30 years of working in the field of polymer additives and other fatty chemical derivatives. He was involved in technical service for polymer additives worldwide for many years and was author or coauthor of many papers for SPE and other technical
organizations.
Jorge O. Bühler-Vidal studied Mechanical Engineering at the University of Buenos Aires and at the
Catholic University of America in Washington, DC. He joined Union Carbide in 1976 and was at the
Bound Brook laboratories from 1978 to 2001. He is currently the Director of Polyolefins Consulting LLC
and President of the Petrochemical Consulting Alliance which provides consulting services to global petrochemical and plastics industries.
Ian Burdett graduated from Manchester University Institute of Science and Technology, U.K. with a
PhD In Chemical Engineering in 1976. The majority of his working career was spent with Union Carbide
Corporation and subsequently The Dow Chemical Company. He worked in the area of Polyolefin
Process Technology for 31 years and became Director of Process R&D for Gas Phase Polyolefins. He was
awarded Union Carbide’s Chairman’s Award for development of the UNIPOLTM Gas-Phase Process for
Polypropylene in 1992.
Thomas I. Butler is retired from The Dow Chemical Company. While at Dow, he was a Technical Leader in
Polyolefins and Elastomers R&D, Technical Service and Development. He was responsible for fabrication
technology, including blown film processes, extrusion, and coextrusion. Tom graduated from New Mexico
State University in Chemical Engineering. He was involved in product and application development for
LDPE film, LLDPE film, ethylene acrylic acid copolymer film, and polyolefin plastomers.
xv
xvi
List of Contributors
Gina Butuc is the Technical Development Manager for Crosslinking Peroxides and Polymer Additives at
AkzoNobel, Functional Chemicals. Gina has more than 15 years of experience in the polymer field, either
polymer synthesis or application development and polymer testing, ranging from bench chemist to business management. She is the inventor or co-inventor on six U.S. and international patents and the author
of four technical articles and a book chapter. Gina holds a MSc in Organic Chemistry from University of
Bucharest, Rumania.
Amit K. Chaudhary is a Research Scientist in the Materials & Parts Processing Group at The Dow
Chemical Company, Midland, MI. He joined Dow in 2010 after completing a BSc (Chemistry) and BTech
(Chemical Engineering) from The University of Calcutta, India and a PhD (Chemical Engineering) from
Michigan State University. He has performed fundamental research in twin-screw extrusion based processes, developed composite materials and processes for application in transportation industry for mass
savings, and developed processes and compositions for water borne polymer dispersions for various industrial applications.
Bharat Chaudhary is a Fellow of the Society of Plastics Engineers and a Principal Research Scientist at
The Dow Chemical Company. He obtained his PhD and MSc from Imperial College, London, and a BEng
from the University of Benin, Nigeria, all in Chemical Engineering. He has 28 years of experience leading
research and development in a variety of areas related to polymer modification. Bharat has received several
awards for his work and he is author of 33 journal papers and 21 conference/technical presentations, as
well as inventor on 63 U.S. granted patents, 28 European granted patents and numerous other patents and
patent applications.
Alexander Chudnovsky is a UIC Distinguished Professor Emeritus and Director of the Fracture Mechanics
and Materials Durability Laboratory, The University of Illinois at Chicago. He has broad expertise in thermodynamics, fracture mechanics, materials science, stress analysis, probability and statistics. He is widely
recognized for his pioneering work in theoretical and experimental fracture mechanics, assessment of
materials durability, accelerated testing of lifetime as well as forensic studies.
Tim Clayfield has a BSc and PhD in Polymer Physics from the University of Manchester Institute of
Science and Technology. He joined Dow in 1989 in Stade, Germany in product and process R&D for the
TYRIN group and moved to DuPont Dow Elastomers (DDE) in 1996. At DDE, he became Business
Development Manager with projects across a range of markets. Returning to Dow and TS&D in 2005,
he worked in the NORDEL EPDM space where he was the Global Application Technology Leader and
Product Technology Leader for 8 years. He is currently a Technical Manager at Resinex.
Jeffrey Cogen is a Fellow at The Dow Chemical Company. He has 26 years of experience working with polymers, polymer additives, and polymer compounds. His expertise includes materials for communication and
power cables. He is an expert in formulation of polyethylene and in polymer additive mechanisms and applications. He is an inventor on 43 issued U.S. patents and has published more than 40 papers and two invited
book chapters. Dr. Cogen received a PhD in Organic Chemistry from the University of California, Berkeley.
Rongjuan Cong received a PhD in Chemical Engineering from McMaster University, Canada in 2001. She
joined The Dow Chemical Company in 2005, where she led the development of various new separation
techniques to characterize the microstructure of polyolefins and quantify the additives. She is currently a
Research Scientist in the Performance Plastics Characterization Group. She has published 40 peer-reviewed
papers. She is an inventor of 14 granted patents in new polyolefin products and polyolefin separation science.
Len Czuba is the President of Czuba Enterprises, Inc. a 2001, Chicago-based medical device product
development consultancy. He holds 15 U.S. patents and has given more than 100 presentations or seminars
and has more than 25 publications. In 2004, Mr. Czuba was one of the 100 ”Medical Device & Diagnostic
Industry” Notable Persons in the medical device industry. He has been a member of the Society of Plastics
Engineers since 1975 and is very active in the Medical Plastics Division. He was the Society’s International
President during SPE’s 2005 – 2006 year and is now a Distinguished, Honored Service, Fellow of SPE.
List of Contributors xvii
Samuel Devisme graduated as a Chemical Engineer from INP-ENSIACET, Toulouse, France and received
his PhD from École Nationale Supérieure des Mines de Paris in 2006. He is employed by Arkema as the
R&D Fluoropolymer Europe Manager and has worked in various areas of technical polymers for 10 years.
Christopher J. B. Dobbin is a Senior Scientist in the Polyethylene Product Research and Development
Team at NOVA Chemicals in Calgary, AB, Canada. He completed his PhD in Organic and Polymer
Chemistry at the University of Waterloo. In 1998 he joined NOVA Chemicals Corporation as a member
of its solution-phase single site catalysis program. Here he has played a fundamental role in the molecular
design, development and commercialization of the SURPASS family of dual reactor single site catalyzed
ethylene-octene copolymer resins.
Joseph Dooley received his PhD in Mechanical Engineering from Eindhoven University of Technology,
the Netherlands. He is currently the President of JDooley Consulting LLC, an organization that provides
consulting services on polymer processing operations, equipment design, and R&D to the plastics industry. Joe retired from The Dow Chemical Company after 35 years in Dow’s Research and Development organization. He has published more than 100 technical papers as well as authoring articles in the Encyclopedia
of Polymer Science and Technology and chapters in textbooks on polymer processing instabilities and coextrusion. Joe is a Fellow of the Society of Plastics Engineers.
Darin L. Dotson obtained his PhD degree in Chemistry from Virginia Polytechnic Institute and State
University in 1996 at which time he began working for Milliken and Company in Spartanburg, South
Carolina. His primary area of focus is the development of new nucleating agent additives for polypropylene (PP) and polyethylene (PE). He is the inventor or co-inventor on more than 40 U.S. and international patents.
Ronald S. Eisinger, a Chemical Engineer, has spent 25 years of his industrial career developing polymerization processes at Union Carbide Corporation and The Dow Chemical Company. His involvement in
the UNIPOLTM gas-phase, fluidized-bed process encompassed polyethylene, polypropylene, and ethylenepropylene rubber. He currently works on novel separations and polymerization processes at MATRIC
(Mid-Atlantic Technology, Research & Innovation Center).
Mohamed Esseghir is a Principal Research Scientist with The Dow Chemical Company. Mohamed holds
a PhD in Mechanical & Aerospace Engineering from Rutgers University and joined Union Carbide’s
Polyolefin Research then Dow Chemical Wire & Cable R&D. Mohamed has over 30 years of research
in polymer blends, including reactive modification and processing. He is currently responsible for new
material research for the Telecommunications segment in Dow’s Electrical & Telecommunications R&D.
Mohamed holds 20 granted patents, is inventor on numerous pending applications and the author of 25
journal publications and technical presentations.
Joe Fay currently works for BASF Corporation in the Applications Laboratory within the Plastic
Additives Business Unit, Tarrytown, NY. He was appointed a Technical Fellow in 2006 in recognition of
his significant contributions to the technology and commercial application of polymer stabilizers. His
current responsibilities include customer-focused technical service and applications development for
BASF’s plastic additives product range including antioxidants, processing stabilizers, light stabilizers
and other specialty additives. Joe obtained his PhD in Polymer Science and Engineering from Lehigh
University, PA.
Johannes Karl Fink is Professor Emeritus of Macromolecular Chemistry at Montanuniversität Leoben,
Austria. His industry and academic career spans more than 35 years in the fields of polymers, and his
research interests include characterization, flame retardancy, thermodynamics and the degradation of
polymers, pyrolysis, and adhesives. Professor Fink has published many books on physical chemistry and
polymer science including 9 with Wiley-Scrivener, amongst which are A Concise Introduction to Additives
for Thermoplastic Polymers (2009), Polymeric Sensors and Actuators (2012), The Chemistry of Biobased
Polymers (2014), and Metallized and Magnetic Polymers (2016).
xviii
List of Contributors
Tad Finnegan is the Technical Manager for the Plastic Additives Regional Application Center at BASF
Corporation, Tarrytown, New York. He is responsible for technical service and application development
for antioxidants, light stabilizers, non-halogenated flame retardants, and polymer modifiers for North and
South America. Tad obtained his BS in Chemical Engineering from Washington University in St. Louis,
and did further graduate studies in chemical engineering at the University of Delaware.
David Gillespie received his BS degree in Chemical Engineering from Drexel University in 1989. He has
spent the last 18 years with The Dow Chemical Company in their Plastics Characterization Laboratory
focusing on polyolefins analysis and in their Robotics Group focusing on high-throughput systems. David
has authored numerous papers in peer-reviewed journals, has numerous patents granted, and continues
to work actively within the chromatographic community to challenge, evaluate, and expand equipment,
detector, and software workflows for polymers.
Thomas Giovannetti holds a degree in Mechanical Engineering from The University of Tulsa and for the
last 21 years has worked for Chevron Phillips Chemical Company as a Technical Service Engineer for
injection molding. Tom started his plastics career by designing various injection molded products for the
chemical industry including explosion proof plugs and receptacles. Tom also holds a patent for design of
a polyphenylene sulfide sleeve in a nylon coolant cross-over of an air intake manifold and is a Certified
Plastic Technologist through the Society of Plastic Engineers.
John Gray LRSC Organic Chemistry spends his career in the UK, Germany and the USA working in surface chemistry in syntheses, production and applications. His expertise in polymer additives, personal care
skin surfaces and surfactants reflects his attention to molecular performance at the surface interphases in
these industries. The chemistry and dynamics of surface modifications and its application to real world
solutions are what has held his interest over many years. His current mission is to recognize these opportunities and apply new surfactant technology to challenging applications in today’s specialty surfaces for
industrial clients.
A. Willem deGroot received his PhD in Fiber and Polymer Science from North Carolina State University
in 1986 after which he began working for The Dow Chemical Company in their Plastics Characterization
Laboratory in Freeport, TX. He is currently a Fellow and is the Technical Leader for Dow’s Global Performance
Plastics Characterization and Testing Laboratory. He is on the Governing Board of the International
Symposium of Polymer Analysis and Characterization (ISPAC) and he has 31 years of industrial experience in Plastics Characterization focusing primarily on polyethylene molecular structure determination.
Bryan Hauger holds a PhD in Inorganic Chemistry from Indiana University. Bryan now operates a plastics consulting firm from Longmont, Colorado after a long career with Phillips Petroleum and Chevron
Phillips Chemical companies. He is an expert in plastic piping, materials and systems. Dr. Hauger is the
author of two ASTM standards and has received a Special Service Award from ASTM. He is currently
serving a three-year appointment to the ASTM Committee on Standards. He is currently the ANTEC Technical Program Chair for the Plastic Pipe and Fittings Special Interest Group.
Theresa Healy has been with Reedy Chemical Foam since 1997. She has more than 20 years’ experience
in the plastics industry. Her technical knowledge and sales management background helped to create
several Reedy programs which involved various business agreements and patented technologies. Since
2015, she has been working as the Technical Sales Director. She serves on the Society of Plastics Engineers,
Thermoplastic Materials & Foams Board of Directors, Secretary and Communications Chair.
Jim Hemphill was formerly with Dow Elastomers division of The Dow Chemical Company as their global
product technology leader for ENGAGE polyolefin elastomers and application technology leader for
elastomers used in thermoplastic polyolefins (TPOs). He joined Sadara Chemical Company in 2017 as
their Quality Director. Jim holds a bachelor’s degree in Chemical Engineering from New Mexico State
University.
List of Contributors xix
Wes Hobson is a Sr. Engineer in Dow Chemical’s Performance Plastics Technical Service and Development
group. Wes has 22 years’ experience in the plastics industry focusing on cast, blown and sheet extrusion.
Wes holds a Master’s Degree in Packaging Science from Michigan State University.
Sue te Hofstee is a Researcher at AkzoNobel Polymer Chemistry and currently active in the area of polymer
additives. After obtaining a Bachelor degree in Organic Chemistry, she started to work within Diosynth in
2003 and then joined AkzoNobel in 2011.
Nicholas Horstman is a Research Technologist in the Materials & Parts Processing Group at The Dow
Chemical Company, Midland, MI. He has been involved in research with various extrusion platforms since
2006 with his most notable accomplishments in the areas of reactive extrusion and process development.
Nick holds a BS in Chemistry from Central Michigan University.
Wenyi Huang is an Associate Research Scientist in the Materials & Parts Processing Group at The Dow
Chemical Company, Midland, MI. He received his PhD in Polymer Engineering from the University of
Akron and joined Dow in 2012. His expertise includes reactive extrusion, microcapillary technology,
porous films, polymer synthesis and nanotechnology. He is a certified Six Sigma Black Belt.
David Hukill is a Registered Professional Engineer and is the Pipe Applications Technical Service Engineer
in the Global Polyethylene Applications Group at Chevron Phillips Chemical Company. David holds a
Bachelor of Science in Civil Engineering from the University of Missouri and an MBA from Washington
University in St. Louis. He has over 24 years of experience in plastics R&D, plastics related construction
and standards work.
Ted Kampen is Account Manager at AkzoNobel Polymer Chemistry and acts as Global Key Account
Manager for several multinational customers in the polyolefin, PVC, rubber and crosslinking markets.
Before this he was active as Technical Development Manager Polymer Additives EMEIA. Ted has a background in Chemical Engineering and has obtained an MBA degree.
Gregory W Kamykowski is a Senior Applications Scientist for rheology at TA Instruments. He received
his PhD in physical chemistry from the University of Wisconsin – Madison, where he studied under the
renowned rheologist, Professor John Ferry. He has been the President of the Chicago Section of SPE and
the Chair of the Polymer Analysis Division and the ANTEC Technical Program Chair for the Polymer
Analysis Division. He is also active in ASTM and the Society of Rheology.
Larry Keefe has experience in the plastics additives markets and associated technology including product
design, new business development, and processing. Industries Larry has supported are thermoplastic and
thermoset compression and injection molding, structural foam molding, profile extrusion, and cast extrusion. Larry is a graduate of Clarkson University with a degree in Mechanical Engineering and has earned
an MBA from the State University of New York at Buffalo.
Kimberly Kern is a Market Manager for the PVC division and former manager of the applications laboratory with emphasis on additives for polyolefins at Baerlocher USA. Before coming to Baerlocher, Kimberly
held several R&D positions related to polymer formulations including: adhesives at Bioformix, composites
at the Air Force Reach Laboratory, and coatings at Sherwin Williams. Kimberly received her MS in polymer chemistry from Wright State University.
Hyunwoo Kim is a Research Scientist in Core R&D for The Dow Chemical Company. In this role, he
is responsible for leading product and process development efforts focused on polymer processing, die
design, extrusion, coextrusion and several other converting processes. Hyunwoo received a PhD in
Chemical Engineering from the University of Minnesota in 2009. His graduate research was on processing,
structure and properties of graphene reinforced polymer nanocomposites. He holds 2 U.S. patent applications and has authored 35 Dow internal publications and 18 external publications.
xx List of Contributors
Yoshitaka Kimura earned a BS degree in Applied Chemistry at the Ehime University, Japan. Since 1992, he
has been with The Japan Steel Works, Ltd. (JSW) as a process and mechanical engineer. His main responsibilities included planning of trials and designing processes for compounding, devolatilization, and dewatering applications. Since 2009, he has been with JSW America in Detroit Michigan. His responsibilities
included managing the Detroit Technical Center and expanding sales in the U.S. market. Since 2015, he
has been a group manager of the Extrusion Process Technology group, Plastics Machinery Department in
JSW Japan.
Rick King, is a Senior Technical Manager for the Plastic Additives Business Unit at BASF Corporation,
Tarrytown, NY. Rick did his graduate work at UCLA focusing on synthesis, characterization and chemistry
of transition metal based organometallics. After working for Union Carbide and Ciba-Geigy he became
part of BASF Corporation in 2009 when BASF purchased Ciba. He has been involved with additives for
plastics and stabilization technology for the last 30 years. Over the last 28 years, Rick has presented more
than 80 conference papers and is an author of 35 U.S. patents.
Roger Kipp’s contributions to the plastics industry include hands-on innovation in tooling and processing,
furthering education initiatives, and creating successful business development models. The Society of Plastics
Engineers has honored Kipp with the 2002 Outstanding Achievement Award, a Lifetime Achievement
Award in 2003, the Honored Service Member Award in 2008 and the Thermoformer of the Year Award in
2010 and Presidents Cup in 2017. In 2012 Kipp was inducted into the Plastics Pioneers Association and is
currently the Chairman of the SPE Foundation where he continues his support of education.
Yury V. Kissin received his PhD degree in Polymer Chemistry in 1965 at Institute of Chemical Physics in
Moscow. From 1960 until 1977 he worked in Institute of Chemical Physics studying kinetics of polymerization reactions of ethylene, propylene and higher α-olefins and the structure of polyolefins and ZieglerNatta catalysts. He emigrated to USA in 1979 and worked at Gulf Research and Development Company
in Pittsburgh, PA, Edison Research Center of Mobil Chemical Company in Edison, NJ, and at Engelhard
Research Center in Iselin, NJ. Since 2000 he is a Visiting Scientist at Department of Chemistry of Rutgers
University in NJ. He has authored four books, twenty articles in chemical and polymer encyclopedias,
~250 scientific articles, and ~70 patents in the field of synthesis of Ziegler-Natta and metallocene catalysts.
Nikhil Koganti is a Pipe Applications Technical Service Engineer in the Global Polyethylene Applications
Group at Chevron Phillips Chemical Company. Nik holds a BS in Chemical Engineering from the
University of Oklahoma, received in 2014. He has over 3 years of experience in polyethylene pipe technical service, which includes product development, general research and design, manufacturing, and
commercialization.
Kurt A. Koppi received a BS in Chemical Engineering at the University of Illinois and a PhD in Chemical
Engineering at the University of Minnesota. Since 1993 he has been employed at The Dow Chemical
Company in Midland, MI. He currently works as a Sr. Research Scientist and leads the Applied Rheology
and Die Design technology area within Core R&D at Dow.
Glenn Larkin holds an MBA and a BS in Business Management from the University of Phoenix. He joined
Chevron Phillips Chemical Company in 1984 and now provides technical services for rotational molding.
He also holds experience in electronics assembly, and biotechnology, along with the operation of polyethylene, polypropylene pilot plants. Glenn became a Certified Plastics Technologist through the Society of
Plastics Engineers in 2000 and served as the President of the Oklahoma Section.
Jennifer Lynch-Branzoi is on the research faculty in the Materials Science and Engineering Department
at Rutgers University. Dr. Lynch develops novel processing methods for thermoplastic polymer composites
and nanocomposites that impart enhanced mixing and superior performance. She has been involved with
the development, processing, characterization, and application of reinforced thermoplastic composites
sourced from recycled plastics for 20 years and these Rutgers patented, commercially successful materials
have been used as railroad ties, pilings, I-beams, and high load capacity bridges.
List of Contributors xxi
Lisa Madenjian is the Global Product and Application Technology Leader for INFUSE and INTUNE
OBC and Dow Elastomers. She has been responsible for development and growth of new to world polymers, marrying materials science to application development. She has successful leveraged technology
and applications around the globe. She has been with Dow since 1989 in R&D/TS&D. She has worked on
Polycarbonate and Blends, Polypropylene, Elastomers and New Business Development. Lisa graduated
with a MS in Inorganic Chemistry from the University of Missouri.
Pamela Maeger is the Pipe Applications and Technical Service Supervisor in the Global Polyethylene
Applications Group at Chevron Phillips Chemical Company. Pam holds a BS in Chemical Engineering
from Louisiana Tech University. She has over 39 year of experience in polyethylene product development,
blow molding technical service, geomembrane technical service and pipe technical service. She has 6 U.S.
and foreign patents and has published over 20 journal and conference papers.
Sanjiv R. Malkan has been the Associate Engineering Manager of Textiles, Nonwovens & Adhesives at
Hunter Douglas since 2005. He received his PhD in Textiles and Polymer Science, an MBA in Finance,
and the CText FTI title conferred by The Textile Institute, Manchester, UK. He has 3 U.S. patents, 4 design
patents, and more than 50 publication and presentations.
Jill Martin is a Fellow in the Packaging and Specialty Plastics business of The Dow Chemical Company.
For the past 22 years, she has served in both research and development and technical service for the North
America Performance Plastics business. Her responsibilities include new product development and technology exploration for the packaging value chain building strategic relationships with downstream partners. Martin is past chair of the South Texas’ division’s International Polyolefins Conference. She received
her Doctorate in Polymer Engineering and Science from Case Western Reserve University in Cleveland,
Ohio and holds over 20 patents.
Bryonna Martinez has a Bachelor’s of Science in Chemical Engineering from the University of Oklahoma.
She graduated in 2014 and began work at Chevron Phillips Chemical in the same year. Currently, Bryonna
works as the liquid food packaging blow molding technical service representative at Chevron Phillips
Chemical in the Global Polyethylene Applications Group.
Kevin P. Menard received his PhD from the Wesleyan University and has worked in refining, polymers,
aerospace, and the instrumentation industry for over 40 years. He has over a dozen patents, over 100
publications, and authored “Dynamic Mechanical Analysis: A Practical Introduction.” He currently
works for Mettler-Toledo LLC as a technical application consultant and part time for Veritas Testing and
Consulting, LLC.
Noah R. Menard received his BS from University of Texas at Dallas. He has worked in thermal analysis
for over a decade including research at University of North Texas, Xavier University, Colorado School of
Mines, and Southern Methodist University and has interned for PerkinElmer Inc. He is a cofounder of
Veritas Testing and Consulting, LLC, where he is currently lab manager.
Randy J. Minton joined Reedy Chemical Foam in October of 2012 as New Product Development
Coordinator. In 2014 Randy became Chief Scientist in charge of Research, Development & Engineering.
He holds a BS in Chemistry from Virginia Polytechnic Institute and State University and oversees quality
control including ISO compliance.
János Móczó studied chemical engineering at the Budapest University of Technology and Economics,
Hungary, with a special emphasis on polymer science and received his PhD in 2005. He works in the
Institute of Materials and Environmental Chemistry, Research Centre for Natural Sciences, Hungarian
Academy of Sciences. His research field includes interfacial interactions in polymer composites, natural
and biodegradable polymers, micromechanical deformation of polymer composites.
xxii List of Contributors
Adriana I. Moncada is an Associate Research Scientist in the Materials & Parts Processing Group at The
Dow Chemical Company, Midland, MI. She joined Dow in 2010 after completing a PhD in Chemistry
from Texas A&M University. She has performed fundamental research in the areas of polymer synthesis, polymer modification, reactive extrusion and coating process technology. She is a certified Six Sigma
Green Belt.
Barry A. Morris is a Technical Fellow in the Ethylene Copolymers business of DuPont. He received
his PhD in chemical engineering from Princeton University. He also has a MBA from the University of
Delaware. During his career at DuPont he has held a variety of positions in R&D, technical service, and
application and new business development with a primary focus on packaging and polymer processing. He
has 10 patents and over 100 publications primarily in the use of ethylene based polymers in flexible packaging.
Cliff Mure is a Product Technology Leader with Univation Technologies, a wholly-owned subsidiary of
The Dow Chemical Company. He has more than 25 years of experience in the polyethylene industry starting with Union Carbide and then Dow Chemical in 2001. He has worked in R & D and technical service
for HDPE products and applications including blow molding, rigid packaging, and pipe extrusion. Cliff is
an inventor on numerous patents on polyethylene materials technology and earned BS and MS degrees in
Mechanical Engineering from Rutgers University.
Bjarne Nielsen has a chemistry background and has more than 20 years’ experience with polymer additives. He works for Palsgaard in Denmark having worked previously for DuPont and Danisco. His main
focus has been antistatic and antifog solutions for a wide variety of polyolefin applications resulting several
patents covering the use of novelty chemistries. His most recent work has focused on the development of
new bio-based material for several different applications, including phase change materials, lube oil and
fuel additives and solutions to improve slip additives that is used in packaging to reduce food waste.
Thomas J. Nosker is the Principal Investigator of Rutgers’ Center for Advanced Materials via Immiscible
Polymer Processing (AMIPP.) His research interests were initially focused on developing plastics recycling and related technologies, based on modifications made to traditional equipment to develop useful
products with new compositions, resulting in over a dozen technologies that have all been licensed to
industries. More recently, he has focused on virgin engineering plastics and composites, developing new
and advanced materials with new compositions, especially graphene-polymer composites made directly
from polymers and graphite. Dr. Nosker has been honored as one of only 20 Rutgers Revolutionaries, as
part of Rutgers’ 250th anniversary. This is a list of “Rutgers people and innovations that have changed lives
around the world.”
Ivan Ong is the V.P. of R&D at Microban International, Ltd., a global leader in providing customized
antimicrobial solutions for consumer and commercial products. For the past 20 years he has participated
in active research in profiling numerous antimicrobial technologies for microbiological activity and suitability for application in polymers, textiles and ceramics to satisfy performance and sustainability needs in
key areas such as health care, food, commercial, institutions and home. Ivan obtained his PhD at the Johns
Hopkins University in Materials Science and Engineering in Baltimore, Maryland.
Tor Palmgren, independent consultant, holds a PhD in Polymer Chemistry from The University of
Connecticut. His career of over thirty years includes 9 years as Global Product Compliance Manager at
Chevron and 6 in water treatment with Schlumberger. Patents, publications and presentations are on polyolefins, polymerization and water treatment. He has been active with SPI/HSE and chaired ASTM subcommittee, spearheading development of elemental standard for polyolefins.
Rajesh Paradkar received his PhD in Analytical Chemistry from Clemson University in 1996. Rajesh
joined The Dow Chemical Company in 2002 and is currently a Research Scientist in the Performance
Plastics Characterization Group. His area of expertise is vibrational spectroscopy and microscopy and he
has applied these techniques for approximately 15 years to solve a wide range of R&D, TS&D and manufacturing problems for Dow’s polyethylene business.
List of Contributors xxiii
Charles Park works closely with A. Schulman (ASI) Research & Development Engineers to develop new
products as well as improve current products. He then transfers this information to specific ASI manufacturing sites globally and oversees initial production efforts. In addition, Charles assists ASI manufacturing sites with respect to process optimization and troubleshooting. Charles has a Bachelor of Science in
Chemical Engineering from the University of Akron.
Rajen Patel joined Polyolefins R&D of The Dow Chemical Company in June 1991 in the Materials Science
group and transferred to Polyolefins TS&D of The Dow Chemical Company in November 2007. He is currently an Associate Research Fellow in the Packaging and Specialty Plastics technical service and development (TS&D) group in The Dow Chemical Company leading packaging applications development.
He has co-authored 22 technical peer-reviewed journal publications and 7 book chapters. He is also a coinventor of 51 granted U.S. patents. In 2009, he was elected as Fellow of Society of Plastics Engineers (SPE)
in recognition of his technical and commercial accomplishments in Polyolefins and especially in single-site
catalyzed (metallocene) Polyolefins.
Timothy Person is a Principal Research Scientist at The Dow Chemical Company. He obtained his PhD in
Chemical Engineering from the University of California at Berkeley and has 21 years in research and development of materials for use in power transmission and distribution applications, with particular focus
on polymer compounding, polymer crosslinking, electrical aging and breakdown phenomena. Tim is the
author of 55 publications/presentations, has 10 granted patents, and is a recipient of the R&D 100 award.
He is also a member of the IEEE Power Engineering Society, Insulated Conductors Committee, IEEE
Standards Association and CIGRE, where he has made contributions to various working groups in the
development of industry guides and standards.
Rudolf Pfaendner studied chemistry at the University of Bayreuth in Germany and received his PhD in
1985. He is currently Division Director Plastics within the Fraunhofer Institute of Structural Durability
and System Reliability LBF in Darmstadt, Germany. Rudolf has over 30 years of experience in polymer
synthesis, plastics and coatings additives, stabilizers, recycling, reactive extrusion, nanocomposites, flame
retardants and innovation management including numerous patents and publications.
Jason Pomante is the North American Market Manager for Arkema. He graduated from Lafayette College
with a Chemical Engineering degree and has an MBA from Saint Joseph’s University, Philadelphia. Jason has
been with Arkema for 15 years most recently in the Fluoropolymers group working in various roles. Jason
is an active member of the Society of Plastics Engineers as a past President and current Treasurer of the SPE
Philadelphia Section.
Benjamin Poon currently works for Baxter Healthcare Corporation in R&D. Prior to Baxter, he was
a Senior Research Engineer at The Dow Chemical Company from 2003 to 2009 working in polyolefin
research. Dr. Poon received his PhD in Macromolecular Science and Engineering from Case Western
Reserve University He holds 9 U.S. patents and is author of 26 journal articles. Dr. Poon also serves on the
board of the Medical Plastics Division of the Society of Plastics Engineers since 2013.
Michele “Mick” Potenza is responsible for the Sales of Polymer Additives at Corbion. He has been working in the specialty chemicals and polymer business for the past 20 years. Before joining Corbion, he held
Business and Sales Management positions at companies such as DuPont and Ciba Specialty Chemicals
among others. He is the author of various articles in technical magazines. Mick has a BS in Chemical
Engineering from the University of Rome “La Sapienza” and an MS in Financial Management from the
University of Turin.
Béla Pukánszky has been a full professor since 1994 and full member of the Hungarian Academy of
Sciences since 2010. He works in the Laboratory of Plastics and Rubber Technology, Department of
Physical Chemistry and Materials Science, Budapest University of Technology and Economics, as well as in
the Institute of Materials and Environmental Chemistry, Research Centre for Natural Sciences, Hungarian
Academy of Sciences. He has published more than 300 papers and has received approximately 4700 independent citations so far.
xxiv List of Contributors
N.S. Ramesh is the Senior Engineering Fellow in Product Care R&D of Sealed Air Corporation. Before
this role, he was the Director of R&D for BRIC region and Foams Expert at Sealed Air involved in
Kevothermal and Dow Ethafoam M&A activities. Dr. Ramesh holds 25+ U.S. patents, developed and
commercialized several new products, and has co-authored 3 books and written more than 100 technical
publications in international journals. He was inducted into Sealed Air’s Inventor Hall of Fame in 2006.
He has won the SPE Best Paper Award for Foams R&D from the Society of Plastics Engineers 3 times. He
serves in SPE’s Thermoplastic Materials & Foams Board and was elected to be SPE Fellow for technical
contributions to foams industry and Honored Service Member for promoting the knowledge of plastics
worldwide.
Jon Ratzlaff holds a BS in Chemical Engineering from Oklahoma State University (1985). He has former
experience in blown and cast film and has worked for Chevron Phillips Chemical Company since 1989.
He currently is a manager for technical services for rotational and injection molding. Jon holds a patent
on polymer compositions, authored over a dozen papers and co-authored the “Injection Molding Toolbox”.
He is a Past President of SPE (2013) and of the Rotational Molding Division and Oklahoma Section. He is
a long standing member of the Injection Molding Division.
Roger A. Reinicker earned a BS in Chemical Engineering from Carnegie-Mellon University and an MChE
from the University of Delaware. He spent 41 years in the plastics industry, chiefly in the pigmentation
and coloration of plastics and customer technical support, retiring in 2014 from BASF. He wrote numerous
papers for the Society of Plastics Engineers Color and Appearance Division, and became a Fellow of the
Society in 2016.
Vivek Rohatgi is in the Global Pipe Technical Service role with Chevron Phillips Chemical Company.
He holds MS and PhD degrees in Chemical Engineering from The Ohio State University. He has over 18
years of experience in plastics R&D, manufacturing and commercialization. He has 8 published patents, 14
journal and conference papers and 2 book chapters in polymers and processing.
Peter Schroeck joined Reedy International Corporation in 2004 as a Sales Account Manager. Peter was
named Technical Sales Director in 2007 and was responsible for training in-house sales representatives,
agents and distributors in Mexico, China, Canada and South America. In 2013 Peter was promoted
to Chief Business Development Officer, and in 2014 was named President and CEO. Peter is an active
member of the Society of Plastics Engineers, Plastics Industry Association, and American Chemical
Society.
Kalyan Sehanobish is a Sr. Research Fellow in Dow Performance Plastics business at The Dow Chemical
Company, Midland, MI. His career started in 1987 in Core R&D function after completing his PhD from
CWRU in Polymer Science & Engineering followed by MS in Chemical Engineering. He was one of the
pioneering team that began structure-property-performance understanding of polyethylene to define
application space for Dow metallocene polyethylene. Numerous product intellectual assets are associated
with him in his 30 years career and key to some of their launches.
David A. Seiler received his BS in Chemical Engineering from Penn State University in 1983. He is
employed by Arkema Inc as the Americas Business Manager Industrial / Global Technical Advisor
Fluoropolymers and has worked in the area of Fluoropolymers for 34 years. David is currently a voting
committee member on ASME-BPE polymers group subcommittee & ASTM E – Fire Standards main and
sub groups.
Robert Sherman is the Technical Director for the Special Additives division at Baerlocher USA. Before
coming to Baerlocher, Robert held several R&D positions related to polymer additives and stabilization in
different resins including: acrylics at Aristech Acrylics, polyethylene at LyondellBasell, and polyesters at
GE Plastics. Robert received his PhD in Chemistry from Oklahoma State University, did Post-Doctoral
research at Texas A&M, and BS/MS in Chemistry at Southern Illinois University.
List of Contributors xxv
Sami Siddiqui holds a Master’s Degree in Chemical Engineering with specialization in Polymer Processing.
He has been working with Ford Motor Company for last 20 years and has been involved in injection molding and blow molding CAE simulations, process optimization and troubleshooting. He is currently leading
the Manufacturing CAE section of Ford Powertrain Installations. Before joining Ford, he worked at BASF
as a polymer processing engineer.
Steven R. Sopher has over 30 years of experience in the plastics field. He has held a variety of technical
positions at Raytheon, Dow/United Technologies, ARCO Chemical, and JSP where he is currently the
Technical Director. He is also an instructor for TMS (The Minerals, Metals & Materials Society) teaching Plastics and Statistics. Steve holds Bachelor’s and Master’s degrees in Mechanical Engineering from
Northeastern University in Boston, MA. Steve is the author of over 30 papers in the area of polyolefin
foamed plastics and holds 2 patents in the area of polyolefin foam energy management.
Mark A. Spalding is a Fellow in the Materials & Parts Processing Group at The Dow Chemical Company,
Midland, MI. He joined Dow in 1985 after completing a BS from The University of Toledo and a MS
and PhD from Purdue University, all in Chemical Engineering. He has performed fundamental research
in single-screw extrusion, developed methods to measure resin physical properties that are important to
polymer processing, developed numerous techniques to debug and increase the rates of extrusion lines,
and has developed new mathematical models for extrusion simulation.
Gea Spijkerman is a Researcher at AkzoNobel Polymer Chemicals and is currently active in the area of
thermoset chemicals. Before the move to thermosets, she spent 10 years working in the field of crosslinking
peroxides and polymer additives. After obtaining her Bachelor degree in Chemical Technology she started
her employment with AkzoNobel in 1989.
Mohammad Usman has been at Ford Motor Company for the last 26 years. Currently he is leading CAE,
Plastics Manufacturing and Material Department of Ford Powertrain Installations. He has extensive experience in design and manufacturing of automotive plastics components, and specifically blow molded
products. His automotive experience at Ford includes introduction of multi-layer fuel tanks, twin sheet
fuel tanks, split parison new generation fuel tanks, use of 3D blow molding for hoses and ducts, and blow
molding simulation software development at CNRC. Dr. Usman holds an MS in Civil Engineering, an MS
in Mechanical Engineering, a PhD in Mechanical Engineering and an MBA from Ross School of Business,
University of Michigan.
Kim Walton was a R&D Fellow for The Dow Chemical Company. He is a polymer materials scientist with
a proven track record of commercial innovation for over thirty years. He has a wide range of market and
application experience including packaging, transportation, and the construction industry. He successfully commercialized polymers and formulations tailored to market needs. He is an author or coauthor of
45 U.S. patents and over 20 articles and conference presentations. Kim graduated from the University of
Southern Mississippi with a BS in Polymer Science. He received a MS in Macromolecular Science from
Case Western Reserve University.
Scott Wasserman is an Associate R&D Director and Intellectual Capital Manager with The Dow Chemical
Company. He has more than 23 years of experience in the study of polyethylene, with 19 of those years
applied to product and technology development for wire and cable. He has spent many years on the leadership team for the International Wire & Cable Symposium, with roles as Chairman of its Symposium
Committee and presently the Vice-Chairman of its Board of Directors. Dr. Wasserman earned his PhD
in Chemical Engineering from Princeton University in 1993 and his B.Ch.E. with Distinction from the
University of Delaware in 1988. He is an inventor for 10 granted U.S. patents and the author of numerous open literature publications and is a member of the Society of Plastics Engineers and the Society of
Rheology.
xxvi
List of Contributors
Jeff Wooster is the global sustainability director for Dow Packaging and Specialty Plastics. In this role, he
collaborates with the entire value chain to promote and improve the sustainability value of plastic packaging. He is the past president of AMERIPEN and is on the Board of Directors for GreenBlue. He serves
on the steering team for the Ocean Conservancy’s Trash Free Seas Alliance signature initiative on marine
debris and as Co-Chair of the World Business Council for Sustainable Development’s Roadmap to Curb
Ocean Waste (ROW) initiative. Additionally, Wooster serves as chair for American Chemistry Council’s
Plastics Division Packaging Team and is a member of the Flexible Packaging Association’s Sustainability
Task Force. Wooster joined Dow in 1988 and holds 45 U.S. and foreign patents and has published more
than 100 technical papers and presentations. Wooster received a Bachelor of Science degree in Chemical
Engineering from Iowa State University of Science and Technology.
Z. Zhou received his PhD in Engineering from Kyoto University, Japan in 1997. After working as a postdoc
in McMaster University of Canada and an NMR facility manager in Louisiana State University, he joined
The Dow Chemical Company in Analytical Science Department in Freeport, TX in 2002. He has worked
on using NMR to characterize the molecular structural features of polyolefins for approximately 15 years.
He is currently a Research Scientist and has been leading the NMR technology developments for polyethylene characterization. He has published over 100 external peer-reviewed and Dow internal research papers
related to polyethylene microstructure characterization.
Feng Zuo is a Senior Scientist in the Plastic Additives Business Unit at BASF Corporation, located in Tarrytown,
NY. Feng graduated from State University of New York - Stony Brook with a PhD in Chemistry, and did his
Post-doc study at University of Minnesota, Minneapolis. He joined BASF in 2012 and has been working on
various aspects of polymer stabilization and new product development with focus on automotive applications.
Part 1
PRINCIPLES AND PROPERTIES
OF POLYETHYLENE
1
An Industrial Chronology of Polyethylene
Christopher Dobbin
NOVA Chemicals Corporation, Calgary, Alberta, Canada
Contents
1.1 Overview ........................................................................................................................................ 4
1.2 The Early Years .............................................................................................................................. 4
1.3 High Pressure Polyethylene ......................................................................................................... 5
1.4 The Advent of High Density Polyethylene ................................................................................. 8
1.5 Product and Process Proliferation ............................................................................................ 11
1.6 Single-Site Catalysts Arrive ........................................................................................................ 14
1.7 The Future of LDPE .................................................................................................................... 21
References .............................................................................................................................................. 21
Abstract
Polyethylene (PE) is the single largest volume polymer produced worldwide. With global production estimated at 84 million tons per year and with a total population exceeding 7 billion
individuals, the annual consumption of PE resin globally is about 23 pounds per person per year.
Since its discovery and commercialization in the years immediately preceding World War II,
the story of PE is one of serendipity, innovation, dedicated research and development, marketing, and new business development. Landmarks along the way have involved stepwise improvements in the polymerization processes with lower capital costs and milder operating conditions,
and the development of catalysts with vastly superior activity and specificity. New technologies
continue to delineate and redefine the state of PE products and processes. The invention and
commercialization of metallocene/single-site catalysts in particular have provided the ability to
tailor PE resins to an extraordinary degree. The future of PE appears to be bright.
Keywords: History, discovery, Ziegler-Natta, metallocene, plastomer, ICI process, solution
phase, homogeneous catalyst, heterogeneous catalyst
Corresponding author: cjb.dobbin@gmail.com
Mark A. Spalding and Ananda M. Chatterjee (eds.) Handbook of Industrial Polyethylene and Technology, (3–24)
© 2018 Scrivener Publishing LLC
3
4
Handbook of Industrial Polyethylene and Technology
1.1 Overview
Polyethylene is a fundamental part of everyday life. It is used in bulk and retail food
packaging where it helps to extend shelf life and to provide a high level of food safety.
It is an essential component in electronics, high voltage power cabling, fiber optic systems, and digital communication networks. Other infrastructure applications include
pressurized potable water tubing, irrigation tubing, pipe liners and coatings, and drainage tile. PE is extremely versatile and can be found in household durable goods, children’s toys, consumer packaging, beverage caps and closures, trash bags, and can liners.
It is also widely used in industrial packaging systems, storage tanks, reusable pallets,
composite decking materials, sealants for solar panels, and artificial turf.
Perhaps not surprisingly, PE is the highest volume polymer manufactured worldwide. Global PE consumption for 2014 has been estimated at 84 million tons, or about
168 billion pounds, an increase of 3.8 percent over the preceding year.
Since the inception of the product over 80 years ago, PE manufacturing processes
have undergone a tremendous evolution. The polymer was observed experimentally
several times during the late 19th and early 20th century, but it wasn’t until 1933 that it
was accidently “discovered” by researchers at Imperial Chemical Industries (ICI) who
managed to commercialize what was later to become known as low density PE (LDPE).
The product was highly branched and exhibited limited physical properties, but nevertheless played a key role in Britain’s World War II combat effort. Over the ensuing years,
new combinations of catalysts and processes led to the development of essentially linear
high density products with improved physical properties. In the 1960s and 70s, new
catalysts promoted the incorporation of alkene comonomers, that allowed the further
customization of product density, flexibility, and toughness. Once again, the range of
products was expanded.
New technology developments continued to delineate and redefine the state of PE
products and processes. In the last decades of the 20th century, the invention and
subsequent commercialization of metallocene/single-site catalysts provided a level
of specialized products ranging in density from 0.857 to 0.967 g/cm3 and exhibiting properties from elastomeric to rigid gas barrier materials. Commercial processes
have also emerged that use multiple reactors to tailor molecular weight distribution
characteristics to remarkable degrees.
1.2 The Early Years
The first known laboratory synthesis of PE was reported by the German chemist Hans
von Pechmann in 1898. While performing an experiment in the laboratory that involved
the heating of diazomethane, von Pechmann observed a white, waxy decomposition product [1]. Subsequent analysis by colleagues Eugene Bamberger and Friedrich
Tschirner indicated a basic hydrogen and carbon composition with CH2 repeat units
[2]. They dubbed the new material “polymethylene.” In principle, the newly described
polymethylene could contain any number of carbons in the polymer backbone while
PE consists of CH2 CH2 paired repeat units. Despite this technicality, von Pechmann
and colleagues are credited with the first documented discovery of PE.
An Industrial Chronology of Polyethylene 5
In 1929, M.E.P Friedrich and Professor Carl S. Marvel reported a low molecular weight form of the same waxy white material while investigating the reactions
of n-butyl lithium and ethyl lithium with ethylene gas at the University of Illinois
(Urbana) [3]. The solid material was not analyzed further, but the authors recognized
that the alkyl lithium likely resulted in the polymerization of ethylene to give nongaseous products.
At about the same time, Marvel was invited to become a consultant to the DuPont
Experimental Station in Wilmington where his responsibilities included working with
a small research group led by one of his former students, Wallace H. Carothers [4].
Carothers and his team were subsequently credited with the discovery of neoprene
synthetic rubber, high molecular weight polyesters, and polyamides. In 1930, however, Carothers was interested in the stability of linear alkanes, and constructed a distillation apparatus to isolate linear alkanes generated by condensation reactions with
up to 70 carbons in the chain. Carothers observed the increase in melting point with
chain length and speculated that materials with molecular weights as high as 200,000
were certainly possible and would display interesting properties [5]. The condensation
approach to PE production was industrially impractical, but Carothers’ experiments
likely laid the ground work for DuPont’s later involvement in the PE industry.
1.3 High Pressure Polyethylene
The first industrially significant synthesis of PE was conducted (again, serendipitously)
by Eric W. Fawcett and Reginald O. Gibson at the ICI Winnington facility in Northwich,
England. The high pressure experiments that eventually led to commercial PE were
begun on March 24, 1933. The work was part of a program designed to study the effects
of high pressures (>1000 atmospheres) on chemical reactions. Although the experimental work was being carried out at Winnington, the home of the ICI Alkali Division,
the project was being sponsored by the company’s dyestuff division, in the hopes that it
would provide new routes to organic intermediates.
Applying extremely high pressure (1700 to 2000 atmospheres at 170 °C) to a mixture
of ethylene and benzaldehyde produced a small amount of white, waxy solid in the ethylene inlet to the reaction vessel. Fawcett reportedly identified the material as a polymer
of ethylene. The material could be repeatedly melted and drawn into threads. Because
the reaction had been initiated by trace oxygen contamination in their apparatus, the
experiment was extremely difficult to reproduce.
In conversations published in 1983 on the 50th anniversary of their auspicious
experiments [6], Fawcett and Gibson reported that they began their first experiment
with ethylene and benzaldehyde, principally because the materials were on hand. ICI
had scheduled five of the six proposed high pressure reactions with carbon monoxide,
but the CO tanks hadn’t arrived on site yet. A tank of ethylene was available however,
so they began their evaluations with that. Gibson and Fawcett pointed out that the
equipment was primitive and difficult to work with. When the pressure in the apparatus
decreased overnight, they feared that a leak had developed. After several experiments,
Fawcett noticed scraps of white material on parts of the reactor tubing. Over the following four months, the two chemists obtained enough of this material to positively
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Handbook of Industrial Polyethylene and Technology
identify it as a polymer of ethylene. In subsequent experimentation, the reaction was
plagued by poor reproducibility and frequent ethylene decompositions. After four
months of work, the project moved on to other high pressure experimentation and PE
was virtually forgotten.
Fawcett announced the discovery of the new polymer at a scientific meeting at
Cambridge in 1935, but the work was essentially ignored. During the 1983 interview,
Fawcett remarked that this was not unexpected as “it was a polymer so unlike the polymers known at the time that no one could envisage a use for it.”
When the high pressure experimentation produced little interest, the Dyestuff
Division eventually withdrew its support for the venture. Fawcett, who was seconded to
the project, was moved on to other work.
Michael Perrin, a young researcher who had just returned to Winnington from
Amsterdam took Fawcett’s place. Gibson was subsequently moved on to other work
and Perrin took charge of the high pressure experimental program in late 1935. Perrin
was interested in fundamental reaction mechanisms and decided to reexamine ethylene on the basis of its physical and chemical characteristics.
The work was actively pursued, but required the installation of safer laboratory equipment. Working with an improved high pressure reactor, Perrin and his team were able
to produce several grams of PE. By 1936, he and his colleagues had made significant
strides in understanding the role of trace oxygen impurities in initiating the polymerization, in reproducing the polymerization, and in limiting the ethylene decomposition
that had beleaguered Fawcett and Gibson’s work.
A British patent application was submitted by the ICI research team and accepted
as a provisional specification in early 1936. The final patent document was issued in
1937 [7]. The patent described the oxygen-initiated polymerization of ethylene to high
molecular weight PE and also anticipated the production of lower molecular weight
oligomers for lubricating greases. The free-radical initiated polymerization process
relied on low levels of oxygen (<10 ppm) in the ethylene feed gas. Perrin and his
research group had successfully developed this “experiment gone wrong” into a reproducible high pressure synthesis for PE.
With laboratory samples in hand, Perrin began to explore the commercial possibilities for the new material. Serendipity once again played an important role here. As a
matter of course, a sample of PE was passed on to the ICI Dyestuff Division, who had
supported the original program several years before. The sample fell into the hands
of a chemist who had recently joined ICI after working for a firm that made submarine telephone and telegraph cables. He recognized its similarity to gutta-percha, a
naturally occurring latex rubber that was the standard jacketing for cables at the time.
Subsequent testing showed that PE was a tough, resilient material with a high dielectric
constant and low loss factor, excellent insulation properties, and exceptional moisture
resistance, which made it far superior to gutta-percha. The cable manufacturer placed a
100 ton order for the new material—and this was sufficient inducement for ICI to build
a production facility on the Winnington site.
The world’s first PE plant came on stream in the UK on September 2, 1939. Because
of its low loss factor at ultrahigh radio frequencies, PE rapidly became a strategic wartime materiel. It played a pivotal role in the development and rapid deployment of
radar during World War II, allowing the Royal Air Force Command to detect incoming
An Industrial Chronology of Polyethylene 7
German aircraft while still over the English Channel and to position their defenses
accordingly. PE is said to have allowed for such weight reductions in radar equipment
that it facilitated the development of airborne radar for aircraft patrolling the Channel
and the North Atlantic.
The initial Winnington PE plant was essentially a large pilot plant or semi-works scale
facility, capable of producing about 100 tonnes of material in the first year of operation.
Towards the end of the war, the British production of PE had grown to about 1500 tonnes
per annum. The strategic importance of PE was not lost on the United States. By 1943,
and with the support of the US Government, both Union Carbide Chemical and the
E.I. du Pont de Nemours and Company (DuPont) had licensed the ICI autoclave process and were starting up production facilities in North America. Within several years,
Union Carbide had also developed and commercialized a high pressure tubular process
for PE. By the early 1950s the Dow Chemical Company also licensed the ICI autoclave
process and joined the ranks of North American low density PE producers.
Although initially sourced from ethanol, the post-war years brought cheaper ethane
and ethylene sourced from large-scale petrochemical refining. This in turn brought the
price of PE down and helped to trigger subsequent growth of the polymer in industrial
and consumer markets worldwide.
As the ICI autoclave process matured, the use of oxygen as the free-radical reaction
initiator was supplanted by peroxyradical initiators designed with specific decomposition temperatures. The ICI autoclave design transitioned to multi-zoned, stirred reactors. By the early 1950s, it was also becoming clear that high pressure PE exhibited a
complex architecture with a wide variety of branching structures. Growing experience
with the new initiators, pressure controls and reaction zone temperatures allowed operators to control branching content, density, overall molecular weight, and molecular
weight distribution. By the mid-1950s, the high pressure product gained the designation LDPE (low density PE) to differentiate it from new higher density homopolymer
PE products starting to arrive on the market.
ICI continued to develop the autoclave process and experimented with a range
of polar co-monomers such as vinyl acetate, methyl acrylate, and n-butyl acrylate.
Although ethylene-vinyl acetate (EVA) copolymers did not demonstrate the best thermal properties, they offered one singular benefit. ICI (and others) realized that in the
high pressure autoclave environment, the ethylene and vinyl acetate monomers were
equally reactive to the growing macro free-radical chain end. As a result, a set ratio
of ethylene and comonomer in the feed stream produced an identical composition in
the polymer and left the same composition of monomers in the recycle stream with no
need for differential feed makeup. Furthermore, the comonomer incorporation was
randomly assimilated into the polymer chain. Hence, ICI introduced EVA products on
a limited scale to European markets in the late 1950s. Other manufacturers were also
involved in the development of ethylene vinyl acetate copolymers, with DuPont filing
a US patent in 1956 and introducing its commercial ELVAX EVA materials in 1960.
Today, commercial EVA grades contain from 2 w/w% to about 40 w/w% vinyl acetate. As the vinyl acetate content increases, the level of crystallinity decreases from
about 60% to around 10%. This yields EVA products with properties ranging from
those of typical LDPE to flexible thermoplastic elastomers. Clarity, flexibility, toughness and solvent solubility increase with increasing vinyl acetate content. Of particular
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Handbook of Industrial Polyethylene and Technology
note is the flexibility and impact resistance of EVA rubber grades at temperatures down
to −70 °C. Because they are copolymers, problems associated with plasticiser migration
observed in flexible PVC are not experienced. Being thermoplastic, EVA can be formed
by extrusion, injection molding, blow molding, calendaring, and rotational molding.
Cross-linking with peroxides can produce thermoset products. Applications are diverse
and include flexible shrink wrap, footwear soles, hot melt adhesives and sealant resins,
flexible toys, tubing, wire coatings, medical gloves, cross-linked microcellular foamed
tires, and sheet materials.
Although a number of manufacturers in addition to ICI worked with alternate
comonomers over the years, DuPont in particular continued to explore and successfully commercialize other highly polar ethylene copolymers. The polar comonomers
included methyl acrylate, ethyl acrylate, methyl methacrylate, acrylic acid, methacrylic
acid, and others. The high pressure autoclave process was well suited for this endeavour,
as many of these comonomers “poisoned” low pressure catalyst systems. In some cases,
reactor metallurgy had to be adjusted to resist the acidic nature of the comonomers.
Co – and terpolymers of ethylene containing these various species have found specialized markets such as adhesives, impact modifiers, and blend compatibilizing agents. In
some cases the acid copolymers can be saponified in a secondary process, creating very
tough Li, K and Na ionomers for specialty cheese and meat packaging and high toughness sporting goods applications.
Ethylene-n-butyl acrylate copolymers (EnBA) were introduced by companies such
as Exxon Chemical and Chevron Chemical in the late 1980s. Although the EnBA product consistency was initially uneven due to the different reactivity ratios of the comonomers, multiple feed strategies improved the clarity and performance of these reactor
products. Their primary advantage was a much higher thermal stability than conventional EVA products.
AT Plastics (formerly Canadian Industries Limited, the Canadian subsidiary of ICI)
and Mitsubishi Chemical both introduced moisture cross-linkable copolymers of ethylene and vinyl trimethoxysilane (VTMOS) for wire and cable and potable water tubing
applications. Because of the moisture sensitivity of the product, special moisture barrier
packaging was required.
ICI and the other earlier adopters of the high pressure autoclave and tubular processes demonstrated a pattern that would repeat several times in the evolution of the
industry. The commercialization process inevitably started with PE homopolymers and
advanced to the incorporation of comonomers to tailor and improve product performance. These innovations both encouraged and serviced the growing list of product
applications and customers.
1.4 The Advent of High Density Polyethylene
Although ICI continued to license the autoclave LDPE process globally, it was clear that
the autoclave process presented challenges to scale-up. Both the ICI stirred autoclave and
the Union Carbide tubular processes ran at high temperatures and pressures; the reactor vessel, piping and compressor systems were constructed of extremely heavy walled
materials and suffered from high installation and maintenance costs. Furthermore, the
An Industrial Chronology of Polyethylene 9
ethylene polymerization reaction is exothermic, and heat removal became a limiting
factor in reactor sizing. Explosive ethylene decompositions to hydrogen, carbon, and
methane (i.e., “decomps”) were not uncommon if reactor instabilities developed.
As the PE industry expanded in the post-war years, there was great interest in the
development of a lower pressure process that would overcome the high capital costs
and rate limitations associated with high pressure PE.
In 1953, partially as a result of the post-war boom in catalyst development, two separate research groups on different continents discovered “high density PE” (HDPE)
almost simultaneously. In the United States, a group led by John P. Hogan and Robert
L. Banks at Phillips Petroleum Company developed a chromium trioxide-based catalyst
on a silica support. At the Max Planck Institute in Mülheim, Germany, Karl Ziegler and
his team of researchers were working on aluminum alkyl activated titanium-based catalysts. Both groups demonstrated the feasibility of producing almost completely linear
HDPE at virtually the same time.
In 1951, the team at Phillips Petroleum in Bartlesville Oklahoma was initially investigating the reactions of chromium salts and propylene for the production of synthetic lubricating oils. The benchtop reactor became fouled with a waxy solid, later identified as atactic
polypropylene (PP). Banks and Hogan continued working with catalyst variants and at one
point were able to produce and isolate crystalline PP. Further experiments with ethylene
feedstock led to the production of a high molecular weight HDPE. To their credit Phillips
Petroleum, who were not active participants in the burgeoning polymer industry, recognized the potential for this catalyst and performed more experiments with an emphasis on
developing a suitable low pressure polymerization process to run the new catalyst.
The amorphous silica supported chromium trioxide catalyst that emerged from this
effort was capable of producing a high molecular weight PE product. The density of the
material was typically between about 0.956–0.965 g/cm3 and the high level of crystallinity indicated that the polymer backbone contained virtually no long-chain branching.
The Phillips reaction was run in inert low-boiling alkane diluents such as iso-pentane
or n-butane at modest 5 to 25 atmosphere pressures and temperatures in the 80–100 °C
range [8, 9]. The diluent media acted as a heat sink to dissipate effectively the heat of
reaction. The supported chromium catalyst provided a template for the growing PE
chain, which was essentially insoluble in the reaction media. Despite this, the slurried
particles had to be circulated at high speeds to prevent particles sticking to the reactor
walls and forming sheets of intractable polymer. Each catalyst particle produced a single large granule of HDPE resin that had to be filtered from the slurried reaction media,
dried and extruded into pellets for commercial sale. The chromium catalyst exhibited a
high activity and could be left in the finished HDPE resin.
Phillips Petroleum, however, had difficulties with their new slurry process and catalyst combination. Warehouses were reportedly filled with off-specification resins. In
1957, the introduction of a new toy, the hula hoop, allowed them to empty their reserves
and gave a much needed push to the developing polymer platform.
The polymerization process ultimately evolved into the vertical Phillips slurry loop
reactor system introduced in 1960. The slurry loop process offered low capital costs and
a vertical format that reduced the plant footprint significantly.
During the period from 1955 to 1975, the chromium-based catalyst became the preferred production route to HDPE. HDPE resins manufactured with the Phillips process
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Handbook of Industrial Polyethylene and Technology
exhibited a high molecular weight and a broad molecular weight distribution. The grades
were slightly higher in density than competitive products, indicating a higher level of
linearity. The high density produced a high modulus well suited for stiff, crush-resistant
containers. Phillips’ MARLEX HDPE resins were central to the development of rigid
commercial and industrial packaging such as the now ubiquitous HDPE milk jug.
The Phillips slurry loop catalyst system continued to develop and evolve over time.
A second generation catalyst was introduced in 1975 to produce lower molecular
weight resins. Third and fourth generation catalyst systems were developed in 1983 and
1990, respectively, allowing resin designers to broaden molecular weight distributions
for improved processability and melt elasticity [10]. Finally, in the late 1990s Phillips
Chemical (later to become known as Chevron Phillips Chemical) commercialized
metallocene catalyzed resins made on the slurry loop process. The Phillips slurry loop
process has been extensively licensed and can be found across the globe.
In the mid-1970s Union Carbide developed its own chromium-based catalysts using
different support media and precursor materials for use in the new UNIPOL gas phase
process for the manufacture of both HDPE and linear low density PE (LLDPE) products.
At about the same time that Hogan and Banks were investigating the reactions of
metal salts with propylene and ethylene at the Phillips Petroleum facility in Bartlesville
Oklahoma, Professor Karl Ziegler and a team of graduate students were researching triethylaluminum (TEAL) catalyzed ethylene polymerization reactions at the Max Planck
Institute in West Germany. Several fortuitous observations with Ni-TEAL complexes
led the research group to start systematically investigating the potential for transition
metal-TEAL complexes to catalyze the polymerization of ethylene. In October 1953,
one of Ziegler’s newest graduate students, Heinz Breil, succeeded in producing a linear
PE using a zirconium complex as a catalyst. Infrared spectroscopy demonstrated that
the polymer was in fact linear. One of the team’s senior researchers, Heinz Martin, subsequently was able to generate a linear HDPE under mild laboratory conditions using
a titanium tetrachloride-TEAL compound. Karl Ziegler patented the “composition of
matter” of this catalyst and of the new form of HDPE in 1953 [11] and left the task of
process development to his industrial licensees.
Both Farbwerke Hoechst AG of West Germany and Hercules Powder Company in
the US showed immediate interest in the new HDPE. Compared to LDPE, the Ziegler
HDPE offered improved hardness, higher modulus, and a higher melting point, making
it suitable for a range of new applications. In addition, the Ziegler products exhibited
slightly lower densities than the Phillips HDPEs and were ideal for applications requiring higher resistance to environmental stress cracking. Like the competing Phillips
process, the Ziegler catalyst promised lower capital plant cost and safer operational
parameters than the high pressure autoclave and tubular LDPE processes.
Hoechst was the first to take the new Ziegler catalyzed polymer to market with a
slurry process in 1955. Hercules Powder, who shared development with Hoechst,
opened the first US plant in 1957. Resin density was typically in the 0.940–0.950 g/cm3
range, indicating that the Ziegler product contained some limited levels of sparse longchain branching. As a result, these products were more flexible than the Marlex HDPEs.
The Ziegler products were narrower in molecular weight distribution, but could be tailored over a broad range of molecular weights. As a result, these resins found utility in
injection molding, blown film, and small part blow molding applications where lower
An Industrial Chronology of Polyethylene 11
melt viscosities were beneficial. Interestingly, the Phillips and Ziegler HDPE products
were somewhat unique and seldom competed with each other.
Meanwhile, Karl Ziegler and his research team were not idle. Shortly after patenting
the transition metal chemistry for PE polymerization, Ziegler moved on to the problem
of PP. Propylene does not respond to free-radical initiation to produce high molecular
weight polymers. Ziegler hoped to use the transition metal catalyst approach to produce isotactic PP, an isomeric form where the methyl side groups are all located on one
side of the polymer backbone which in turn creates a high degree of crystallinity. They
attacked the PP polymerization problem but met with only limited success.
Giulio Natta, a Professor of Chemistry at the Milan Polytechnic and a consultant
to the Italian chemical company Montecatini was aware of Ziegler’s success with the
polymerization of ethylene. A contractual agreement between Ziegler and Montecatini
made it possible for Natta to set up a joint research program. Early in 1954, the Natta
group successfully developed a TiCl4-based catalyst system that produced a mixture of
isotactic and atactic PPs. In the atactic form, the pendant methyl side groups in PP are
randomly distributed on either side of the polymer backbone. This effectively limits
crystallization so that the resin is sticky and soft and has no material strength. Natta
found that the atactic material could be readily extracted with solvent however, leaving
the desired isotactic form.
Karl Ziegler and Giulio Natta were awarded the Nobel Prize in Chemistry in 1963 for
their work on the polymerization of ethylene and propylene. Transition metal catalysts
of this type are used to this day and are commonly referred to as Ziegler-Natta (or ZN)
catalysts.
Concurrent with the Phillips Petroleum HDPE work and the discoveries of Ziegler
and Natta, Standard Oil of Indiana also introduced an HDPE polymerization system
based on a supported molybdenum oxide catalyst. Commercialization efforts fell short
however and the process never gained much acceptance.
As a footnote to this period, the commercialization of PP was ultimately hindered by
a conflict between Phillips Petroleum and Montecatini over the rights to the composition of matter for isotactic PP. The issue was not resolved until 1980 when the Federal
District Court of Delaware awarded the composition of matter patent to Phillips, based
on the original work of Hogan and Banks back in 1951–52. Ziegler’s colleague Heinz
Martin wrote a book that detailed the events, from the original bench work to the final
court resolution nearly 30 years later [12].
The first generation ZN catalysts used in PE manufacture were not supported and
were not soluble in the solvent or diluent systems used in the various commercial processes. The catalysts themselves contained multiple active sites, but the nature of the
catalytic complex was such that it was readily obscured by the growing polymer chain.
Activity levels were low and catalyst residues often had to be removed by filtration or by
passing over absorption beds.
1.5 Product and Process Proliferation
During this early development period, both Dow Chemical and DuPont were actively
engaged in experimentation with ZN catalyst systems that would function at high
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Handbook of Industrial Polyethylene and Technology
temperatures. Their aim was to develop new solution phase polymerization processes
for manufacturing HDPE in the US.
DuPont Canada, the Canadian subsidiary of DuPont, had different objectives however. At the time, it was operating a small LDPE blown film plant that was hampered
by tariffs on imported PE resins. They realized that with the new ZN catalysts, there
was an opportunity to design and build a single solution phase plant with the capability of producing both HDPE and lower density materials. The goal was to incorporate
short-chain branches on the PE backbone in order to disrupt the overall crystallinity and lower the product density as desired. Typical comonomer alkenes of interest
included 1-butene, 1-hexene and 1-octene. Pilot plant work suggested that 1-butene
was the comonomer of choice, based on economics, separation and distillation capabilities, and impact on product properties (1-octene was included some years later).
The new materials were dubbed “linear low density PEs” (LLDPE) to distinguish them
from high pressure LDPE products. The world’s first LLDPE plant went into production
near Sarnia, Ontario in early 1960 [13]. The new LLDPE product was named SCLAIR
after the nearby St. Clair River.
One of the unexpected benefits of the new high temperature solution process was
the short residence time in the reactor. This allowed for rapid grade changes that better served the small and fragmented Canadian PE markets. By 1970, the SCLAIR line
included products with densities ranging from 0.915 to 0.960 g/cm3 and with melt indices from less than 1 to 100 dg/min (190 °C, 2.16 kg), serving blown film, liquid packaging, and injection molding markets all from a single production plant. In about 1980,
DuPont Canada made the decision to license the SCLAIRTECH process globally.
Despite the versatility of the ZN catalyst variants in use around the world, PE producers worked diligently to improve their proprietary manufacturing positions. The
goal was to improve the catalytic activity of the ZN complexes as much as possible.
By 1970, a second generation of ZN catalysts on supported media began to appear.
Although numerous supports were evaluated, magnesium chloride systems largely won
out as the optimum support media and were widely used in the ensuing years. The supported catalyst effectively spread the ZN reactive sites over a large surface area, increasing the activity by several orders of magnitude over the older unsupported ZN systems.
This high activity negated the need for a catalyst removal step. These high activity ZN
catalyst systems provided the necessary tools for truly large-scale commercialization of
new ethylene-α-olefin copolymerization systems.
Dow Chemical was an early adopter of the new high activity ZN systems and
launched their ethylene-octene solution process under the trade name DOWLEX
resins in about 1970. The new ethylene-octene copolymers demonstrated outstanding
toughness, tear resistance, and clarity. Today, ethylene-octene LLDPE resins are available from a number of producers and are still considered to be the preferred choice for
high performance films for specialty consumer and industrial packaging.
In 1972, DSM opened a solution PE plant in the Netherlands focused on medium
density PE (MDPE) and HDPE products. In 1981, the plant capability was extended to
include octene-based LLDPE products sold under the STAMYLEX tradename. The
proprietary solution process was referred to as the COMPACT technology and has
since been licensed worldwide.
An Industrial Chronology of Polyethylene 13
In about 1968, Union Carbide learned how to use silica-supported Ziegler-Natta
and chromium-based catalysts in a gas-phase fluidized bed reactor to produce HDPEs.
In 1977, Union Carbide introduced high activity silica-supported ZN catalysts for the
production of 1-butene – (and later 1-hexene-) based LLDPEs. This Union Carbide gas
phase system, known as the UNIPOL process, was characterized by relatively low
capital and production costs. Innovations such as condensing mode, which appeared
in the mid-1980s and super-condensing mode in 1990 further increased the process
rates by removing heat from the reaction mass. The grade capability is determined by
the specific catalyst selected; as a result specific reactor platforms are usually rationalized to a limited product base. Grade transitions are difficult and lengthy, so long runs
of specific products are usually encouraged. The Union Carbide UNIPOL process has
been licensed worldwide, initially by Union Carbide alone, and later through a joint
venture with Exxon Chemical called Univation Technologies.
The gas phase process itself has also been modified and patented by other companies.
Examples include the Innovene technology developed by BP and Rhone-Poulenc
and now owned by Ineos and the SPHERILENE gas phase process now owned by
LyondellBasell Industries.
In the early 1970s, slurry technology was also taking advantage of the advances
in high activity supported ZN catalyst. Hoechst continued to advance their original
Hostalen stirred tank slurry process. Nissan introduced a stirred tank slurry process in Japan in 1976. This was subsequently acquired by Equistar and eventually
LyondellBasell Industries (LBI) and is the basis for the Matagorda TX plant built in
1983. LBI has also acquired the Hoechst technology and is licencing it worldwide.
Mitsui Chemicals introduced their own stirred tank CX slurry process for HDPE in
the mid 1970s.
The period between 1970 and 1990 witnessed an unprecedented increase in the use
of PE worldwide and the breadth of products and specialized processes expanded to
meet the varied needs of the customer. Some of these developments included:
Introduction of cross-linkable PEs – usually formulated with 1 to 2 wt%
of reactive peroxide such as dicumyl peroxide (e.g., DICUP dicumyl
peroxide). Used in medium and high voltage cable insulation, rotationally molded tanks, and foamed footwear applications (1970).
Free-radical extrusion grafting of vinyl trimethoxy silanes for moisture
curable wire and cable jacketing and PE cross-linked (PEX) tubing.
SIOPLAS process from Dow Corning/Owen Illinois for producing
improved environmental stress crack resistance (ESCR) and thermal
resistant (1970–73) materials. Also MONOSIL process (Maillefer/
UCC – 1975–85).
Introduction of homopolymers and ethylene copolymers of 4-methyl1-pentene, produced from ZN catalysts, for use in high temperature, high
clarity medical applications (a technology developed by ICI, but produced and marketed by Mitsui Petrochemical as TPX resins – 1975).
Electron beam and radiation modification of molecular weights for specialty EVA (1985).
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Handbook of Industrial Polyethylene and Technology
Gas phase very low density PE (VLDPE) resins with densities below
0.910 g/cm3 (UCC FLEXOMER resins – 1985).
Solution polymerization of high vinyl acetate-ethylene copolymers,
subsequently saponified to produce ethylene-vinyl alcohol (EVOH)
high oxygen barrier resins for multilayer food packaging. Pioneered by
Kuraray under the EVAL trade name (1975–1985).
BYNEL and FUSABOND maleic anhydride modified PEs for interlayer bonding in coextruded film, hot melt adhesives from DuPont,
PLEXAR maleic anhydride modified PEs for tie layer adhesives from
Chemplex/NorChem/Equistar (now LBI), PRIMACOR
ethyleneacrylic acid (EAA) resins and maleic anhydride modified PE tie layer resins from Dow Chemical, Reichold POLYBOND EAA resins. Bonding
to EVOH, polyamides, polyesters, aluminum foils etc. (1975–85).
LOTADER ethylene-ethyl acrylate-maleic anhydride (E-EA-MAH)
resins and ethylene-butyl acrylate-maleic anhydride (E-BA–MAH) terpolymers made by high pressure autoclave for hot melt adhesives, tie resins and coupling agents – CdF Chemie (1987).
Commercial introduction of DYNEEMA high-strength ultrahigh
molecular weight PE (UHMWPE) oriented fibers and strands gel spun
through a spinneret. Strength-to-weight ratios for these materials are
8 to 15 times higher than steel. DYNEEMA technology was invented
by Albert Pennings and Ron Koningsveld [14–16] in 1963, but made
commercially available by DSM in 1990. Honeywell (formerly Allied
Chemical) also produces UHMW PE gel-spun fibers and composites in
US under the trade name SPECTRA [17].
Throughout this period, many PE producers worked diligently to customize and tailor the molecular weight distributions of their products. This led many manufacturers
to develop multi-reactor processes operating in series (“cascade”) or parallel modes.
In these arrangements, each reactor is typically run under different conditions producing different molecular weights and comonomer compositions in each. The list of
manufacturers adopting multi-reactor strategies is almost endless but includes Hoechst
Hostalen and the Nissan/Equistar slurry processes (now operated by LyondellBasell
Industries ), Dow Chemical Company’s ELITE resins, NOVA Chemicals Advanced
SCLAIRTECH process, Mitsui Petrochemicals’ EVOLUE process (now Prime
Polymers) and the Borealis BORSTAR slurry loop/gas phase hybrid process.
1.6 Single-Site Catalysts Arrive
In early 1968, Clay Elston of DuPont Canada filed a US patent application describing
a process for the preparation of homogeneous copolymers of ethylene with 1-butene
and 1-octene. The catalyst system under investigation was a vanadium-based transition metal complex. In the disclosure, homogeneous copolymers were defined “as those
in which not only is the comonomer randomly distributed within a given molecule,
but all of the copolymer molecules have the same ethylene/comonomer ratio” (i.e.,
An Industrial Chronology of Polyethylene 15
composition). Furthermore the molecular weight distributions were described as being
extremely narrow compared to equivalent products prepared with conventional ZN
catalysts. The patent was issued in early 1972 [18]. While Elston’s vanadium catalyst
lacked sufficient activity for commercialization, this was the first patent disclosing the
discovery of single-site catalyzed PEs—products with exceptionally narrow molecular
weight distributions and uniform comonomer incorporation.
Several years later in 1976, Walter Kaminsky and his supervisor Hansjorg Sinn,
working at the University of Hamburg, were examining reactions of biscyclopentadienyl-titanium dimethyl and biscyclopentadienyl-zirconium dimethyl with trimethylaluminum (TMA). While preparing NMR samples, they noted a new peak that they had
not seen previously [19]. When scaled-up and exposed to ethylene, the new complex
consumed the reactive monomer at much higher rates than expected. They systematically explored the possible contaminants and finally traced this activity increase to the
inadvertent exposure of the reactive TMA to water vapour. The research team found
that the polymerization rate with ethylene was optimized when the water to TMA molar
ratio was 1:1. The new activation species was investigated at length and finally christened “methylaluminoxane” or MAO. By 1979, Kaminsky and coworkers had examined
MAO-activated zirconium and titanium metallocene complexes with both propylene
and ethylene monomers and a patent was subsequently filed [20].
The metallocene MAO catalysts were 10 to 100 times more active than the supported
ZN catalysts of the time. They were soluble in solution systems where they worked as
a molecular catalyst with a single active polymerisation site. Furthermore Kaminsky
demonstrated that these active species could be adsorbed on inorganic supports for use
as a drop-in catalyst for gas phase and slurry polymerization systems.
The race was clearly on. In mid-1990, Exxon Chemical opened the first commercial
production facility for metallocene-catalyzed PE at its Baton Rouge LA plant site. The
process used a modified high pressure autoclave technology developed by Mitsubishi
and licenced to Exxon for use with their new EXXPOL metallocenes [21]. The
Mitsubishi process showed that high pressure autoclaves and tubular reactors could
be modified to use soluble, fully homogeneous metallocene catalysts in much the same
way that solution polymerization platforms could. Gas phase and slurry loop reactors
would require supported catalyst technologies however [22, 23]. Exxon Chemical was
the first company to introduce metallocene-based ethylene-butene and ethylene-hexene copolymers marketed under the EXACT trade name in 1991. The initial EXACT
“plastomer” grade lineup included provisional products ranging between 0.870 and
0.920 g/cm3 density [24]. Early experimental offerings also included ethylene-butenehexene terpolymers, but these appear to have gained little commercial interest.
In 1995, Exxon announced a major expansion of its EXACT plastomer capacity by
retrofitting part of its Paxon Polymer Co. plant in Baton Rouge [25]. In the same year,
Exxon formed a joint venture with DSM in Geleen Netherlands dubbed DEX Polymers
to utilize one of their COMPACT solution reactors to produce EXACT ethyleneoctene plastomers. In 2013, the DEX facility was acquired by Borealis AG, who now
produces the octene plastomers under the QUEO trade name.
In 1987 Exxon also began a collaborative research program with Mitsui Petrochemical
to develop a metallocene-catalyzed gas phase process. It is not clear how long the relationship lasted, but by the spring of 1995, Exxon Chemical had received a significant
16
Handbook of Industrial Polyethylene and Technology
patent on gas phase operation and was running an EXXPOL metallocene catalyst in
super-condensed mode [26, 27]. Exxon began to sample gas phase metallocene-catalyzed ethylene-hexene resins to the marketplace later that year. Unlike the EXACT
products, these new EXCEED grades were aimed squarely at conventional LLDPE
blown and cast film markets. The product line became fully commercial in about 1997–
98. EXCEED resins exhibit a narrow molecular weight distribution and are more challenging to process than conventional ZN grades. LDPE is often blended in to improve
bubble stability in blown film applications. Because the metallocene catalyst for the gas
phase process is supported, it does not act strictly as a single-site molecular catalyst; the
comonomer incorporation in EXCEED resins is not entirely uniform and this provides
some interesting benefits in film sealing performance.
In 2008, Exxon, by now ExxonMobil Chemical, introduced ENABLE resins, a
new gas phase metallocene/single-site product which incorporated low levels of longchain branching. ENABLE resin grades are easier to process because of the long-chain
branch content, and to some extent this also allows for the production of lower melt
index grades. ExxonMobil claims that these materials can successfully replace LDPE
in EXCEED/LDPE film blends. The new EXCEED/ENABLE blends and co-ex film
structures offer superior down gauging capabilities, improved stiffness, and improved
optics [28].
In September 1993, The Dow Chemical Company launched its own line of singlesite ethylene-octene plastomers. These were manufactured in a solution process using
Dow’s new INSITE catalyst technology. The technology was widely promoted as a constrained geometry catalyst (CGC) that incorporated octene readily but was also capable
of incorporating vinyl-terminated macromolecules as long-chain branches [29, 30].
The initial polyolefin plastomer (POP) product line was trademarked AFFINITY
resins and was promoted for a wide range of blown and cast film, “breathable” fresh
cut produce packaging, sealant, and lamination applications. AFFINITY resin product
densities initially ranged from about 0.870 to 0.920 g/cm. The product line was characterized by narrow molecular weight distributions and uniform octene comonomer
incorporation as a function of molecular weight. Individual grades contained varying
degrees of long-chain branching, as evidenced by their “I10/I2“ values (a ratio of melt
indices carried out with 10 kg and 2.16 kg piston loadings), shear thinning behaviour
and general processability in extrusion processes.
In February 1994, the company followed with a line of lower density ENGAGE polyolefin elastomer resins (POE) for use in profile extrusion, shoe soles, wire and cable
jacketing and insulation, gaskets, polymer modification, and rubber replacement.
Despite the market positioning of these two product lines, many of the early ENGAGE
resins appeared to be rebranded AFFINITY resin grades.
In 1995, presumably to gain a broader access to rubber and elastomer replacement applications, Dow Chemical formed a joint venture (JV) with DuPont known as
DuPont Dow Elastomers. The new company marketed a number of Dow solution products of that included ENGAGE and NORDEL EPDM and EPR elastomers [31]. With
the active promotion of the new ENGAGE and AFFINITY resin product lines, Dow
Chemical reportedly converted 250 million pounds per year of ZN-based DOWLEX
solution process capacity in Freeport, Texas to produce the new single-site catalyzed
polymers. A 125 million pound per year plant in Tarragona, Spain was also converted.
An Industrial Chronology of Polyethylene 17
In early 2005, the JV with DuPont was dissolved and Dow absorbed the ENGAGE
product line back into its portfolio. By this time, the ENGAGE resin product offering
had diverged significantly from AFFINITY resin products and now included ethylenebutene copolymers and a number of specialty grades. Product densities now ranged
from 0.857 to 0.915 g/cm and melt indices between about 0.3 and 30 dg/min (190 °C,
2.16 kg).
Over the two decades since their commercial introduction, Dow has broadened
the scope of the original octene-based AFFINITY product line. This includes several
grades for extrusion coating, four hot melt adhesive grades initially developed and copatented with adhesives manufacturer H.B. Fuller, and a rotational molding grade that
appears to have met with only limited success. There is also evidence of several distinctly bimodal AFFINITY POP products over the years.
In 1997, Dow Chemical launched the ELITE product line, another new line of
enhanced PE resins made using the INSITE catalyst technology. The products were promoted as high performance resins for commodity blown and cast film markets. They
offered an improved balance of stiffness and toughness with improved sealability. Over
time, it became clear that these products were the result of a dual reactor mixed catalyst strategy—a homogeneous CGC single-site resin component made in one reactor
and a Ziegler-Natta heterogeneous resin component produced in a second reactor [32].
ELITE products have expanded to include several higher density “barrier” food packaging grades and lower density products for stretch hooder and sealant applications. A
new line of Advanced Technology ELITE resins was also introduced in early 2012.
Dow Chemical also employed its single-site constrained geometry catalyst to develop
INDEX ethylene-styrene copolymer elastomers in 1999. Using INSITE Technology
and an optimized solution process, Dow was able to vary the ratio of ethylene and styrene monomers to provide a broad range of composition and performance attributes.
On the high-ethylene content end of the spectrum (E series), INDEX interpolymers
appeared to provide increased flexibility, abrasion resistance, toughness, paintability
and creep resistance. On the high-styrene content end (S series), INDEX interpolymers
offer increased deadfold, paintability and sound management capabilities [33]. Despite
apparent commercial interest, the INDEX product line was withdrawn from the market
in about 2002.
In 2006, Dow Chemical launched another innovative new line of olefin-based elastomers. INFUSE Olefin Block Copolymers (OBC) are ethylene-octene products
claimed to offer superior property performance compared with competitive elastomeric materials. The chain-shuttling technology uses two distinct single-site catalysts
and diethylzinc as the shuttling agent in one or more continuous solution reactors.
One single-site catalyst produces hard ethylene-based blocks while the second hybrid
single-site catalyst produces soft octene-rich blocks. Blocks form on the two catalysts,
detach from them and park on diethylzinc particles before reattaching to new monomer blocks forming on the catalysts. The length of block segments is controlled by the
ratio of diethylzinc to ethylene monomer, with a higher ratio making finer blocks, while
the overall hard-to-soft ratio is controlled by the relative amounts of the two catalysts
[34]. According to Dow, the unique ethylene-octene block architecture results in resins
with high flexibility at varying temperatures, faster setup time, greater elastic recovery,
strong compression set and improved abrasion resistance. Targeted applications include
18
Handbook of Industrial Polyethylene and Technology
elastic films, flexible molded goods, adhesives, and foams. Dow claims that the INFUSE
products have been particularly successful in replacing EVA and styrene block copolymers in athletic footwear applications, where they offer reduced weight and improved
softness and resiliency [35].
Dow unveiled the INTUNE resin line, a PP-based OBC in October 2013. The new
material is made using the same proprietary chain shuttling catalyst technology used
for the INFUSE product line and is claimed to “combine PE and PP to give previously unattainable performance” [35]. INTUNE resin is expected to find a home in
packaging, consumer durables, transportation, and other markets. The new material
can combine PE, polyolefin elastomers, ethylene vinyl alcohol, or polyamides with PP
while minimizing trade-off properties. The product also acts as a compatibilizer for
PE and PP in compounding and recycling applications [36]. The material is not commercially accessible at this time, but will be apparently be available in injection molding and extrusion grades and is not intended to replace the firm’s INFUSE olefin block
copolymer.
The slurry loop process poses some unique challenges for metallocene catalyst
development, since the catalyst system must be engineered to remain rigorously
insoluble in the light hydrocarbon slurry medium. Phillips Chemical (now Chevron
Phillips Chemical) was able to resolve this issue and developed a suitable supported
metallocene catalyst in the mid-1990s. In 1998, it commercialized its line of metallocene-catalyzed LLDPE products under the trade name mPACT . In the years that
followed, enhancements in both the slurry loop process and metallocene activation
strategies resulted in further improvements in operability, product quality and cost
structure [37]. Current mPact ethylene-hexene resin offerings have product densities ranging from 0.914 to 0.933 g/cm with melt indices in the 0.9 to 1.4 dg/min
range and are targeted squarely at performance blown film markets. The mPACT
metallocene grades exhibit narrow molecular weight distributions, but the comonomer distribution is non-uniform across the molecular weight distribution. Like the
ExxonMobil EXCEED gas phase metallocene resins, this provides distinct benefits
in film sealing performance. All of the mPACT products display exceptional film
clarity and gloss.
Major western companies including Exxon and Dow have also formed strategic joint
ventures with Japanese manufacturing companies such as Mitsui, Asahi and Mitsubishi
for the development of PE technologies and in some cases, marketing of products. A
number of Japanese polymer producers have also licensed or developed their own production platforms. With the advent of metallocene technologies, the polyolefin industry in Japan has undergone a number of mergers and joint ventures. Mitsui Chemical
and Sumitomo Chemical are among the largest independent producers of PE products
in Japan at the moment [38].
TAFMER products represent a line of polymeric modifiers that was first introduced by Mitsui Petrochemical Industries in 1971. These unique ethylene-based elastomeric products were widely promoted in Asian and North American markets for
use as impact modifiers, adhesion promoters, sealant resins, film, and sheet lamination
components and for use in multi-layer coextruded structures [39–41]. TAFMER resins
were originally produced with vanadium-based Ziegler-Natta catalysts. Over the next
two decades, the TAFMER product line grew to include:
An Industrial Chronology of Polyethylene 19
TAFMER A – low crystallinity ethylene-butene copolymers with densities in the 0.880 to 0.890 g/cm3 range.
TAFMER P – “zero” crystallinity ethylene-propylene copolymers with
densities in the 0.880 to 0.860 g/cm range.
TAFMER XR – low crystallinity propylene-butene copolymers with densities in the 0.88 to 0.89 g/cm range.
In 1997, Mitsui Chemicals announced an advance in metallocene chemistry, and
commissioned the first Japanese metallocene-based manufacturing facility dedicated
specifically to elastomers [38]. The metallocene catalysis narrowed the molecular weight
and comonomer distributions of the products, improving product consistency, clarity, and blend miscibility. Mitsui immediately began to convert its Japanese TAFMER
A and P products from ZN catalysts to the newer metallocene catalyst technology. An
overseas facility with two TAFMER resin production lines was subsequently opened in
Singapore in 2003. The Singapore facility was further expanded in 2010. Plant output
was believed to be in excess 200 kilotons per annum (ktpa) after this expansion was
completed. Yet another expansion was announced in 2012, with an aim to increase
output by another 8 ktpa.
Mitsui has also continued to expand the product family with the addition of TAFMER
DF, H, XM, PN BL and M series resins. DF and H are ethylene-butene and (most likely)
ethylene-hexene-based respectively. The BL series are butene-based. XM appears to
have replaced the older XR products. TAFMER PN resin is a propylene-based material
with controlled morphology and TAFMER M resin is an acid modified polyolefin elastomer. TAFMER products are believed to be manufactured on a proprietary solution
polymerization platform.
In 1996, Mitsui Petrochemical Industries, Ltd. and Sumitomo Chemical Industries
Ltd. established Evolue Japan Co. to produce metallocene-catalyzed LLDPE resins in
Japan. Both the dual cascade gas phase reactor technology and the subsequent bimodal
resins are referred to as EVOLUE resins.
Mitsui Chemicals and Idemitsu Kosan established Prime Polymer in 2005 as a
means to better manage their polyolefins businesses. Mitsui Chemicals currently owns
65 percent of the Prime joint venture, with Idemitsu owning the remaining 35 percent.
Mitsui retains ownership of the EVOLUE process technology and is solely responsible
for process development, licensing, and the associated metallocene catalyst business.
Prime Polymer reportedly takes 80 percent of the product from the EVOLUE resin
plant in Chiba Japan, with the balance allocated to Sumitomo. Prime Polymer continues to develop domestic and export markets for Evolue resins, focusing on high-value
applications.
Purportedly, Mitsui’s primary objective in developing Evolue was to produce LLDPE
with good processability, excellent mechanical properties (including impact and tear
strength) and good optical properties. Mitsui believed that a bimodal approach using
its metallocene catalyst would achieve these properties. They opted for two in-series
cascade gas-phase reactors as the most efficient means of producing high volume
fractional melt index and high flow commodity resins without sacrificing production
rate. The Evolue dual reactor cascade gas-phase process was co-developed with Exxon
Chemical in the early 1990s.
20
Handbook of Industrial Polyethylene and Technology
Since Mitsui Chemicals already had a significant commercial and technical position
in PE with its elastomer and LLDPE solution, HDPE slurry, and LDPE autoclave processes, it dedicated the EVOLUE resin plant to the manufacture of metallocene-based
LLDPE resins. The process uses 1-hexene as the comonomer and currently produces
resins in the 0.903 to 0.946 g/cm density range. Melt indices between 0.8 to 3.8 dg/min
are offered in the current grade slate, although grades as high as 10 dg/min have been
manufactured in the past. In addition to bimodal LLDPE grades, the process is also
capable of producing unimodal LLDPE resins for applications where processability is
not critical or narrow molecular weight distributions are preferred (e.g., cast stretch
films, lamination films and rotational molding products).
Mitsui Chemicals claims that its EVOLUE resins have advantages over single reactor metallocene-based grades, which include very high impact, machine direction (MD)
tear and environmental stress crack resistance (ESCR), lower seal initiation temperature
(SIT), and low haze. It also claims that because of the good melt strength, the processability is better than long-chain branched metallocene/single-site LLDPE. Mitsui/Prime has
also claimed that EVOLUE resins compete successfully against octene-based solution
single-site resins and it has positioned it in Asian markets as a premium-priced product.
The original plant in Chiba has been debottlenecked and expanded several times and
is now believed to have a nameplate capacity of about 250 ktpa. A new 300 ktpa plant
was built in Singapore with commercial operation starting in 2016.
In 2001, NOVA Chemicals completed the construction of its new Advanced
SCLAIRTECH dual reactor solution platform at its Joffre site in Western Canada. The
technology was a fundamental extension of the original DuPont Canada SCLAIRTECH
technology acquired some years before. The initial phase of the operation used a new
high-performance supported ZN catalyst produced in a continuous in-line manufacturing process linked directly to the reactor. The Advanced SCLAIRTECH ZN ethylene-octene product line was marketed under the SCLAIR trade name.
Several years later, NOVA Chemicals commercialized a proprietary single-site catalyst for use on its dual reactor AST solution platform. The multimodal SURPASS
product line currently ranges from 0.914 to 0.967 g/cm in density. The process and
catalyst are extremely versatile and are capable of producing resins over a range from
0.6 to 150 dg/min melt index. The dual-reactor system allows for the precise tailoring of molecular weight distributions extending from very narrow to extremely broad
and bimodal, with complete separation of the two reactor modes if desired. Octene
comonomer distribution can also be controlled, producing a diverse range of resin
architectures. Reactor residence time is of the order of minutes and grade transitions
are extremely fast. The Advanced SCLAIRTECH platform now operates with both the
Ziegler-Natta and the proprietary SSC catalyst systems.
The Advanced SCLAIRTECH plant has been debottlenecked several times over
the course of its first decade and the production capacity is now about 550 ktpa. The
Joffre plant is currently the only one of its kind worldwide, but NOVA Chemicals has
announced that it is evaluating options for a second Advanced SCLAIRTECH facility
in a location yet to be determined.
INEOS is one of the latest entrants in the metallocene LLDPE field, starting up in late
2007. Its ELTEX products are reportedly made using a proprietary constrained geometry-type supported single-site catalyst in the INNOVENE gas phase process. Other
An Industrial Chronology of Polyethylene 21
new players include Total Petrochemicals with its LUMICENE and LUMICENE
SUPERTUF second generation HDPE and MDPE film resins. South Korea’s SK
Innovation has also announced its entry into the metallocene-catalyzed ethyleneoctene LLDPE and VLDPE markets with its SMART resins made using NEXLENE
dual reactor solution technology.
1.7 The Future of LDPE
In 1983, on the 50th anniversary of the Gibson and Fawcett experiment, ICI presented
the original high pressure reactor used by the team to the Science Museum of London.
The museum collection also includes an extruded strand of PE made at the Wallerscote
Works, Winnington on December 12, 1938. The sample is labeled as “part of the 1016th
kg of PE” to be produced.
Ironically, five years after this momentous anniversary ICI, the company that started
the PE revolution, had divested or closed virtually all of its PE manufacturing facilities
world-wide, opting to focus instead on its paints and coatings divisions and its growing
pharmaceutical business.
Despite this apparent setback, many of the early ICI autoclave polymerization facilities built in the mid-1950s continue to operate safely and profitably. The autoclave process is still hampered by high capital costs and scale-up challenges and the last new
autoclave reactor installation (rated at about 50 ktpa) was constructed in the mid1990s. Tubular high pressure technology has become the preferred source of new LDPE
capacity and lines with 350 ktpa capacity are not uncommon. The long-chain branching
structure of tubular LDPE grades is subtly different from that of the autoclave LDPE,
and autoclave products are still preferred for applications such as extrusion coating
where neck-in performance is critical.
Almost from the beginning, the stated goal of most new PE technologies was to
render the original LDPE obsolete. In spite of the high costs associated with high pressure LDPE processes and the “inferior” physical properties of LDPE resins, the product
continues to occupy an important spot in the portfolio of PE resins available to current
customers. It offers high melt strength and bubble stability and is a common blend
ingredient in many monolayer and multilayer film formulations. LDPE markets are
highly competitive and the material is far from obsolete.
Reginald Gibson and Eric Fawcett would be proud. The rumors of LDPE’s passing
have been greatly exaggerated.
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Handbook of Industrial Polyethylene and Technology
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2
Catalysts for the Manufacture of Polyethylene
Yury V. Kissin
Rutgers University, New Brunswick, New Jersey, USA
Contents
2.1 Introduction ................................................................................................................................. 26
2.2 Synthesis of Low Density Polyethylene .................................................................................... 27
2.2.1 Peroxide Initiators ........................................................................................................... 28
2.2.2 Chemistry of Radical Polymerization Reactions ........................................................ 28
2.2.3 Types and Degree of Branching in Low Density Polyethylene Resins ..................... 32
2.3 Catalytic Synthesis of Polyethylene Resins .............................................................................. 32
2.3.1 Commercial Technologies of PE Manufacture ............................................................ 32
2.3.2 Chromium-Based Catalysts ........................................................................................... 33
2.3.2.1 Chromium Oxide Catalysts (Phillips Catalysts) ........................................... 33
2.3.2.2 Organochromium Catalysts ............................................................................ 35
2.3.3 Titanium-Based Ziegler-Natta Catalysts ...................................................................... 36
2.3.3.1 General Features of Commercial Ziegler-Natta Catalysts ........................... 36
2.3.3.2 Catalysts Produced from Soluble MgCl2 Complexes ................................... 38
2.3.3.3 Catalysts Produced by Synthesis of MgCl2 .................................................... 40
2.3.3.4 Ziegler-Natta Catalysts Used in Solution Polymerization Processes ......... 40
2.3.4 Metallocene Catalysts ..................................................................................................... 41
2.3.4.1 Metallocene Complexes and Cocatalysts for Them ..................................... 42
2.3.4.2 Supported Metallocene Catalysts.................................................................... 44
2.3.5 Post-Metallocene Ethylene Polymerization Catalysts ................................................ 45
2.3.6 Binary Transition Metal Catalysts ................................................................................. 46
2.3.6.1 Binary Ziegler-Natta/Metallocene Catalysts ................................................. 46
2.3.6.2 Binary Metallocene and Post-Metallocene Catalysts ................................... 47
2.4 Chemistry of Catalytic Polymerization Reactions .................................................................. 47
2.5 Uniformity of Active Centers .................................................................................................... 51
2.5.1 Uniformity of Active Centers with Respect to Molecular Weight of Polymers ...... 51
2.5.2 Uniformity of Active Centers with Respect to Copolymerization Ability ............... 51
References .............................................................................................................................................. 54
Corresponding author: ykissin@rci.rutgers.edu
Mark A. Spalding and Ananda M. Chatterjee (eds.) Handbook of Industrial Polyethylene and Technology, (25–60)
© 2018 Scrivener Publishing LLC
25
26 Handbook of Industrial Polyethylene and Technology
Abstract
Polyethylene (PE) resins of various types are produced in industry in two principally different
types of polymerization processes. The processes of the first type are radical polymerization
reactions of ethylene under high pressure. These reactions are initiated by peroxide compounds.
The processes of the second type are catalytic polymerization reactions of ethylene and its copolymerization reactions with -olefins. The catalysts employed in these reactions are based on
derivatives of three transition atoms: titanium, chromium or zirconium. This chapter describes
the principal radical initiators and the transition-atom catalysts which are used for polymerization of ethylene.
Keywords: Catalysts for PE manufacture, catalysts, Ziegler-Natta catalysts, chromium oxide
catalysts, metallocene, polymerization initiators
2.1 Introduction
All semicrystalline polymers derived mostly from ethylene and used as commodity
plastics are called PE resins. These materials are produced in radical polymerization
reactions under high pressure and in catalytic polymerization reactions. Some PE resins consist of strictly linear polymer chains; their chemical formula is (CH2 CH2)n ,
where n is usually a very large number ranging from ~1,000 to ~100,000. Two such
PE resins are produced commercially, high density PE (HDPE) with n from ~1,000 to
10,000 and linear PE of ultrahigh molecular weight (UHMWPE) with n above 30,000.
Other PE resins contain branches in their molecular chains. The structure of all such
resins can be schematically represented by a general formula:
(CH2 CH2)x CH2 CH(Branch) (CH2 CH2)y CH2 CH(Branch)
(CH2 CH2)z CH2 CH(Branch)
The CH2 CH2– units in such polymer chains come from ethylene and the CH2
CH(Branch)– units can come from two different sources. They are either formed spontaneously in the course of radical polymerization reactions or they are introduced
deliberately by copolymerizing ethylene with -olefins over transition-metal polymerization catalysts.
When ethylene is polymerized at high pressure via the radical mechanism, branches
of many different types are formed in parallel due to peculiarities of the radical reactions,
as described in Section 2.1.2. These branches are linear or branched alkyl groups. Their
length varies widely within each polymer molecule. Two types of such branches are usually distinguished, short-chain branches, from the methyl group to the isooctyl group,
and long-chain branches, which can contain up to several thousand carbon atoms.
When ethylene is copolymerized with -olefin molecules CH2=CH R over transition metal catalysts, all the branches are the same; they are alkyl substituents R in the
-olefin units CH2 CHR . If the -olefin is 1-butene, the R branch is the ethyl group
CH2 CH3; if the -olefin is 1-hexene, the R is the butyl group (CH2)3 CH3; if the -olefin is 1-octene, the R is the hexyl group (CH2)5 CH3; etc. In principle, all such PE
materials should contain only short-chain R branches derived from the -olefin units.
However, polymerization reactions utilizing some metallocene catalysts and chromium
Catalysts for the Manufacture of Polyethylene 27
oxide catalysts can also introduce long-chain branches in the PE chains. Their formation mechanism is described in Section 2.1.4.
Some metallocene catalysts can copolymerize ethylene with cycloolefins, such as
cyclopentene, cyclooctene, norbornene, etc. These PE resins can also be viewed as
branched; the branches in this case are either small cycles containing from 5 to 10
carbon atoms, or two fused cycles. These copolymers are also produced commercially;
they form a special class of PE resins called cycloolefin copolymers (COC).
In polymer science, the fraction of an -olefin in an ethylene/ -olefin copolymer
is represented as the molar content of the -olefin, CMcopol (mol%). The content of
an -olefin in an ethylene/ -olefin copolymer is usually measured using 13C nuclear
magnetic resonance (NMR) or infrared (IR) spectroscopic methods. The content of
-olefins in commercially manufactured ethylene/ -olefin copolymers varies in a
wide range, from less than 0.5 to over 20 mol%. Depending on the content of -olefin,
these products are called medium density PE resins (MDPE, copolymers with CMcopol
from 1 to 2 mol%), linear low density PE resins (LLDPE, copolymers with CMcopol from
2.5 to 3.5 mol%), or very low density PE resins (VLDPE, copolymers with CMcopol over
5 mol%). The VLDPE resin group is further divided into two subgroups, PE plastomers
with a crystallinity degree of 10 to 20% and completely amorphous ethylene elastomers.
In industry, the fraction of an -olefin in a PE resin is represented by a parameter
called “the branching degree.” It is defined as the number of branches per 1000 carbon
atoms, Branch/1000C. Because all such branches usually end with a methyl group, the
branching degree is often represented by the symbol CH3/1000C. The ratio between
the content of an -olefin in an ethylene/ -olefin copolymer, CMcopol (mol%), and the
branching degree is given by:
CH3 /1000 C = CMcopol/(0.2 + 0.001 · k · CMcopol)
(2.1)
Here k is equal to 2 for ethylene/1-butene copolymers, 4 for ethylene/1-hexene copolymers, and 6 for ethylene/1-octene copolymers. In the case of PE resins prepared in
radical polymerization reactions (i.e., LDPE resins), the combined number of all shortchain branches is also expressed as a CH3/1000C value.
2.2 Synthesis of Low Density Polyethylene
Low density polyethylene resins are PE polymers which are synthesized in radical
polymerization reactions of ethylene at high pressure. The first materials of this type
were produced in 1933 in Great Britain. At the present time, over 40% of all the PE
resins are produced in polymerization reactions of this type. Production of LDPE is
carried out at a temperature ranging from 100 to 300 °C and under very high pressure
ranging from 1200 to 3000 atmospheres (120 to 300 MPa).
Two continuous high-pressure technologies are currently used for the commercial
manufacture of LDPE resins, polymerization in a tubular reactor and polymerization
in an autoclave.
Tubular reactors are made of a folded high-strength steel tube with an inside diameter of 2.5 to 6.5 cm and a very large length, from 500 to 1500 meters. Compressed
28 Handbook of Industrial Polyethylene and Technology
ethylene is continuously injected into the tube and a mixture of liquefied ethylene and
molten PE is continuously discharged from it. The ethylene polymerization reaction is
exothermic and the heat of the reaction must be removed by external cooling. Radical
initiators are injected at several points along the tube length and, as a result, the temperature along the reactor varies in a significant range. The PE/ethylene mixture flows
through the tube at a very high linear velocity of about 20 m/s. Although the residence
time is short, from 30 to 70 s, it is sufficient to polymerize ethylene to a 20 to 30%
conversion. The hot reaction mixture exiting the reactor is immediately separated into
ethylene, which is recycled, and molten PE, which is extruded and pelletized.
The second technology of LDPE synthesis employs vertically positioned tall highpressure reactors with a volume of up to 2 m3. A typical reactor is equipped with a
stirrer and separated into several zones. The polymerization process is also continuous.
Ethylene and a radical initiator are injected into the reactor and a mixture of ethylene
and molten PE (at a conversion from 15 to 20%) is continuously discharged from it. The
average residence time in this process is also short, ranging from 20 to 80 s. In spite of
the modest reactor volume in both high-pressure processes, the production capacity is
very high, up to 20 to 25 ton/h, because of the short residence time.
2.2.1 Peroxide Initiators
Radical polymerization reactions of ethylene are initiated by chemical compounds
called organic peroxides. The principal feature of all these compounds is the peroxide bond between two organic moieties, R–O–O–R . The peroxide bond is unstable at
increased temperatures, and it readily dissociates with the formation of two radicals,
R–O and R –O Organic peroxides decompose at different rates depending on the
type of R and R Each peroxide is characterized by a special kinetic parameter which
describes its effectiveness as an initiator in an ethylene polymerization reaction. This
parameter is the temperature at which the half-life of the peroxide, τ0.5 (the time after
which the concentration of a peroxide is reduced by 50%) is equal to one hour. Highly
unstable peroxides used for LDPE synthesis have τ0.5 = 1 h at temperatures below
100 °C, peroxides of average stability have τ0.5 = 1 h at temperatures between 100 and
130 °C, and relatively stable peroxides have τ0.5 = 1 h at temperatures above 130 °C.
The initiators for each zone in a polymerization reactor are selected among numerous
commercially available peroxide compounds; the choice depends on the temperature
of the PE-ethylene mixture in the zone, the ethylene concentration in it, and a desirable
conversion degree in the zone.
Table 2.1 gives several examples of peroxide initiators used in the industry for the
synthesis of LDPE and the respective temperatures of τ0.5 = 1 h.
2.2.2 Chemistry of Radical Polymerization Reactions
A radical polymerization reaction of ethylene is a chain reaction. The reaction sequence,
as in all chemical chain reactions, contains the steps of chain initiation, chain propagation, and chain termination. The reaction scheme given below also contains several
steps typical for chain polymerization reaction in general, such as isomerization reactions of a polymer chain and chain transfer reactions.
Catalysts for the Manufacture of Polyethylene 29
Table 2.1 Examples of peroxide initiators for the synthesis of LDPE [1].
Temperature, C, for τ0.5 = 1 h
Chemical name
tert-butyl peroxyneodecanoate
64
tert-amyl peroxypivalate
72
tert-butyl peroxy-2-ethylhexanoate
91
tert-butyl peroxydiethylacetate
93
1,1-di(tertamylperoxy)cyclohexane
106
tert-butyl peroxyacetate
119
tert-butyl peroxybenzoate
122
2,5-dimethyl-2,5-di(tert-butylperoxy)hexane
134
tert-butyl cumyl peroxide
136
di-tert-butyl peroxide
141
1,1,3,3-tetramethylbutyl hydroperoxide
160
tert-amylhydroperoxide
190
Chain initiation reaction (decomposition of a peroxide molecule with the formation
of R–O radicals; the rate-controlling reaction step):
R–O–O–R
R–O + R –O
(2.2)
Chain propagation reactions (growth of a polymer chain)
i) the addition of the first ethylene molecule to the R–O radical:
R–O + CH2=CH2
R–O–CH2–CH2
(2.3)
Growth of a polymer chain:
R–O–CH2–CH2 + n CH2=CH2
R–O–(CH2–CH2)n–CH2–CH2
(2.4)
Isomerization reactions of polymer chains
Because stability of any primary radical R–CH2 is lower than stability of a secondary radical R1–CH –R2, growing chains in radical polymerization reactions of ethylene
occasionally isomerize.
Formation of a methyl branch:
R–O–(CH2–CH2)n–CH2–CH2
R–O–(CH2–CH2)n–CH –CH3
R–O–(CH2–CH2)n–CH –CH3 + m CH2=CH2
(2.5)
30 Handbook of Industrial Polyethylene and Technology
R–O–(CH2–CH2)n–CH–(CH2–CH2)m 1–CH2–CH2
(2.6)
CH3
Formation of an ethyl branch:
R–O–(CH2–CH2)n–CH2–CH2–CH2–CH2
R–O–(CH2–CH2)n–CH2–CH –CH2–CH3
(2.7)
R–O–(CH2–CH2)n–CH2–CH –CH2–CH3 + m CH2=CH2
R–O–(CH2–CH2) n–CH2–CH–(CH2–CH2)m 1–CH2–CH2
(2.8)
CH2–CH3
Other types of short branches are formed in similar chain isomerization reactions.
The rate of all such isomerization reactions does not depend on ethylene concentration
whereas the rate of the chain growth (the rate of Reaction 2.4) does. As a result, the
degree of branching increases at lower ethylene pressures in the reactor and at lower
temperatures.
Chain transfer reactions
Chain Transfer to Ethylene:
R–O–(CH2–CH2)n–CH2–CH2 + CH2=CH2
R–O–(CH2–CH2)n–CH=CH2 + CH3–CH2
(2.9)
Transfer to a special chain transfer agent R–H:
R –O–(CH2–CH2)n–CH2–CH2 + R–H
R –O–(CH2–CH2)n–CH2–CH3 + R
(2.10)
Propane is frequently used as a chain transfer agent in these reactions [2, 3]:
R –O–(CH2–CH2)n–CH2–CH2 + CH3–CH2–CH3
R –O–(CH2–CH2)n–CH2–CH3 + CH3–CH –CH3
(2.11)
All the chain transfer steps, Reactions 2.10, are followed by the growth of new polymer chains:
R + n CH2=CH2
R–(CH2–CH2)n 1–CH2–CH2
(2.12)
Chain Transfer within the Same Polymer Molecule (abstraction of a hydrogen atom)
Catalysts for the Manufacture of Polyethylene 31
R–(CH2–CH2)i–CH2–CH2–CH2–CH2–CH2
R–(CH2–CH2)i–CH2—CH2
CH2
\
CH2
\
/
CH2–CH2
R–(CH2–CH2)i–CH2—CH
\
CH3 CH2
\
/
CH2–CH2
(2.13)
Chain growth after Reaction 2.13 continues:
R–(CH2–CH2)i–CH2—CH + CH2=CH2
\
CH3 CH2
\
/
CH2–CH2
R–(CH2–CH2)i– CH2–CH–CH2–CH2
CH3–CH2–CH2–CH2
(2.14)
Chain Transfer to Another Polymer Molecule (abstraction of a hydrogen atom):
R–O–(CH2–CH2)n–CH2–CH2 + R1–(CH2–CH2)x–CH2–CH2–(CH2–CH2)y–R2
R–O–(CH2–CH2)n–CH2–CH3 + R1–(CH2–CH2)x–CH –CH2–(CH2–CH2)y–R2
(2.15)
A continuation of the chain growth after Reaction 2.15 leads to the formation of long
branches in LDPE molecules:
R1–(CH2–CH2)x–CH –CH2–(CH2–CH2)y– R2 + z CH2=CH2
R1–(CH2–CH2)x–CH–CH2–(CH2–CH2)y–R2
(2.16)
CH2–CH2–(CH2–CH2)z 2–CH2–CH2
The degree of long-chain branching increases at higher temperatures and at lower
ethylene pressures.
Chain termination reaction (radical disproportionation reaction):
R–O–(CH2–CH2)n–CH2–CH2 + CH2–CH2–(CH2–CH2)m–O–R
R–O–(CH2–CH2)n–CH2–CH3 + CH2=CH–(CH2–CH2)m–O–R
(2.17)
32 Handbook of Industrial Polyethylene and Technology
2.2.3 Types and Degree of Branching in Low Density Polyethylene Resins
Molecular chains of LDPE resins contain various short-chain branches. They are
formed in Reactions 2.5, 2.6, 2.13, 2.14 and in other similar chain isomerization reactions. The total amount of short-chain branches and their types are usually determined
by the 13C NMR method. The average number of short-chain branches in LDPE resins ranges between 20 and 25 CH3/1000C. The most abundant among the short-chain
branches are n-butyl branches formed in Reactions 2.13 and 2.14 (their content is
~5 to 6 per 1000C), ethyl branches produced in Reaction 2.6 (~4 per 1000C), n-amyl
branches formed in reactions similar to Reactions 2.13 and 2.14 via a six carbon atom
cycle (~2 per 1000C), whereas the contents of n-hexyl and longer linear branches are
lower, ~1 to 2 per 1000C [4, 5].
Long-chain branches in LDPE chains are generated in chain transfer reactions to
dead polymer chains, Reactions 2.15 and 2.16. The content of long-chain branches can
be evaluated by comparing results of size exclusion chromatography in combination
with on-line viscosity measurements or in combination with on-line low-angle light
scattering measurements [4, 5]. The average number of long-chain branches in LDPE
molecules ranges from 0.5 to 1 per 1000C.
2.3 Catalytic Synthesis of Polyethylene Resins
Because PE resins produced in catalytic polymerization reactions have no or small
amounts of long-chain branching, they are called linear PE resins to distinguish them
from LPPE resins produced in radical polymerization reactions. Commercial synthesis
of linear PE resins is carried out with the use of transition metal catalysts of different
types [6–9]. Three groups of such catalysts are especially important, chromium-based
catalysts, titanium-based catalysts (Ziegler-Natta catalysts), and metallocene catalysts.
2.3.1
Commercial Technologies of PE Manufacture
Several types of reactors are used for ethylene polymerization reactions with transition
metal catalysts [10, 11]:
Gas Phase Polymerization Technology: A gas-phase polymerization reactor is a large
cylindrical tower with a height of up to 25 m and with a length-to-diameter ratio of ~7.
It usually operates at a pressure of 1.5 to 2.5 MPa (15 to 25 atm) and at a temperature
from 70 to 100 °C. The reactor is half-filled with a bed of polymer particles, which
is vigorously agitated and mixed by a high-velocity gas stream. The gas stream enters
the reactor through a perforated distribution plate at the reactor’s bottom; it fluidizes
the bed of the polymer particles and removes the heat of polymerization. The stream
includes ethylene, an -olefin (in copolymerization reactions), hydrogen, which is used
for molecular weight control, and nitrogen (an inert component). The gas mixture exits
the reactor at its top; it is compressed, cooled, its composition is reconstituted, and the
gas is returned to the reactor.
A solid catalyst in the form of small particles is continuously fed into the reactor. On
average, these particles remain in a reactor for 2.5 to 4.0 hours. As soon as a catalyst
Catalysts for the Manufacture of Polyethylene 33
particle enters a reactor, it starts polymerizing ethylene and, as a result, the size of an
average polymer particle gradually increases 15 to 20 times compared to the diameter
of the original catalyst particle. The particulate resin is continuously removed from the
reactor.
Solution Polymerization Technologies: Two types of a high-temperature solution
polymerization technology are used for the LLDPE manufacture. Processes of the first
type utilize hydrocarbon solvents and the processes of the second type are carried out in
mixtures of supercritical ethylene and molten PE. Both solution processes usually operate at 130 to 180 °C, their reaction pressure varies from 3 to 20 MPa (30 to 200 atm),
the ethylene content in the reactors ranges between 8 and 10%, and the residence time
varies from 5 to 15 minutes. These processes can accommodate both particulate catalysts and soluble metallocene catalysts. The product stream containing up to 30–35 wt%
of a molten polymer leaves the reactor and is discharged into a strip vessel where the
polymer is separated from unreacted monomers and the solvent and then pelletized.
Slurry (Suspension) Polymerization Technology: Most slurry reactors for the production of HDPE and LLDPE resins are built as large folded loops containing vertically
positioned long runs of pipe 0.5 to 1 m in diameter connected by short horizontal
stretches of the same pipe. The loop reactors operate at a pressure of up to 3 MPa
(30 atm) and at a temperature from 60 to 75 °C. The reactor is filled with slurry of polymer particles suspended in a low-boiling solvent, usually isobutane or isopentane. An
internal pump forces a high-speed circulation of the suspension through the loop. The
concentration of polymer particles in the slurry is maintained at ~20 to 25 wt%. The
residence time of polymer particles in the reactor ranges between 1.5 and 3 hours and
the ethylene conversion in the process is usually very high, 97 to 98%.
2.3.2 Chromium-Based Catalysts
Several types of chromium-based catalysts are used in ethylene polymerization reactions. Two of them, chromium oxide catalysts and supported organochromium catalysts, have great commercial importance. Taken together, these two catalysts account
for nearly 40% of all the PE resins manufactured throughout the world.
2.3.2.1 Chromium Oxide Catalysts (Phillips Catalysts)
Chromium oxide catalysts are primarily used for the manufacture of ethylene homopolymers, HDPE resins. Most materials produced with chromium oxide catalysts have
practically linear polymer chains and a broad molecular weight distribution [6, 7, 9],
whereas some chromium oxide catalysts prepared with special grades of silica produce
PE resins with a significant content of long-chain branching [12, 13]. Chromium oxide
catalysts are also used for the synthesis of ethylene/α-olefin copolymers. These copolymers are usually designated as low density linear PE resins (LDLPE resins) to distinguish them from more common LLDPE resins.
Chromium oxide catalysts for the manufacture of HDPE resins were originally developed by Phillips Petroleum Company in USA in 1968 [6, 7, 15–17]. These catalysts are
supported on inert porous substrates, mostly on silica, although other materials such as
silica-alumina, alumina and amorphous alumophosphate, AlPO4, are also used as the
34 Handbook of Industrial Polyethylene and Technology
supports [6–9, 14–26]. The average particle size of the supports is 180–250 μm; their
specific surface area is very high at over 300 m2/g. The structure and porosity of silica
in the catalysts have a strong influence on their activity in ethylene polymerization
reactions and on the character of the polymer they produce [12–14]. Synthesis of silica
particles with various ratios of large and small pores affords the production of chromium oxide catalysts suitable for the synthesis of PE resin with a controlled amount of
long-chain branching [27].
The synthesis procedure of a chromium oxide catalyst consists of two main steps.
First, a catalyst support is impregnated with a source of chromium, such as aqueous
solution of chromic acid or chromium nitrate, or with alcoholic solution of chromium
acetate. The solvent is removed at an elevated temperature and then the solid material is
dehydrated at ~200 °C. In the second step, the chromium species in the catalyst precursor are activated by calcination of catalyst particles in a fluidized-bed reactor at 500 to
850 °C in a dry oxidizing environment.
The calcination step results in anchoring the chromium species to the surface of
the support and their oxidation to various CrVI complexes [18–25]. The Cr content in
Phillips catalysts is usually low, from 0.5 to 1 wt%, which corresponds to 2 to 4 Cr atoms
per a 10 nm2 area on the support surface. Depending on the activation conditions, such
a catalyst contains from 50 to ~100% of its CrVI species in the form of silyl monochromates and bichromates [15, 22–25, 28–31]:
O
O
2 SiOH + CrO3
Cr
O
Si
Si
(2.18)
O
The formation of the chromates on the silica surface is practically complete when
the calcination temperature reaches 500 °C. However, the activity of the catalysts can
be further increased by increasing the calcination temperature due to a gradual elimination of remaining OH groups from the silica surface [15]. The maximum activity is
reached after the catalysts are calcined at 925 °C [15, 18–25].
When such a catalyst is introduced into an ethylene polymerization reactor, the chromate species formed in Reaction 2.18 are reduced to anchored CrII species [15, 22–25]:
( Si–O)2CrO2 + C2H4
( Si–O)2Cr + 2CH2O
(2.19)
Chromium oxide catalysts can be also pre-activated with carbon monoxide at
300–350 °C [22–25, 28–31] or with organomagnesium compounds [32]. Both these
treatments produce the same effect; they reduce the CrVI species to CrII and CrIII.
As typical for many catalysts, the synthesis of chromium oxide catalysts is a combination of science and art. For example, the average molecular weight of HDPE resins produced with the catalysts and the content of long-chain branches in them are
strongly influenced by the physical strength of the silica particles (which fragment during the polymerization reactions) and by the size of pores in the support [12–14]. Silica
with wide pores produces easily fragmented, highly active chromium oxide catalysts
Catalysts for the Manufacture of Polyethylene 35
whereas silica with narrow pores produces a different type of catalysts which yield PE
resins with a significant fraction of long-chain branches.
The performance of the catalysts can be further improved by their chemical modification prior to the calcination step. The most successful modifications include the
treatment of the catalysts with Ti(OR)4 or with (NH4)2SiF6 [15, 30, 32], as well as the
treatment of the alumina support with small amounts of phoshates [26].
The productivity of chromium oxide catalysts under typical polymerization conditions (at ~80–90 °C) is very high, ~500 kg/(mol Cr∙atm) in one hour. The catalysts
produce PE resins with a broad molecular weight distribution. The molecular weight of
these resins is controlled mostly by the reaction temperature. An important difference
between chromium oxide catalysts and Ti-based Ziegler-Natta catalysts (Section 2.3.3)
is nearly complete insensitivity of chromium oxide catalysts to the presence of hydrogen in a polymerization reaction.
2.3.2.2
Organochromium Catalysts
The second family of chromium-based catalysts for the synthesis of HDPE resins utilize organochromium compounds, such as bis(triphenylsilyl)chromate,
[(C6H5)3Si–O]2CrO2, or bis(cyclopentadienyl)chromium, (C5H5)2Cr. Several supported
organochromium catalysts are commercially significant [6, 7, 9, 33–35].
A typical catalyst of this family is prepared in two steps. First, [(C6H5)3Si–O]2CrO2
is deposited into the pores of pre-calcined silica suspended in a hydrocarbon slurry
[36]. The chromate molecules form complexes with isolated OH group on the silica
surface [37]:
[(C6H5)3Si–O]2CrO2 + H–O–Si
(C6H5)3Si–O–Cr(O2)–O–Si[(C6H5)2]–Ph…H–O–Si
(2.20)
In the second step, diethylaluminum ethoxide, (C2H5)2Al–O–C2H5, is added to the
slurry. It reduces the CrVI species and produces various CrII and CrIII species covalently
bonded to the silica surface,
(C6H5)3Si–O–Cr(O2)–O–Si[(C6H5)2]–Ph…H–O–Si + (C2H5)2Al–O–C2H5
LxCr–O–Al(C2H5)–O–Si
(2.21)
where L can be O= or (C6H5)3Si–O–.
This organochromium catalyst and chromium oxide catalysts described in Section
2.3.2 are inter-convertible. For example, if the organochromium catalyst produced in
Reaction 2.21 is calcined under vacuum or in a stream of oxygen, it is converted into
a chromium oxide catalyst [37]. Alternatively, a conventional chromium oxide catalyst
can be converted into a silyl chromate catalyst by reacting it with (C6H5)3Si–OH under
moderate conditions [35]:
( Si–O)2CrO2 + [(C6H5)3Si–OH
Si–O–Cr(O2)–O–Si(C6H5)3 + Si–OH
(2.22)
36 Handbook of Industrial Polyethylene and Technology
Another type of an organochromium catalyst uses a similar approach, supporting
an organochromium compound on a silica surface, but it employs a different chromium source, chromocene (C5H5)2Cr. A silica support is calcined at 600 °C and then
(C5H5)2Cr is added. It reacts with the remaining OH groups on the surface of the support in a toluene slurry [33, 34]:
(C5H5)2Cr + H–O–Si
(C5H5)Cr–O–Si + cyclopentadiene
(2.23)
Supported chromocene catalysts are employed in ethylene polymerization reactions
at 90 to 110 °C, they are highly active (productivity at 80 °C varies from 5 to 7 kg PE/g
cat in one hour) and produce PE materials with a medium molecular weight distribution. Molecular weight of these resins is readily controlled with hydrogen, in contrast
to chromium oxide catalysts. Supported chromocene-derived catalysts are predominantly used for the synthesis of low molecular weight HDPE resins suitable for injection molding.
2.3.3 Titanium-Based Ziegler-Natta Catalysts
All Ziegler-Natta catalysts consist of two components. The first component, a solid
powder or a liquid, is called “catalyst”; it contains a derivative of titanium. The second
component is called “cocatalyst”; it is an organoaluminum compound, most often, triethylaluminum (Al(C2H5)3). The first catalyst system of this type was developed by Karl
Ziegler in Germany in 1953. In the industry, titanium-based Ziegler-Natta catalysts
are employed at temperatures between 80 and 95 °C and at the ethylene pressure of 0.7
to 1.5 MPa (7 to 15 atm). Ziegler-Natta catalysts are used for the synthesis of several
grades of PE, both homopolymers (HDPE) and copolymers of ethylene with α-olefins
(LLDPE and VLDPE resins). Over 75% of all LLDPE resins are produced with these
catalysts. Most of Ziegler-Natta catalysts currently used in industry are supported. A
variety of techniques were developed for supporting titanium compounds, usually titanium tetrachloride (TiCl4), on specially prepared supports such as silica and microcrystalline magnesium dichloride (MgCl2) [6–7, 9, 14].
2.3.3.1
General Features of Commercial Ziegler-Natta Catalysts
The two most important commercial processes for the synthesis of PE resins with
Ziegler-Natta catalysts are polymerization in gas-phase reactors and polymerization in
slurry reactors utilizing light hydrocarbon solvents. The selection of a solid titaniumbased catalyst for the use in a particular ethylene homopolymerization or an ethylene/
α-olefin copolymerization reaction is based on four technical criteria:
(a) The catalysts should be able to produce PE resins with a desirable combination of end-use properties, such the mechanical strength of film
manufactured from the HDPE or the LLDPE resin, cling properties of
the LLDPE film, rigidity of articles made of HDPE, the content of an
extractable material, etc. [38].
Catalysts for the Manufacture of Polyethylene 37
(b) Physical features of the catalyst which determine its behavior in a particular polymerization reactor. These features include the shape of the
catalyst particles (a spherical shape is preferred), the average particle
size, and the absence of very small catalyst particles (catalyst fines).
(c) Technological features of the catalyst: moderate activity, high reactivity
toward α-olefins, a sufficient response to the presence of hydrogen (the
agent of molecular weight control), etc.
(d) Processability of resins produced with the catalyst, the function of the
resin’s molecular weight and molecular weight distribution. This factor is important for manufacturers of such large-volume products as
LLDPE film or HDPE containers.
The gross recipe of most commercial supported Ziegler-Natta catalysts can be represented by a common scheme:
Active titanium ingredient/Support/Carrier
The active titanium ingredient is nearly always either TiCl4 or titanium trichloride
(TiCl3). The support in the majority of the catalysts is microcrystalline MgCl2. The
principal role of the support is to increase the activity of the titanium ingredient. This
increase can be quite high. For example, the activity of a typical TiCl3/MgCl2 catalyst
(expressed as the polymer yield per gram of catalyst) can be 10 to 30 times higher compared to the activity of TiCl3 used alone.
The third component of a commercial Ziegler-Natta catalyst is the carrier. A carrier is a catalyst component that does not affect the catalyst performance as such but
determines the size and the shape of the catalyst particles and the polymer particles
produced from them. The carriers usually used in Ziegler-Natta catalysts designed for
ethylene polymerization are either silica particles or particles of MgCl2, which in this
case serves both as a support and a carrier.
Spherical particles of amorphous silica have a diameter ranging from 30 to 40 μm.
They are highly porous (pore volume 1 to 3 mL/g) and, as a consequence, have a high
specific surface area, 250 to 400 m2/g. Commercially produced silica has a high water
content. To remove water and to reduce the concentration of silanol (Si-OH) groups
on the internal surface of the particles, silica is calcined at a temperature from 500
to 700 °C [14]. The calcination step removes all physically absorbed water and most
silanol groups, as shown in Table 2.2.
Only two types of silanol groups are left on the silica surface after the calcination,
isolated Si-OH groups and vicinal Si-OH groups, two OH groups bound to adjacent Si
atoms. In some cases, these remaining silanol groups should be completely eliminated.
Table 2.2 Concentration of silanol groups in calcined silca as a function of temperature [14].
Calcination temperature, °C:
300
400
500
600
700
800
[Si-OH], mmol/g SiO2:
>1.6
~1.3
~1.1
0.7–0.8
0.5–0.6
0.4–0.5
~0.4
~0.6
~0.7
~0.8
Fraction of isolated Si-OH:
38 Handbook of Industrial Polyethylene and Technology
Treating calcined silica with an organometallic compound, usually with Al(C2H5)3, can
be used for the purpose [6–7, 9, 18–25]:
Si-OH + Al(C2H5)3
Si−OH + Al(C2H5)3
Si
Si-O-Al(C2H5)2 + C2H6
Si−O−Al(C2H5)2 + C2H6
Si
OH
O
O
Al
O
+ Al(C2H5)3
Si
(2.24)
Si
OH
C2H5 + 2C2H6
(2.25)
O
Another reaction often used to remove the remaining silanol groups from calcined
silica is its treatment with highly concentrated TiCl4 solution or with neat TiCl4:
Si-OH + TiCl4
Si−OH + TiCl4
Si
Si
+ TiCl4
OH
(2.26)
Si−O−TiCl3+ HCl
OH
O
Si
Si-O-TiCl3 + HCl
O
O
Si
O
(2.27)
Cl
+ 2HCl
Ti
Cl
All synthetic recipes of Ziegler-Natta catalysts based on silica as a carrier can be
subdivided into two classes depending on the method which is used to introduce
the MgCl2 support into the pores of the silica. Some catalyst recipes use for this purpose commercially available anhydrous MgCl2. The main difficulty in this case is to
make MgCl2 soluble in an organic solvent, the step necessary for the impregnation of
porous silica. As an alternative, to avoid this difficult step, MgCl2 can be synthesized
directly inside the silica pores during the catalyst preparation. The sources of magnesium in such synthetic routes are organomagnesium compounds Mg(OR)2, MgR2,
mixtures of MgR2 and AlR3, or Grignard reagents. A source of chlorine is needed to
convert Mg(OR)2 or MgR2 into MgCl2; it is usually TiCl4, HCl, or a chloro-organic
compound.
The number of patented Ziegler-Natta catalysts and the number of the catalysts used
commercially is quite large. Below, several examples of their typical synthetic procedures are described.
2.3.3.2 Catalysts Produced from Soluble MgCl2 Complexes
Catalyst recipes based on TiCl4, MgCl2, and tetrahydrofuran
Two components for the catalyst synthesis are prepared separately, solution of an
active ingredient and a silica carrier:
(a) The active ingredient in these catalysts is a Ti-Mg complex, which is
formed in a reaction between TiCl4 and MgCl2 in a tetrahydrofuran
(THF) solution [6–7, 9, 39–42]. Both TiCl4 and MgCl2 are dissolved in
THF and then the two solutions are combined.
Catalysts for the Manufacture of Polyethylene 39
(b) The silica support is calcined at 600 °C and then treated with Al(C2H5)3
to remove the remaining silanol groups in Reactions 2.24 and 2.25
[40].
To prepare the catalyst, the TiCl4/MgCl2/THF solution is added to the slurry of the
pre-calcined silica in THF. After the solvent is removed by evaporation at 50–60 °C,
microcrystals of the Ti-Mg ionic complex are formed inside silica pores [39–42]:
TiCl4  (THF)2 + 2 MgCl2  (THF)2 —(THF)
[Mg2( -Cl)3  (THF)6]+ [TiCl5  THF]
(2.28)
This complex, when combined with Al(C2H5)3 cocatalyst, forms a very active ethylene polymerization catalyst. However, it has several deficiencies including relatively
poor reactivity toward α-olefins. The performance of this catalyst can be improved if
a part of THF is removed from the complex formed in Reaction 2.28. This is achieved
by treating the catalyst with a mixture of Al(C2H5)3 and Al(C2H5)2Cl prior to its use in
polymerization reactions [43].
Catalyst recipes based on MgCl2/alcohol complexes
MgCl2 forms numerous complexes with alcohols. The [alcohol]:[MgCl2] ratio in the
complexes depends on the type of alcohol and the temperature at which the complex
is formed [44, 45]. These complexes have two advantages as semi-products for the synthesis of supported Ziegler-Natta catalysts, they melt at 80 to 100 °C, at a much lower
temperature than MgCl2 itself, and they are readily soluble in alcohols.
One recipe for the preparation of a catalyst of this type uses MgCl2 both as the support and as the carrier. The procedure consists of two stages, synthesis of the MgCl2
support/carrier and its impregnation with TiCl4 [44–48]. To prepare the support/
carrier, MgCl2 is mixed with a small amount of ethanol and with an inert dispersant,
paraffin oil. After the mixture is heated to 120 °C, MgCl2 forms the MgCl23C2H5OH
complex, which melts and forms a fine emulsion in the paraffin oil. The emulsion is
rapidly dispensed into a large volume of a cold alkane. This step leads to crystallization of the MgCl23C2H5OH complex into spherical particles that are 60 to 100 mm in
diameter.
Two different procedures can be used to convert the particulate MgCl23C2H5OH
complex into a titanium-based supported catalyst. The principal goal in both recipes is
to remove ethanol from the complex and to turn it into porous particles of MgCl2. In
the first procedure [46], ethanol is removed from the complex with neat TiCl4 by heating the slurry of the MgCl23C2H5OH particles in a mixture of an alkane and TiCl4 from
20 to 130 °C [6–7, 9]. This treatment directly produces the final TiCl4/MgCl2 catalyst
containing ~7.5 wt% of titanium.
In an alternative procedure, ethanol is removed from the MgCl23C2H5OH complex
by a treatment with Al(C2H5)3 at a high temperature [44, 47–48]. First, spherical particles
of the complex are combined with Al(C2H5)3 solution and heated from 50 to 150 °C until
the [alcohol]:[MgCl2] ratio in the complex decreases from 3 to ~1. Then the remaining
40 Handbook of Industrial Polyethylene and Technology
ethanol is removed at 60 °C by repeated treatment with Al(C2H5)3. Finally, the generated
spherical MgCl2 support is reacted with the source of titanium, Ti(OC4H9)4.
Both these catalysts have high activity in ethylene polymerization reactions,
~100 kg/g Ti, and produce large spherical polymer particles.
2.3.3.3 Catalysts Produced by Synthesis of MgCl2
A large variety of these catalysts are described in the literature. Two examples below
provide the gist of the procedures usually employed for their preparation.
The first recipe uses MgCl2 both as the carrier and the support in the finished catalyst
[6–7 ,9, 44, 49–50, 51]. MgCl2 particles are prepared from a dialkylmagnesium compound,
MgR2. Its solution in an alkane is mixed with di-iso-amyl ether at an [ether]:[MgR2] ratio
of ~0.4, leading to the formation of an MgR2-ether complex. Then the solution of the
complex is vigorously stirred and heated to 50 °C. Tert-butyl chloride is slowly added to
it at a [tert-C4H9Cl]:[MgR2] ratio of ~2 leading to precipitation of MgCl2.
MgR2 + 2 tert-C4H9Cl
MgCl2 + 2 tert-C4H9-R
(2.29)
The MgCl2 generated in Reaction 2.29 forms spherical porous particles that are 25
to 50 mm in diameter with a specific surface area of ~40 m2/g. In the second step,
the MgCl2 support is reacted with TiCl4 and with AlR2Cl. This step is carried out in a
hexane-ether slurry at 50 to 80 °C. The principal chemical reaction at this step is the
reduction of TiIV species to TiIII:
TiCl4 + AlR2Cl
TiCl3 + AlRCl2 + R R
(2.30)
The second group of such catalysts uses magnesium compounds as a support and silica as a carrier [6–7, 9, 52–55]. Spherical amorphous silica is calcined at 600 °C and then
its slurry in an alkane is treated with Mg(C4H9)2 at 50 to 60 °C. The organomagnesium
compound migrates into the silica pores and partially reacts with the silanol groups,
similarly to Reactions 2.24 and 2.25. In the next step, an organic compound (a modifier)
is added to the slurry to form a catalyst precursor. If the modifier is CCl4 or CCl3CH3, it
converts Mg(C4H9)2 into MgCl2, and if the modifier is an alcohol or Si(OR)4, it converts
Mg(C4H9)2 into Mg(OR)2. Finally, the catalyst precursor is treated with TiCl4 at 50 to
60 °C. The main advantage of this catalyst preparation scheme is its versatility. Catalysts
with different reactivity toward α-olefins are produced merely by changing the nature of
the modifying organic compound or by varying its amount. These catalysts can be used
with two different types of cocatalysts, AlR3 and AlR2Cl [52–54]. Examples in Table 2.3
demonstrate inter-dependencies between the catalyst recipes (the modifying compound
and the type of cocatalyst), and the productivity of the catalyst systems and the molecular weight distribution of ethylene/1-hexene copolymers (LLDPE resins) they produce.
2.3.3.4
Ziegler-Natta Catalysts Used in Solution Polymerization Processes
A special type of Ziegler-Natta catalysts is employed in industry in high-temperature
polymerization processes for the synthesis of ethylene/1-octene copolymers, both
Catalysts for the Manufacture of Polyethylene 41
Table 2.3 Performance of catalysts containing synthesized MgCl2 in ethylene/1-hexene copolymerization reactions* [9, 52–54].
Modifying organic
compound
Cocatalyst
Productivity,
g/(g cat∙h)
Molecular weight
distribution, Mw/Mn
none
Al(C2H5)3
1250
~10
n-C4H9OH
Al(C2H5)3
2130
5.5
CCl4
Al(C2H5)3
3400 to 4660
~6
CCl4
Al(C2H5)2Cl
2610
6.9
CCl3CH3
Al(C2H5)3
3080
6.7
Si(OC2H5)4
Al(C2H5)3
3900
5.0
Si(OC2H5)4
Al(CH3)3
3460
3.6
Si(OC2H5)4
Al(CH3)2Cl
1570
5.6
*Reactions at 80 °C; reaction conditions are selected to produce copolymers containing 1.5 to 1.7 mol%
of 1-hexene.
LLDPE and VLDPE resins [10–11]. These polymerization reactions are carried out
in high-boiling hydrocarbon solvents, and the PE remains dissolved in such solvents
at 130 to 200 °C throughout the whole polymerization cycle [8, 56]. Therefore, such
requirements to the catalysts as the spherical shape of catalyst particles and their size
become irrelevant. All Ziegler-Natta catalysts rapidly loose activity at temperatures
above 100 °C and, for this reason, a different requirement to the catalysts becomes
paramount; high initial activity.
One such catalyst uses Ti(Oiso-C3H7)4 as a source of titanium and MgCl2 as a support [57]. The support is prepared by reacting a solution of Mg(C4H9)(C2H5) with gaseous HCl in a heavy hydrocarbon solvent at 25 to 50 °C:
Mg(C4H9)(C2H5) + 2 HCl
MgCl2 + C4H10 + C2H6
(2.31)
Slurry of microcrystalline MgCl2 generated in Reaction 2.31 is reacted, in sequence,
with Al(C2H5)Cl2, Ti(Oiso-C3H7)4, and with Al(C2H5)3. The resulting product is used as
a single-component catalyst stream to copolymerize ethylene and 1-octene in solution
at 180 to 200 °C and at a residence time from 15 to 20 minutes.
Catalysts of the highest activity are formed directly in a polymerization reactor [58–60].
Combining an organomagnesium compound such as Mg(C4H9)2 and a source of chlorine atoms, tert-C4H9Cl, at a high temperature produces a finely dispersed MgCl2 support
(Reaction 2.29). The active ingredient in the catalysts of this type is TiCl4 and the cocatalyst is Al(C2H5)2OC2H5; they both are also directly injected into the reactor [58–60].
2.3.4 Metallocene Catalysts
Metallocene catalysts are mostly used for the synthesis of HDPE and LLDPE resins
with a narrow molecular weight distribution and high compositional uniformity (see
42 Handbook of Industrial Polyethylene and Technology
its definitions in Section 2.5). The modern period in the history of these catalysts has
begun in Germany in 1976 when Kaminsky and Sinn discovered highly active catalyst systems based on zirconocene complexes. These catalysts are employed either as
solutions or they are supported on inert carriers. All these catalysts contain two components, the first one is a metallocene complex of a transition metal and the second
component, a cocatalyst, is either a special organoaluminum compound, methylalumoxane (abbreviated as MAO), or a fluoroaromatic boron compound.
2.3.4.1 Metallocene Complexes and Cocatalysts for Them
A large number of metallocene complexes can be used as components of soluble polymerization catalysts. Figure 2.1 shows the structures of several such complexes. Each of
them contains either one or two cyclopentadienyl rings which are π-coordinated to a
transition metal atom. Three metals are typically used, zirconium, titanium, or hafnium.
Complexes I – Bis(cyclopentadienyl) complexes [(Ri C5H5-i)2]MX2 (abbreviated as Cp2MX2) The substituent R in the cyclopentadienyl rings is usually a small alkyl group, from methyl to n-hexyl and i = 1 in the majority
of cases. The X group in these complexes is usually Cl.
Complexes II – Bis(indenyl) complexes Ind2MX2. Each indenyl ligand can
carry alkyl substituents.
Complexes III – Bis(tetrahydroindenyl) complexes (Ind-H4)2MX2.
Complexes IV – Bridged bis-metallocene complexes of the general formula
(Bridge)(Cp )(Cp )MX2, where the two cyclopentadienyl groups, Cp and
Cp , can be the same or different. The bridge connects the two cyclopentadienyl rings; the most often used bridges are –CH2 CH2 , –CH(C6H5)
CH(C6H5) , >C(CH3)2, >C(C6H5)2, >Si(CH3)2, and >Si(C6H5)2.
Complexes V – Bridged bis-indenyl complexes of the general formula
(Bridge)(Ind)2MX2 with a bridge connecting the two indenyl rings.
R
X
X
X
M
M
M
X
X
X
R
I
II
III
X
X
M
Bridge
X
MX2
Bridge
M
Bridge
N
R'
IV
V
VI
Figure 2.1 Metallocene complexes used as catalysts in ethylene polymerization reactions.
X
Catalysts for the Manufacture of Polyethylene 43
Table 2.4 The effect of the [MAO]:[Zr] ratio in an ethylene/1-hexene copolymerization
reaction with the (C5H5)2ZrCl2-MAO system at 40 °C.
Molar [MAO]:[(C5H5)2ZrCl2] ratio
100
200
500
1000
2000
Yield, g/mmol (C5H5)2ZrCl2
6
64
454
459
680
Complexes VI – Bridged monocyclopentadienyl complexes of the general
formula (Bridge)[(Ri Cp)(R A)]MX2. The single cyclopentadienyl group
Ri Cp in these complexes is usually alkyl-substituted; i ranges from 1 to
4. This cyclopentadienyl group is π-coordinated to the metal atom M and
additionally connected to M via a σ-link, –bridge -A(R )-, containing a
heteroatom A, either N or P. The preferred substituent R at the heteroatom
A is a bulky alkyl group, iso-C3H7 or tert-C4H9. The bridge between the
cyclopentadienyl group and the heteroatom can be >Si(CH3)2, –(CH2)n ,
>C(C6H5)2, >C(CH3)CH2CH2 CH=CH2, or –Si(CH3)2 O Si(CH3)2 –
[9, 56, 61–70]. This type of bridged monocyclopentadienyl complexes
is often called “constrained-geometry” complexes. The commonly used
complex of this structure is the titanium complex [(CH3)2Si[(CH3)4Cp]
(t-C4H9-N)TiCl2.
Complexes I through V are widely used in ethylene polymerization reactions as solutions in aromatic hydrocarbons. They can also be used to prepare supported metallocene
catalysts. A common support for such catalysts is calcined silica [14]. Polymerization
catalysts prepared from Complexes VI exhibit significant stability above 120 to 130 °C;
they are preferred for ethylene/α-olefin copolymerization reactions in high-temperature solution processes.
Two types of cocatalysts are especially suitable for converting metallocene complexes
into polymerization catalysts. The first one is MAO [6, 7, 9, 71–74]. MAO is produced
by reacting Al(CH3)3 with water. MAO is a member of a special class of oligomeric
organoaluminum compounds with the alumoxane backbone, [Al(R)–O]m, where R =
CH3 in MAO, iso-C4H9 in isobutyl alumoxane, tert-C4H9 in tert-butylalumoxane, etc.
[9, 71–72], and m ranges from 12 to 20 in MAO [75–78] and from 6 to 9 in tert-butylalumoxane [79–80]. The structure of MAO and tert-butylalumoxane was thoroughly
investigated [75, 79–80]. Both products are mixtures of different oligomers with the
cage molecular structure; each cage is composed of hexagonal and square faces.
The productivity of metallocene catalysts in ethylene polymerization reactions
strongly depends on the [MAO]:[transition metal] ratio: the higher the ratio, the higher
is the productivity per mole of metallocene complex [9, 71, 72, 81, 82]. The following
example (Table 2.4) illustrates the effect of the [MAO]:[Zr] ratio in an ethylene/1-hexene copolymerization reaction with the (C5H5)2ZrCl2-MAO system at 40 °C (5-minute
reaction in toluene) [83].
The second family of cocatalysts suitable for converting metallocene complexes into
ethylene polymerization catalysts is ion-forming cocatalysts [84–97]. The most important among them are two types of organoborane compounds:
44 Handbook of Industrial Polyethylene and Technology
1. Perfluorinated triphenylborane, B(C6F5)3 [93, 94–97]. This cocatalyst is
used in polymerization reactions in combination with Al(CH3)3. The role
of Al(CH3)3 in these catalyst systems is to convert metallocene complexes
Cp2ZrCl2 (here Cp stands for any π-ligand with a cyclopentadienyl ring)
into Cp2Zr(CH3)2.
2. Perfluorinated phenylborate salts such as [R3NH]+ [B(C6F5)4]– or
[C(C6H5)3]+ [B(C6F5)4]– [88, 89].
Metallocene catalysts differ quite noticeably in their ability to polymerize ethylene
and to copolymerize it with α-olefins [9, 81]. The principal features of metallocene complexes that affect these two aspects of catalyst reactivity are the type of the transition
metal atom M, the type and the number of substituents at the cyclopentadienyl groups
(see Figure 2.1), and the presence of bridges between the two rings. One example for
bis-indenyl complexes illustrates this effect in ethylene/1-hexene copolymerization
reactions at 75 °C with MAO as a cocatalyst [9], as shown in Table 2.5.
2.3.4.2 Supported Metallocene Catalysts
All the recipes for supported metallocene catalysts have several common features
[6–7, 9]. The preferred carriers for these catalysts are spherical particulates of silica or
alumina. Several strategies are used to deposit and immobilize MAO in the pores in the
carrier particles or to synthesize MAO inside the pores.
Recipes for supported metallocene catalysts containing immobilized MAO include
two steps, a reaction of the carrier with MAO followed by a reaction of immobilized
MAO with a metallocene complex [73, 98–102]. Commercial MAO always contains
a significant quantity of free Al(CH3)3, and the latter reacts with silanol groups in silica similarly to Reaction 2.24. MAO does not react with the silanol groups directly
but it is adsorbed on the Al(CH3)3-modified silica surface. Silica-immobilized MAO
is a commercially available product. This material can be impregnated with solution
of a metallocene complex, either in a separate catalyst preparation step or directly in
a polymerization reactor prior to introduction of the ethylene [101, 102]. Premixed
MAO-metallocene compositions can also be used for the preparation of supported
metallocene catalysts containing immobilized MAO.
A different approach uses a strong acid as a part of the carrier [103]. A calcined
carrier, silica, or alumina, is treated with a concentrated solution of sulfuric acid in
alcohol, then it is calcined the second time at a high temperature and impregnated
with a metallocene complex. These supported catalysts are activated with common
Table 2.5 Productivity for ethylene/1-hexene copolymerization reactions for bis-indenyl complexes at 75 °C with MAO as a cocatalyst [9].
Metallocene complex
Ind2ZrCl2
C2H4(Ind)2ZrCl2
(CH3)2Si(Ind)2ZrCl2
Complex structure
Complex II
Complex V
Complex V
Productivity, kg/(mmol Zr × h)
155
420
340
Catalysts for the Manufacture of Polyethylene 45
organoaluminum cocatalysts. Acidic centers in the carrier convert the metallocene
complex into a metallocenium cation, which is the active polymerization center (see
Section 2.4).
2.3.5 Post-Metallocene Ethylene Polymerization Catalysts
The expression “post-metallocene catalysts” has a historic rather than any particular
chemical meaning; most of these polymerization catalysts were discovered several
decades after the discovery of metallocene catalysts. In practice, the term “post-metallocene catalysts” embraces a large variety of complexes containing transition metal atoms
and different bidentate and multidentate organic ligands. Figure 2.2 shows several representative samples of such complexes containing early transition metals and Figure 2.3
shows some complexes of late transition metals. The metals in these complexes vary
Complexes with diphenoxy ligands
O
Ti
O
R
O
R
O
Ti
O
R
Ti
R
O
R
R
Complexes with diamide ligands
R
N
X
R = 2,6-iso-(C3H7)2-C6H3, 2,6-(CH3)2-C6H3
X
X = Cl, CH3, CH2-C6H5
Ti
N
R
Complexes with phenoxy-imine ligands
R
R
M= Ti, Zr
O
R’
N
Cl
Ar = C6H5, C6F5, 3,5-F2-C6H3
R’
Ar
M
N
Cl
R = CH3, t-C4H9, halogen atoms
Ar
R’ = H (aldimine complexes),
R’ = alkyl group (ketimine complexes)
O
R
Complexes with saligenin ligands
t-C4H9
Complexes with dialkylcarbamato ligands
R
O
CF3
t-C4H9
O
O
Ti
Cl
CF3
C
Ti
O
R 4
CH3
O
TiCl2
N
C
O
N
CH3 2
R = C2H5,CH3
Cl
Figure 2.2 Examples of complexes of early transition metals used as components of post-metallocene
catalysts.
46 Handbook of Industrial Polyethylene and Technology
Complexes with -diimine ligands
M = Ni, Pd
X = Cl, Br, CH3, NCS
R
R
N
R'
R’ = H, CH3, C2H5
X
R’+R’ =
M
-Dinaphthyl, cyclohexyl
X
R'
R = CH3, iso-C3H7, CF3, C6H5, C6F5
N
R
R
Complexes with bis(imino)pyridyl ligands
M = Co, Fe, V
X1 and X2 = CH3, iso-C3H7, t-C4H9, Cl, Br
R
N
X1
N
X3
R
M
Cl
X2
X3 = H, CH3, C2H5, iso-C3H7
X1
N
R = H, CH
Cl
X2
X3
Figure 2.3 Examples of complexes of late transition metals used as components of post-metallocene
catalysts.
from scandium to lanthanoid and actinoid metals, the ligands coordinate to the metal
atoms through oxygen, nitrogen, phosphorus, or sulfur atoms. As a rule, these complexes are well characterized by x-ray, nuclear magnetic resonance (NMR), and IR
methods [104–110].
The total number of different transition metal complexes that can be converted into
ethylene polymerization catalysts is quite large [110], and several catalysts prepared
from phenoxy-imine complexes (Figure 2.2) have already found commercial application [55, 111, 112]. Some of the catalysts with phosphinimine ligands were also tested
in multi-stage polymerization processes [60].
All post-metallocene complexes used for ethylene polymerization are soluble in
aromatic solvents, and some are soluble in aliphatic hydrocarbons. Similarly to metallocene complexes, post-metallocene complexes shown in Figures 2.2 and 2.3 are converted into polymerization catalysts by reacting them with MAO or with ion-forming
cocatalysts as described in Section 2.3.4.1.
2.3.6 Binary Transition Metal Catalysts
2.3.6.1 Binary Ziegler-Natta/Metallocene Catalysts
Mechanical properties of ethylene homopolymers produced with titanium-based
Ziegler-Natta catalysts significantly improve when their molecular weight increases.
Catalysts for the Manufacture of Polyethylene 47
However, polymers of a very high molecular weight are difficult to process into useful
articles (for example, into film) due to the high viscosity of their melts. The problem
of excessively high viscosity can be overcome if such a polymer is “diluted” with PE
of a relatively low molecular weight so that the average melt viscosity of the mixture
remains within the acceptably low range.
Unfortunately, mechanical blending of two PE resins in the molten state cannot produce such materials because polymer melts with vastly different viscosities do not mix well. An alternative route to the manufacture of such PE blends
involves the use of bicomponent metallocene/Ziegler-Natta catalysts. Every particle
in these catalysts contains active centers of two types, Ziegler-Natta and metallocene. Both catalyst components polymerize ethylene at the same time and under
the same reaction conditions. However, average molecular weights of the polymer material they form are very different: the metallocene active centers produce
a PE component with a low molecular weight and the Ziegler-Natta active centers produce a PE component with a high molecular weight. Blending of the two
polymer components occurs directly during their synthesis and, as a result, binary
metallocene/Ziegler-Natta catalysts form PE products with a bimodal molecular
weight distribution.
Combining two types of catalysts, one a titanium-based Ziegler-Natta catalyst and
another a metallocene catalyst, has proved to be a difficult task, mostly due to mutual
adverse effects of the catalyst components [113–114]. Nevertheless, several successful
bicomponent catalysts of this type were developed. They usually contain combinations
of supported titanium-based catalysts of the TiCl4/MgCl2/silica type and metallocene/
MAO/silica catalysts [55, 113–116]. The skill of catalyst chemists in manufacturing
such bicomponent catalysts is two-fold, (a) to choose a combination of the catalyst
components that produces PE mixtures of a desired average molecular weight, and (b)
to develop a technique for the catalyst synthesis that minimizes the mutual interference
of the catalyst components.
2.3.6.2 Binary Metallocene and Post-Metallocene Catalysts
Combining two metallocene complexes, or two post-metallocene complexes, or one
metallocene and one post-metallocene complex, can bring the same advantages as combining a Ziegler-Natta and a metallocene catalyst: copolymerization reactions of ethylene and an α-olefin in a single reactor with such dual catalysts produce blends of two
materials with different composition, different molecular weight and different molecular
weight distribution [55, 112, 117].
If both catalysts are soluble in a polymerization medium, another option opens: a
transfer (shuttling) of polymer chains between two catalyst species [118–121]. A special
shuttling agent, usually Zn(C2H5)2, separates a growing polymer chain from one active
center and attaches it to a center of a different nature. The transfer reactions are very
fast and occur several times during the time of growth of a single polymer molecule. In
this case, the copolymer product is not a mixture of two individual materials anymore;
it consists of macromolecules of a single type containing blocks of different composition [118–122].
48 Handbook of Industrial Polyethylene and Technology
2.4 Chemistry of Catalytic Polymerization Reactions
Polymerization reactions of ethylene and copolymerization reactions of ethylene and
α-olefins with all classes of transition metal catalysts discussed in Section 2.3 are catalytic
in the same sense as other catalytic reactions: a single active center containing a transition metal atom produces many polymer molecules, one after another, over a period of
time dictated by the polymerization technology, from several minutes in high-temperature solution processes to several hours in gas-phase and slurry processes. When one
describes the mechanism of the catalytic polymerization reactions, one should be aware
that two different chemical processes occur in parallel in the course of each such reaction.
The first of them involves formation and decay reactions of the active centers.
Active centers in supported chromium-based catalysts are formed when the catalysts
are contacted with ethylene in a polymerization reactor at an increased temperature.
Active centers in titanium-based Ziegler-Natta catalysts and in all metallocene and
post-metallocene catalysts are formed in reactions of respective catalysts with cocatalysts, when the original transition metal compounds are transformed into a variety of
secondary chemical compounds, some of which are the true catalytic centers. In both
cases, chromium– and titanium-based catalysts, the active centers are formed relatively
rapidly, within 10–20 minutes. In some catalysts, the formed active centers are very
stable and the polymerization rate remains steady for several hours. In other catalysts,
the active centers are quite unstable, especially at high temperatures, and the reaction
rate rapidly decreases with time.
The real chemical structure of the active centers was established only for metallocene catalysts. The active center in these catalysts is a metallocenium cation Cp2M+-R
containing the M+-C bond (M is Zr, Hf or Ti). The metallocenium cation is formed
in a reaction of a metallocene complex Cp2MCl2 with a cocatalyst (MAO) and with
Al(CH3)3 present in MAO. This sequence of reactions can be schematically represented
as follows [8, 9, 123]:
Cp2MCl2 + 2Al(CH3)3
Cp2M(CH3)2 + [Al(CH3)–O]m
Cp2M(CH3)2 + 2Al(CH3)2Cl
(2.32)
Cp2M+ CH3 + {[Al(CH3)–O]m–CH3}- (2.33)
If an ion-forming cocatalyst, such as a combination of B(C6F5)3 and Al(CH3)3, is used
instead of MAO, a different ion pair containing the metallocenium cation Cp2M+-CH3
is produced:
Cp2M(CH3)2 + B(C6F5)3
Cp2M+ CH3 + [(C6F5)3B–CH3]–
(2.34)
The structure of active centers in supported Ziegler-Natta and chromium-based
polymerization catalysts is not known definitely yet, but it is generally assumed that
these centers are also cationic and contain the [M]+-C bond. Here M is a Ti atom or a
Cr atom, the brackets [ ] signify that the nature of other ligands attached to M+ is not yet
definitely known, and R is an alkyl group, which can be either short (CH3, C2H5, etc.)
or larger (a polymer chain).
Catalysts for the Manufacture of Polyethylene 49
Chemical reactions of the second type, the polymerization reactions themselves,
start immediately after the active centers are formed. The principal chemical step in any
catalytic polymerization reaction of ethylene is the insertion of an ethylene molecule
into the [M]+-C bond:
[M]+ R + CH2=CH2
[M]+ CH2 CH2 R
(2.35)
In polymer chemistry, this reaction is called the chain growth (or the chain propagation) reaction. Each Reaction 2.35 is very fast, thousands of them occur over a period
of a few seconds. This repeated insertion reaction can be written in a general form as an
increase of the polymer chain length while the chain remains attached to the transition
metal atom M in the active center:
[M]+ (CH2 CH2)n R + CH2=CH2 → [M]+ (CH2 CH2)n+1 R
(2.36)
The R group in Reaction 2.36 is called the starting end-group of the polymer
chain. For example, the R = CH3 in the metallocene center formed in Reactions 2.33
or 2.34.
To produce LLDPE and VLDPE resins, ethylene is copolymerized with a small
quantity of an α-olefin. When an α-olefin molecule CH2=CH-R is present in a reactor,
another chain growth reaction can take place:
[M]+ (CH2 CH2)n R + CH2=CH R
[M]+ CH2 CHR (CH2 CH2)n R
(2.37)
Reactions 2.36 and 2.37 are the two principal steps in the formation of ethylene/αolefin copolymer chains. Ethylene is the monomer of the highest reactivity in catalytic
polymerization reactions and Reaction 2.36 always proceeds with a significantly higher
rate constant compared to Reaction 2.37.
Very infrequently, an ethylene molecule reacts with the growing polymer chain in a
different manner than in Reaction 2.36:
[M]+ (CH2 CH2)n R + CH2=CH2
[M]+ CH2 CH3 + CH2=CH (CH2 CH2)n 1 R
(2.38)
Reaction 2.38 is called the chain transfer reaction to a monomer (to ethylene in this
case). It leaves the vinyl CH2=CH– bond as the end-bond in the polymer chain. This
reaction results in disengaging of a polymer chain from the active center. The new center [M]+-CH2-CH3 still has the M+ C bond and it retains the ability to insert ethylene
molecules and to grow a new polymer chain. The reaction of chain transfer to ethylene
is faster when it immediately follows Reaction 2.37:
[M]+ CH2 CHR (CH2 CH2)n R + CH2=CH2
[M]+ CH2 CH3 + CH2=CR (CH2 CH2)n 1 R
(2.39)
50 Handbook of Industrial Polyethylene and Technology
Reaction 2.39 leaves a different type of double bond, CH2=CR , as the end-bond in
the polymer chain. Reactions 2.38 and 2.39 are the principal chain transfer reaction in
ethylene polymerization reactions with chromium oxide and metallocene catalysts. The
frequency of both these reactions strongly depends on reaction temperature: the higher
the temperature the more frequent are the two reactions and, as a result, the lower the
polymerization number n in the formed polymer molecules CH2=CH (CH2 CH2)n 1 R
or CH2=CR (CH2 CH2)n 1 R.
When titanium-based Ziegler-Natta catalysts are used for ethylene polymerization,
Reactions 2.38 and 2.39 occur quite rarely and the polymer chains contain, on average, a very large number of ethylene units; the n value can reach ten to fifty thousand. Polymers with such a high molecular weight are difficult to process; they found
commercial use only as components of HDPE resins produced with binary catalysts
(Section 2.3.5) or as ultrahigh molecular weight PE resins, which require special processing techniques and equipment. In order to reduce the molecular weight of ethylene
polymers prepared with Ziegler-Natta catalysts, a special chemical agent, hydrogen, is
added to the polymerization reactions. It hydrogenates the [Ti]+-C bond in the growing
polymer chain:
[Ti]+ (CH2 CH2)n R + H2
[Ti]+ H + CH3 CH2 (CH2 CH2)n 1 R (2.40)
Reaction 2.40 is called the chain transfer reaction to hydrogen. It produces saturated
polymer chains. The center with the [Ti]+-H bond, which is formed in Reaction 2.40,
remains catalytically active; it can insert the double bond of an ethylene molecule:
[Ti]+ H + CH2=CH2
[Ti]+ CH2 CH3
(2.41)
The polymerization process continues as in Reaction 2.36. Reaction 2.40 proceeds
at a much higher rate than Reaction 2.38, and the molecular weight of ethylene polymers produced with Ziegler-Natta catalysts in the presence of hydrogen is always
significantly lower. Active centers in metallocene catalysts also easily react with
hydrogen:
Cp2Zr+ (CH2 CH2)n R + H2
Cp2Zr+ H + CH3 CH2 (CH2 CH2)n 1 R
(2.42)
However, when chromium oxide catalysts are used to polymerize ethylene, Reaction
2.41 does not take place.
One more chemical reaction is essential for the understanding of mechanical
properties of linear PE resins. Polymer chains formed in Reaction 2.38, CH2=CH
(CH2 CH2)n-1-R, have the same vinyl double bond CH2=CH– as any α-olefin molecule CH2=CH R employed in an ethylene/α-olefin copolymerization reaction, such
as Reaction 2.37. The difference between these two α-olefin molecules is merely in the
size of the alkyl group attached to the vinyl double bond, this alkyl group is small in
α-olefin molecules usually used in the copolymerization reactions (R = C2H5 to C6H13)
but it can be very large in polymer molecules generated in Reaction 2.38. When ethylene
Catalysts for the Manufacture of Polyethylene 51
polymerization reactions are catalyzed by metallocene or chromium oxide catalysts,
such “dead” CH2=CH (CH2 CH2)n 1 R macromolecules can participate in copolymerization reactions with ethylene similarly to Reaction 2.37 [9, 65–66, 124, 125]. They produce long-chain branches in PE molecules. Although long-chain branches are always
present in such materials in a very low amount compared to LDPE resins produced in
radical polymerization reactions, they significantly affect rheology of the polymer melts
and mechanical properties of the resins.
2.5 Uniformity of Active Centers
In the field of chemical catalysis, the term “uniformity of active centers” refers to differences in the reactivity of different active centers in a given catalyst. Two manifestations
of active center uniformity (or their non-uniformity) can be readily observed in ethylene polymerization catalysts:
(a) Differences in average molecular weights of macromolecules they form.
(b) Differences in the ability of active centers to copolymerize ethylene with
-olefins.
2.5.1 Uniformity of Active Centers with Respect
to Molecular Weight of Polymers
The average molecular weight of macromolecules produced by a given active center in
an ethylene polymerization reaction is determined by the ratio of the rate of the chain
propagation reaction (Reaction 2.36) and the sum of the rates of two chain transfer
reactions (Reactions 2.38 and 2.40). If all active centers in a given catalyst have the same
kinetic properties, they produce polymer molecules with a certain type of molecular
weight distribution which is called “the normal distribution.” In general, the molecular weight distribution of a given polymer mixture is characterized by the ratio of two
average molecular weights, the weight-average and the number-average, Mw/Mn. The
normal molecular weight distribution is quite narrow, Mw/Mn = 2. If the Mw/Mn ratio of
a PE resin produced with a given catalyst is noticeably higher than 2, it means that the
active centers in the catalyst are not kinetically equivalent.
Ethylene polymerization catalysts provide examples of both types of center uniformity. Some of the catalysts contain active centers that are kinetically uniform and other
catalysts contain several types (populations) of active centers which significantly differ
one from others in terms of their kinetic features [9, 126]. For example, metallocene
catalysts based on zirconocene Complexes I (Figure 2.1) have kinetically uniform active
centers, metallocenium cations formed in Reaction 2.33. Ethylene polymers produced
with these catalysts are characterized by the Mw/Mn ratio close to 2 [126]. In contrast, all
chromium oxide catalysts and supported Ziegler-Natta catalysts contain several populations of active centers. Each of the centers produces PE molecules of a different average
molecular weight and, as a result, the Mw/Mn ratios for the total polymers are significantly
higher than 2. Table 2.6 gives one example of kinetic nonuniformity of active centers in a
chromium oxide catalyst. The Mw/Mn ratio for this HDPE resin is ~12 [127].
52 Handbook of Industrial Polyethylene and Technology
Table 2.6 Components of a PE homopolymer produced with different populations of active
centers in chromium oxide catalyst [127].
Population of active centers*
I
II
III
IV
V
VI
VII
Average molecular weight
3,200
12,100
38,500
97,000
228,000
600,000
1,850,000
Content of fraction, %
2.6
9.9
23.5
25.2
21.2
14.5
3.1
* Polymerization at 90 °C in a 2-hour reaction.
2.5.2 Uniformity of Active Centers with Respect
to Copolymerization Ability
A significant part of PE resins used in industry are ethylene/ -olefin copolymers with a
small content of -olefins. Compositional uniformity of these copolymers has emerged
as a very important factor which affects their mechanical and end-use properties.
Some ethylene/ -olefin copolymers have a high degree of compositional uniformity.
All macromolecules in such resins have approximately the same content of -olefin.
This type of compositional uniformity is typical for LLDPE and VLDPE resins prepared
with metallocene catalysts in solution polymerization reactions. LDPE resins produced
in radical polymerization processes at a high pressure in the presence of radical initiators, although they are manufactured from a single monomer, ethylene, can also be
viewed as ethylene copolymers with a mixture of various -olefins (see Section 2.2.2).
These materials are also mostly compositionally uniform: any such polymer molecule,
taken at random, has approximately the same number of short-chain branches of the
same types.
In contrast, ethylene/ -olefin copolymers produced with supported Ziegler-Natta
catalysts and with chromium oxide catalysts have a pronounced compositional nonuniformity. Such copolymers, even when they are produced in copolymerization reactions under strictly controlled stable reaction conditions, are mixtures of copolymer
molecules with very different compositions. Some of the molecules in these mixtures
contain very few -olefin units and, to a first approximation, can be viewed as nearly
linear ethylene homopolymers, while other copolymer molecules contain a relatively
large number of -olefin units. For example, if a copolymer is produced with a typical
Ziegler-Natta catalyst and contains 3.0 mol% of an -olefin (a typical composition of a
commercial LLDPE resin), it is in reality a mixture of macromolecules with the -olefin
content ranging from below 0.3 mol% in its high molecular weight fraction to over
20 mol% in the fraction with the lowest molecular weight. This phenomenon can be
demonstrated by fractionating ethylene polymers with a series of solvents, as shown in
Table 2.7 for two ethylene plastomers.
Obviously, the average contents of -olefins in the plastomers in Table 2.7 (10.3 mol%
in the ethylene/1-butene resin and 7.2 mol% in the ethylene/1-hexene resin) are the
Catalysts for the Manufacture of Polyethylene 53
Table 2.7 Fractionation of ethylene/ a-olefin plastomers [8].
Ethylene/1-butene plastomer*
Total
Hot hexane Hot hexane Cold hexane Hot ether** Cold ether**
resin
insoluble
soluble
soluble
soluble
soluble
Fraction, %
100
29.0
32.0
13.6
7.0
18.4
Butene, mol% 10.3
3.2
6.4
12.6
19.8
30.0
Ethylene/1-hexene plastomer*
Total
Hot octane Hot octane Hot heptane Hot hexane Cold hexane
resin
insoluble
soluble
soluble
soluble
soluble
Fraction, %
100
49.0
5.0
19.1
7.7
19.3
Hexene,
mol%
7.2
1.1
2.1
4.4
9.1
~28
* produced with supported Ziegler-Natta catalysts in slurry polymerization reactions
** ether = diisopropyl ether
averages of very broad compositional distributions ranging from fractions with a quite
low -olefin content (these fractions are highly crystalline) to completely amorphous
fractions containing about 30 mol% of the -olefin.
The reason for the formation of copolymer molecules of different composition in
a single monomer mixture is well established: all titanium– or chromium-based supported catalysts contain several populations of active centers with different copolymerization ability [6–7, 9, 127].
Thus, in general, principal differences between different formulations of ZieglerNatta catalysts (Section 2.3.3) and chromium oxide catalysts (Section 2.3.2) are related
to the following three features:
a) The catalysts differ one from another in activity; i.e., in the number of
active centers per weight unit of catalyst.
b) The catalysts differ in the content of populations of active centers that
produce polymer molecules with different average molecular weights
(Section 2.5.1).
c) The catalysts differ in the content of populations of active centers which
have different copolymerization ability.
At the present time, two complimentary fractionation techniques, the analytical temperature-rising elution fractionation (analytical Tref) and the crystallization fractionation (Crystaf) [9], provide the most detailed analysis of compositional uniformity of
ethylene/ -olefin copolymers. Both techniques exploit the same principle: copolymer
molecules of a different composition crystallize from solution at different temperatures.
Compositional uniformity of ethylene/ -olefin copolymers can be also evaluated rapidly, although on a semi-quantitative level, using the differential scanning calorimetry
(DSC) method. If an -olefin content in a copolymer is low, a few percent, all its molecular chains contain long sequences of ethylene units. These linear ethylene blocks form
PE crystals, the same as in ethylene homopolymers. If a catalyst has uniform active
54 Handbook of Industrial Polyethylene and Technology
centers with respect to their copolymerization ability, all copolymer molecules it produces have approximately the same composition. Even if a copolymer of this type has
an -olefin content of merely 2 to 3 mol%, the average length of an ethylene block in its
macromolecules is relatively short and these blocks form thin crystalline lamellae that
melt at a relatively low temperature in comparison to the melting point of strictly linear
PE. As a result, the melting behavior of compositionally uniform copolymers shows a
strong dependence on composition: their melting point decreases from 135 to 140 °C
for the ethylene homopolymer to ~120 °C for copolymers containing 1.5 to 2 mol% of
an -olefin and to ~110 °C for copolymers containing 3.0 to 3.5 mol% of an -olefin.
If a catalyst has nonuniform active centers with respect to their copolymerization
ability, the copolymer materials it produces have a pronounced compositional nonuniformity; i.e., they are mixtures containing copolymer molecules of different composition, from practically linear macromolecules to macromolecules with a quite high
-olefin content (see Table 2.7). Melting of such copolymer mixtures is dominated by
their nearly linear fractions, which are highly crystalline. As a result, melting points of
compositionally nonuniform PE resins are not very sensitive to the average copolymer
composition and are usually in the range of 125 to 128 °C [9, 33, 128].
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Catalysts for the Manufacture of Polyethylene 57
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Catalysts for the Manufacture of Polyethylene 59
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3
Ethylene Polymerization Processes and
Manufacture of Polyethylene
Ian D. Burdett and Ronald S. Eisinger*
Mid-Atlantic Technology, Research and Innovation Center
(MATRIC), South Charleston, West Virginia, USA
Contents
3.1 Introduction ................................................................................................................................. 62
3.1.1 Magnitude of the PE Industry ....................................................................................... 63
3.1.2 Active Processes............................................................................................................... 63
3.1.3 Range of Products ........................................................................................................... 64
3.1.4 Chronology of Development of Processes ................................................................... 65
3.2 Processes....................................................................................................................................... 65
3.2.1 Common Principles of Ethylene Polymerization at Commercial Scale ................... 65
3.2.2 High-Pressure Process Technology............................................................................... 67
3.2.2.1 Autoclave Reactors............................................................................................ 69
3.2.2.2 Tubular Reactor ................................................................................................. 70
3.2.2.3 Operational Considerations and Safety ......................................................... 71
3.2.2.4 Current Operation ............................................................................................ 73
3.2.3 Gas-Phase Fluidized Bed Reactors................................................................................ 73
3.2.3.1 Process Description .......................................................................................... 74
3.2.3.2 Product Capabilities ......................................................................................... 76
3.2.3.3 Process Challenges ............................................................................................ 76
3.2.4 Slurry Reactors ................................................................................................................ 77
3.2.4.1 Slurry Loop Reactors ........................................................................................ 77
3.2.4.2 Slurry CSTR Reactors....................................................................................... 83
3.2.5 Solution Reactors............................................................................................................. 87
3.2.5.1 The Dow Chemical Company’s DOWLEX Process ..................................... 87
3.2.5.2 The SCLAIRTECH Process ............................................................................. 90
3.2.5.3 The Stamicarbon Compact Process ................................................................ 92
3.2.5.4 Equistar Solution Process ................................................................................ 93
3.2.6 Hybrid Processes ............................................................................................................. 93
3.2.6.1 Spherilene Hybrid Process ............................................................................... 93
3.2.6.2 Borstar Hybrid Process .................................................................................... 94
*Corresponding author: ron.eisinger@matricinnovates.com
Mark A. Spalding and Ananda M. Chatterjee (eds.) Handbook of Industrial Polyethylene and Technology, (61–104)
© 2018 Scrivener Publishing LLC
61
62
Handbook of Industrial Polyethylene and Technology
3.3 Resin Property and Reactor Control in Catalytic Polymerization Reactors ....................... 95
3.3.1 Production Rate ............................................................................................................... 95
3.3.2 Catalyst Productivity....................................................................................................... 96
3.3.3 Reactor Pressure .............................................................................................................. 96
3.3.4 Crystallinity...................................................................................................................... 96
3.3.4.1 Ziegler-Natta and Metallocene Catalysts ....................................................... 96
3.3.5 Molecular Weight ............................................................................................................ 96
3.4 Economics .................................................................................................................................... 97
References ............................................................................................................................................ 100
Abstract
Over the past 60 years, applications for polyethylene (PE), catalysts for producing it, and ever
more efficient processes for manufacturing it have blossomed and continue to expand. Today,
this polymer is the most versatile and widely used plastic in the world. This chapter addresses
the manufacturing processes used to produce the major types of PE: high-pressure low density
PE (LDPE), linear low density PE (LLDPE) and high density PE (HDPE). Emphasis is placed on
descriptions of current processes from an engineering viewpoint. The oldest processes, which
operate at high pressure and usually do not employ catalyst, are described first. Discussion of
low-pressure catalytic processes begins with the most widely used process, the gas-phase, fluidized-bed process. Then, those HDPE processes which are carried out in a slurry of PE solids in
a diluent are examined. Solution-based processes for production of LLDPE, which evolved from
the HDPE processes, are then discussed. Recent hybrid processes are also described.
Keywords: Manufacturing processes, slurry, loop, fluidized bed, solution, LDPE, LLDPE,
HDPE
3.1 Introduction
With the possible exception of computers, plastics have been the greatest industrial
advancement of the twentieth century. The introduction of plastics has provided materials with an extraordinary range of properties. Plastic products can be rigid and strong,
flimsy as a thin film, or of any intermediate property. Plastics can be molded, blown into
desired shape, rolled into sheets of specified thickness, or spun into filaments. Products
can be opaque or crystal clear. Many plastics are resistant to strong acids and bases, and
to gasoline and many other organic compounds. As a vital part of the plastics revolution, researchers have furthered the understanding of how to design polymer molecules
to the point where a given plastic can meet numerous specifications. A major part of
this revolution has been the advances made in producing PE, the plastic most widely
used in the world.
This chapter focuses on the processes that have been developed to produce PE over
the past 75 years. In particular, processes used today are discussed. A brief history is
included to understand how processes have evolved. Catalysts are often used to make
PE, but this chapter will not focus on the chemistry. The chemistry of polymerization is
described in Chapter 2. That said, process development engineers understand that the
most important driving force in designing low-pressure processes for PE is the catalyst.
And for high-pressure low density PE processes, it is initiator technology.
Ethylene Polymerization Processes and Manufacture of Polyethylene
63
3.1.1 Magnitude of the PE Industry
Polyethylene is the world’s largest volume plastic. Worldwide production has been variously reported from a high of 96 million tonnes (metric tons) in 2010 [1] to as low as 82
million tonnes in 2013 [2]. This volume is about 37% of all polymer production. The PE
sales in North America in 2013 amounted to 17.5 million tonnes [3]. Commodity PE is
divided into three types: LDPE, LLDPE, and HDPE. Production rates in 2013 for these
types are shown in Table 3.1. The rate of increase in production now follows the rate of
increase of the population. Growth in HDPE and LLDPE is about three times higher
than that for LDPE. Most new production in the past 20 years has been in third-world
countries, especially in Asia.
3.1.2 Active Processes
Between the 1930s, when PE was first synthesized, to the present time, numerous processes have been devised to produce this versatile polymer with improved properties
and at less expense. These production processes are divided into the broad categories
of high-pressure and low-pressure. High-pressure processes operate at pressures ranging from roughly 400 to 3400 bar. Ordinarily, one or more free-radical initiators, such
as oxygen or organic peroxides, are fed to the reactor, but no catalyst is needed. Both
tubular and autoclave reactors are employed. Even though commercial processes were
operational as early as 1939, numerous producers of PE continue to use high-pressure
processes today.
Low-pressure processes, which invariably use catalysts, were first commercialized in
the late 1950s. Catalytic processes were able to make linear chains of PE, in contrast to
highly branched, high-pressure LDPE resins. The earliest catalysts were based on supported chromium oxide. Surprisingly, these catalysts remain among the most widely
used PE catalysts today. Somewhat later, catalysts were categorized as coordination catalysts, most notably Ziegler-Natta (Z-N) catalysts. The active sites of these catalysts are
titanium or vanadium. Titanium is used in the form of TiCl3 or TiCl4, often mixed with
MgCl2, while vanadium is either oxygenated or chlorinated. In either case, a reducing
cocatalyst, typically an aluminum alkyl, accompanies the active species. Beginning in
the 1990s an offshoot of Z-N catalysts called metallocene catalysts were commercialized. These catalysts classically contain a dicyclopentadiene moiety which contains an
active metal. The metal is usually zirconium or titanium, or can be hafnium. These catalysts also require a cocatalyst, the earliest of which was methylaluminoxane (MAO), a
Table 3.1 Sales of PE in North America in 2013 [3].
PE type
Sales in North America, 2013 millions of tonnes
LDPE
3.1
LLDPE
6.3
HDPE
8.1
Total
17.5
64
Handbook of Industrial Polyethylene and Technology
hydrated form of aluminum alkyl. Around the same time, a related class of catalysts
called single-site or constrained-geometry catalysts was introduced commercially.
Low-pressure processes operate at a pressure range of 5 to 40 bar. The numerous processes that operate at low pressure are grouped into three types. The first type is a slurry
process in which polymerization occurs in a hydrocarbon diluent in which the polymer
will not dissolve. This process makes HDPE. The earliest of these processes, the Phillips
process, was introduced in the late 1950s. Originally based on the chromium catalyst,
the Phillips process has now been adapted to use other catalysts as well. It remains the
benchmark among slurry processes. Later, processes were developed that could make
LLDPE. Because it can dissolve in some organic solvents, it became possible to produce
LLDPE in a homogeneous solution. These solution processes, which operate at higher
temperatures and shorter residence times, constitute the second type of catalytic processes. DuPont and The Dow Chemical Company were among the earliest practitioners
of the solution process. The solution process generally cannot make HDPE. The third
type of catalytic process produces PE in the gas phase. The most notable incarnation of
the gas-phase process is the fluidized-bed process commercialized by Union Carbide.
Both HDPE and LLDPE can be made in gasphase processes.
3.1.3
Range of Products
The ability to produce PE over a wide range of crystallinity and molecular weight (MW)
enables its use in thousands of different applications. The range of resin density, a measure of crystallinity, is typically 0.900 to 0.965 g/cm3. At typical MWs, these density
values correspond to crystallinities of 35 to 85% [4]. Low-volume specialty polymers of
very-low- or ultra-low-density PE have resin densities as low as 0.890 g/cm3 and even
lower. Even so, these products still retain some crystallinity.
Molecular weight correlates with viscosity of the polymer melt. Therefore, MW is
inferred by measurement of melt viscosity, which is obtained by measuring the time
required for a given amount of the polymer melt to be forced through an orifice. A
common set of conditions are used. For example, the I2 melt index (MI) is measured
at 190 °C and a force of 2.16 kg. The range of I2 melt index in commodity PE usually
ranges from 0.1 to 100 dg/min. This range is roughly equivalent to molecular weights of
hundreds of thousands down to 5,000 Daltons. PE with an I2 melt index as low as 0.01
is produced [5]. Ultra-high-MW PE, i.e., above one million Daltons, had previously
not been made on a large scale because it is difficult to process the high-viscosity melt
without degrading the polymer. However, such polymers may have desirable end-use
properties, especially for film. It has been known that PE having a very broad molecular weight distribution (MWD) can be easier to process. During the last two decades,
catalyst advances and process developments have enabled unprecedented control of
MWD. The use of two or more reactors in series can produce polymer with a bimodal
or multimodal MWD. For reasons of economy, however, it would be desirable to create
a very broad MWD in a single reactor. That capability is valuable and has become the
focus of recent catalyst and process developments.
A brief overview of applications of PE will illustrate its omnipresence worldwide. Its
first key use as electrical insulation for radar cable was critical for the British in World
War II. The largest use is in film for such purposes as packaging, storage bags, garbage
Ethylene Polymerization Processes and Manufacture of Polyethylene
65
Table 3.2 Chronology of major PE process developments.
Company
Process name
Process type
Product
Year
Imperial Chemical
Industries (ICI)
High-pressure autoclave
LDPE
1939
BASF
High-pressure tubular
LDPE
1940s
DuPont Canada
SCLAIR
Solution
LLDPE
1960
Phillips Petroleum (now
Chevron-Phillips
Chemical)
Phillips (now
MarTECH)
Slurry loop
HDPE
1961
Union Carbide
UNIPOL
Gas-phase fluidized bed
HDPE
1968
BASF
Hostalen
Slurry tank
HDPE
1975
Union Carbide
UNIPOL
Gas-phase fluidized bed
LLDPE
1977
bags, wraps, and greenhouse applications. It is employed in pipe for natural gas and
water. Geomembranes (sheet) are used to line earthen structures. Toys, beginning with
hula hoops, are a common application. A major market is household and industrial
containers to hold foods, household cleaners, detergents, motor oil, and gasoline. Pails,
drums, and crates are important products. Foams are used in headgear. Coatings are a
widespread use. Diapers and fishing nets are other important applications.
3.1.4 Chronology of Development of Processes
The development of major processes to manufacture PE is summarized in Table 3.2.
In recent years, advances have been primarily in catalysts and in the use of reactors in
series to make products with bimodal and multimodal MWDs. Lesser advances include
the means to increase production rate by increasing cooling and to recycle fluids more
effectively.
3.2
Processes
3.2.1 Common Principles of Ethylene
Polymerization at Commercial Scale
A number of engineering issues are applicable to all ethylene polymerization processes.
These issues are introduced here; most of them will be addressed in more detail in the
discussions of the individual processes.
Safety comes first. While major accidents in the PE industry are infrequent, the
capacity for a catastrophe with loss of life and widespread destruction within the plant
is present. All processes employ at least moderately high pressures and temperatures.
All plants use highly flammable monomers and many also use flammable diluents.
Decomposition of ethylene during a runaway reaction in high-pressure reactors is
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Handbook of Industrial Polyethylene and Technology
well known. Reaction rates increase with temperature for some catalysts used in lowpressure processes, which may lead to runaway reactions. Aluminum alkyl cocatalysts
used with Ziegler-Natta and some metallocene catalysts are pyrophoric and have been
a frequent source of fires. Catalyst components may be toxic. For example, chromium
trioxide, the basis for chrome catalysts, is carcinogenic when the chromium is in the +6
oxidation state. A constant state of safety awareness is necessary to maintain a safe plant.
The most important unit of operation in a polymerization unit is removal of the
heat of reaction. Ethylene polymerization is highly exothermic. The heat of polymerization of ethylene is variously reported as 3350 J/g [6, 7] and 3600 J/g [8]. The heats
of polymerization of comonomers, which are normally present at less than 10 mol%,
are substantially lower, probably no more than 2000 J/g. Extraordinary efforts are put
into the design of large-scale plants in order to remove the heat of polymerization.
Methods employed include: 1) heat exchangers on a recycle line, 2) boiling reactors in
which the vapor is removed, cooled, and condensed before returning it to the reactor, 3)
cooling jackets around a tubular reactor, and 4) the use of cold or even chilled ethylene
feeds. Because cooling water is ordinarily used, reactor temperatures above 70 °C are
preferred.
Fouling is the next most important design issue. Streamlined designs of the reactors and any heat exchangers are important. Frequently, there is a minimum velocity
at which fluids must move to minimize fouling. For LDPE high-pressure technologies,
control of axial and radial temperature profiles are critical. In low-pressure processes,
the kinetic profile of the catalyst has a significant effect on fouling, and is often a major
determinant of the catalyst’s viability. Adequate mixing is not only critical for making
homogeneous products, but is also important for avoiding fouling in stagnant areas.
Advances in commercial polymerization processes are typically characterized by incremental, proprietary improvements in antifouling technologies.
Conversion per pass of ethylene can exceed 95% in some processes, but is never
adequate to avoid the need for recycling of monomers and/or diluent for an economical process. In liquid phase processes, including slurry processes, substantial separation
steps are necessary to recycle monomers and diluents.
Feed stream purification is common to all processes, and especially to low-pressure,
catalyzed processes. Impurities may poison the catalyst or react with the cocatalyst,
reducing catalyst productivity. Impurities also can affect reactor fouling. Further, they
can change morphology of the polymer in subtle ways as well as influencing rheological properties. The impurities that must be removed depend on feed stream quality,
type of polymerization catalyst, and sometimes process type. In catalyzed processes,
the most common poisons in monomers and diluent are water and oxygen. Other poisons include carbon monoxide, carbon dioxide, methanol, other oxygenated hydrocarbons, sulfur compounds, and acetylene. Infrequently, ammonia, arsine, and phosphine
may be encountered. The tolerance level toward each poison varies with the catalyst.
Strong poisons usually need to be below 1 ppm (weight basis), and often below 0.2 ppm.
Adsorbents and catalysts used to remove impurities include molecular sieves, alumina, zeolite-modified alumina, and highly dispersed copper- and zinc-based materials. Manufacturers of these materials include BASF, Sud-Chemie (owned by Clariant),
UOP (owned by Honeywell), and Axens. The development of ultra-high-activity catalysts would not have been possible without advances in purification materials.
Ethylene Polymerization Processes and Manufacture of Polyethylene
67
The form of the catalyst or activator determines how it can be fed to the reactor.
Gaseous activators and liquid catalysts are the easiest to feed. However, catalysts often
are impregnated onto a solid support. For example, it was found that supported Z-N
catalyst is much more productive in slurry reactors [1]. Feeding of dry supported
catalyst is more difficult. Such catalysts are more easily fed as a slurry in an appropriate
liquid medium.
In modern production facilities, catalyst productivity is usually high enough that
removal of the catalyst residue from the polymer is unnecessary. However, stabilization
additives including antioxidants are generally added to the polymer in a post-reaction
step, in part to prevent the residue from affecting polymer properties.
Devolatilization of the product is usually necessary for three reasons. First, the
buildup of desorbed hydrocarbons during storage or shipment in air can create a flammable atmosphere. Second, some customer applications, especially those involving
food contact, do not allow for odor in the polymer. Third, monomer and diluent may
be recovered and recycled. Devolatilization of high-pressure LDPE that does not contain comonomer employs a simple degassing step. In low-pressure catalytic processes,
devolatilization is more complicated and may contain more than one step. This operation is often carried out on granular polymer before it is melt-mixed and pelletized
because desorption is much more rapid in small particles. Steaming and gas purging
are frequently used for devolatilization. Gas purging employs continuous countercurrent purging of the granular polymer in a column, usually using nitrogen as the purge
gas. Purging of LLDPE and slurry HDPE granular products require up to several hours
of residence time because of deviation from plug flow of the particles and of the purge
gas in the column.
Nearly all PE is sold in pellet form. The added cost of pelletizing is considerable.
Efforts by Union Carbide in the 1980s to reduce costs by selling the granular reactor
product directly were unsuccessful because customers did not want to change their
handling facilities and because uniform mixing of stabilization additives into the product was more difficult.
3.2.2 High-Pressure Process Technology
Polyethylene was first discovered during experiments with ethylene under high pressure by ICI scientists in England in 1933. The cause was determined to be a free radical
polymerization initiated by trace levels of oxygen. ICI developed a commercial process by 1939 using an autoclave reactor operating between 200 to 300 °C and 1000 to
3000 bar. The product had immediate impact in the rapidly growing need for electrical
insulation on wire such as for newly discovered radar systems. The first commercial
tubular reactor system was developed by BASF during the early 1940s, producing products similar to but not identical to products using the autoclave system.
These two types of reaction systems are still used today but the technology has
advanced significantly through many developments [9, 10]. In particular in the 1970s,
process advances were achieved in reliability, process control, safety, and the range
of operating conditions. Coupled with chemistry advances in reaction initiators and
polymer chain length regulators, these advances improved product quality and the
range of possible products. The development of a variety of organic peroxides as
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Handbook of Industrial Polyethylene and Technology
reaction initiators for use in specific reaction zone conditions enabled higher conversion rates with stable process control. In recent years as computing capabilities have
dramatically advanced, improvements in automation, process simulation, and scaleup have resulted in much larger single-reaction train plants, particularly in tubular
systems. On the mechanical side, the development of high-pressure steel for use in
reactors has increased operating pressures, reduced capital costs, and broadened the
product range [11].
The nature of the PE product depends on process design and reaction conditions.
PE is a statistical product, not a uniquely chemically defined product. The statistical
properties, such as average MW and amount, length, and form of chain branching,
determine the rheology and other properties of the polymer and hence its applicability
to a specific commercial use. These statistical properties can be modified significantly
by process configuration, catalyst, process additives, and basic process conditions, such
as temperature and pressure, in the various sections of the process. In low-pressure processes, catalyst type, and its formulation and preparation are critical in achieving the
desired polymer properties. For high-pressure processes using free-radical polymerization, reaction initiators with chain transfer agents are used rather than catalyst. As a
result, high-pressure processes produce polymers which are structurally very different
from those produced using low-pressure processes. Also, as described later, the process
differences between an autoclave reactor and a tubular reactor impact product properties achievable in each system.
Within a particular high-pressure process configuration, product properties are controlled by adjusting a process condition variable based on its ability to modify a product property and on the ease of control of that variable while having minimal impact
on the operation of the process [10, 12]. The effect of the primary process parameter
on polymer properties is the same between autoclave and tubular reactors. However,
different reactor designs result in different product control strategies such as the ability
in a tubular reactor to modify the temperature profile along the length of the reactor.
One key property is crystallinity, which is inferred from the measurement of polymer
density. A decrease in reactor peak temperature or an increase in pressure decreases
short-chain branching and increases polymer density and crystallinity. The incorporation of a comonomer can be used to reduce density and alter other polymer properties.
The average molecular weight, which correlates with the viscosity measurement MI,
is controlled primarily by the concentration of an added chain transfer agent. Longchain branching formed by intermolecular chain transfer can be increased by raising
reactor temperature, decreasing pressure, increasing reaction initiator concentration,
and lengthening polymer residence time in the reaction system. In high-pressure processes, the breadth of molecular weight distribution (MWD) is most dependent on
reactor type because of differences in polymer residence time distribution and to a
lesser amount by reaction process conditions. It can be broadened to some extent by
the use of lower reaction temperatures and narrowed by higher reaction temperatures
and higher reaction initiator concentrations. Companies with advanced technologies
have sophisticated process and product response mathematical models for the above
relationships, which can be used with laboratory data feedback in computerized process control systems to maintain stable control of process conditions and product properties [9, 13, 14].
Ethylene Polymerization Processes and Manufacture of Polyethylene
69
3.2.2.1 Autoclave Reactors
The initial commercialization of PE manufacturing employed a process using a highpressure autoclave reactor. Although initially limited in reactor production capacity by
the rate of adiabatic heat removal to about 30,000 tonnes a year, it was very successful.
Current technology has now achieved single-train systems up to 150,000 tonnes a year,
operating at 1200 to 3000 bar and 150 to 315 °C with average residence times of 20 to
40 seconds. The capability of these autoclave systems to produce broader molecular
weight products and higher content ethyl vinyl acetate copolymers (up to 40 wt% vinyl
acetate) not possible in tubular reactor systems has kept existing units operating. The
slightly lower operating pressures and temperatures in these autoclave reaction systems
and the large distribution in residence times of growing polymer chains in the reaction
system because of the behavior of continuous stirred tank reactors (CSTR) result in a
polymer with a broader MWD and higher level of long-chain branching than in the
tubular system. This can be seen clearly in Figure 3.1 for an extrusion coating grade
made in either reaction type [15].
In the autoclave process for LDPE [10, 16–20], as shown in Figure 3.2, the purified ethylene is first compressed to about 250 bar in a primary compressor, cooled and
then compressed further in a secondary compressor to the required reactor pressure
of 1200 to 2000 bar. The ethylene feed for the secondary compressor includes recycled
unreacted ethylene which was discharged from the reactor and separated from the molten polymer. Discharge from the secondary compressor is cooled and fed to the reaction system along with peroxide compounds which are used as reaction initiators. This
reactor is a CSTR and in modern systems can have multiple CSTR stages. It can be
modeled as several CSTRs with connecting plug-flow reactor sections. The single-stage
reactor has an aspect ratio (length/diameter, L/D) of about 4. It is capable of producing broader MWD LDPE and can form more long-chain branches in the polymer. The
multistage system can have 3 to 4 CSTRs in series with an overall L/D of about 20 [16].
These stages are controlled at different temperatures, feed concentrations, peroxide
initiators types, and ethylene feed rates. The autoclave reactor is adiabatic and therefore only removes heat from the discharge of the hot molten polymer and unreacted
0.7
Molecular weight distribution
0.6
Tubular
Frequency
0.5
Autoclave
0.4
0.3
0.2
0.1
0
2
3
4
5
6
7
8
Log molecular weight
Figure 3.1 Difference in LDPE molecular weight distribution between tubular and autoclave highpressure reactors [15].
70
Handbook of Industrial Polyethylene and Technology
Primary
compressor
250 bar
Ethylene
Secondary 1400 bar
compressor
HX
HX
Autoclave
reactor
170 °C
240 °C
HX
Peroxide
initiators
280 °C
Wax
separator
Oil
separator
HX
Low pressure
separator
0.5 bar
250 °C
HX
280 bar
250 °C
HX
High pressure
separator
Degassing
Extruder
Product to bulk
loading
Figure 3.2 High-pressure autoclave process for the production of LDPE (adapted from [18]).
ethylene. Thus, ethylene conversion is limited to about 20% as compared to 35% in a
tubular reaction system with wall cooling. However, the autoclave reactor capacity can
be increased by refrigeration of the inlet ethylene. At the reactor outlet, the reaction
fluid is passed through a letdown valve and cooled before entering a high-pressure separator operating at about 250 °C and 250 to 1000 bar. This step quenches the reaction
quickly and prevents formation of high MW polymer which can cause gel contamination in film grade products. Phase-separated ethylene gas from the high-pressure separator is fed back to the inlet of the secondary compressor after further cooling and flash
separation from low MW waxes. The pressure of the molten polymer is further reduced
in a low-pressure separator operating at less than 1 bar. Released ethylene is recycled
back to the primary compressor inlet after cooling and separation from the low MW
oils. The molten polymer is discharged directly into an extruder and pelletizer system.
Additives are also fed into the extruder feed for the purposes of improving polymer
processing during fabrication, polymer stabilization, and sometimes to modify enduse polymer properties. Solid additives can be pre-mixed or fed as individual streams.
Liquid additives are normally directly fed. The pellets are cooled, dried, and conveyed
to polymer silos for blending and storage. These silos are purged with gas to remove
residual ethylene in order to avoid gas-phase explosive mixtures during transportation
or in customer facilities. It should be noted that LDPE film-grade products with high
clarity and low haze can be made by operating autoclave systems at lower pressure
(about 1200 bar). At these conditions, phase separation can occur in the reactor. The
polymer has the desired low amount of long-chain branching.
3.2.2.2 Tubular Reactor
The tubular reaction process for LDPE is very different from the reactor concept in the
previously discussed autoclave reactor technology [8, 10, 12, 18–21]. The flow pattern,
Ethylene Polymerization Processes and Manufacture of Polyethylene
Low pressure recycle
O2
initiator
Primary
compressor
HX
High-pressure recycle
Secondary
compressor
71
HX
Preheater
2400–2900 bar
HX
Reaction zone 1
Purge
Fresh
ethylene
HX
Reaction zone 2
Modifier
Precooler
180–310 °C
HX
Highpressure
separator
Low-
Degassing
Extruder
Product to
bulk loading
280 bar pressure
250 °C
separator
0.5 bar
250 °C
Figure 3.3 High-pressure tubular process for the production of LDPE (adapted from [18]).
as indicated by the process flow diagram in Figure 3.3, is closer to plug flow. Therefore,
the residence time distribution of forming polymers is narrower, resulting in a narrower MWD. Also, the use of a long length of pipe which is jacketed for heat removal
increases the single reactor production capacity compared to an adiabatic autoclave
reactor system. For commercial systems, the reactor consists of long sections of highpressure piping operating in the range of 1500 to 3500 bar. At the point of initial feed
of the reaction initiator component, the reactor temperature is about 150 °C but can
increase to a maximum temperature of 330 °C. Different temperature profiles can be
created in each section of the reactor by varying the feed rates of ethylene, reaction initiator, and chain transfer agents. Figure 3.4 shows the temperature profiles for up to four
feed locations. Different initiator types are also used and are selected based on temperature range required and reaction effectiveness in each specific pipe section. The pipe
is arranged as folded loops with a total length between 1 to 3 km and with an internal
diameter of the piping of 3 to 9 cm. The diameter can be varied along reactor length for
pressure gradient control. In addition, it is normally preferred to keep a relatively steady
concentration of chain transfer agent along the length of the reactor. These process
capabilities in modern tubular reaction systems provide greater control of the polymer
properties and greater stability of operation. Reaction heat is removed by the inlet feeds
of the cooled ethylene and the water-cooled jackets on the reactor piping. The primary
feed ethylene and compressor system and post-reactor systems are similar to autoclave
systems. However, ExxonMobil refers to a variation in their technology in which a portion of the feed ethylene after the primary compressor bypasses the reactor and is fed to
the high-pressure separator to cool and quench the reaction. This unreacted ethylene is
recycled back to the secondary compressor for feeding to the reactor inlet points.
3.2.2.3 Operational Considerations and Safety
A key objective in the operation of the reaction system is to keep the polymer products specification during steady-state operations. Another objective is to limit off-grade
Handbook of Industrial Polyethylene and Technology
300
200
Both #1
feeds
100
Temperature, °C
Temperature, °C
72
Single initiator
and ethylene
feed at inlet
Temperature, °C
Reactor length
300
200
Begin
both #1
feeds
Initiator
#2 feed
100
One ethylene
and two
initiator feed
locations
Reactor length
300
200
100
Both #1
feeds
Initiator and
ethylene
feed #2
Initiator and
ethylene
feed #3
Three initiator
and ethylene
feed locations
Temperature, °C
Reactor length
300
Ethylene #2
feed
Ethylene #3
feed
Both #1
200
feeds
100
Initiator Initiator
#2 feed #3 feed
Four ethylene and
initiator feed locations
Initiator
#4 feed
Reactor length
Figure 3.4 Temperature profiles for different feed locations of initiator and ethylene in the LDPE tubular
reaction system: 1) profile for a reactor with a single feed, 2) profile for two feeds, 3) profile for three
feeds, and 4) the profile for a reactor with four feeds.
material during transitions between products. In addition, process conditions must be
avoided in which contaminants like high MW gels are made which foul the reaction
system, or in the worst case cause runaway ethylene reactions, resulting in reactor shutdowns [8, 10, 13, 14, 16, 18, 20, 21]. Typical causes of process instability are fouling of
the walls and accumulation of impurities from recycling of the ethylene-rich stream.
While reaction times are short, reaction system changes can occur over the time period
of production of a specific grade of polymer. This means that adjustments need to be
made in initiators, ethylene feeds, and for tubular systems in the wall cooling. Typical
examples of operational problems are high temperature excursions which can cause
excessive long-chain branching and unwanted side reactions of chain transfer agents
and of organic impurities in ethylene and the initiators. The impurities can cause problems such as odor and higher haze levels in the final product and reduce the effectiveness of the peroxide initiators. In tubular reactors, temperature excursions are more
likely in the radial and axial directions due to poor control of the initiator concentration. Also, reactors with large tubular internal diameters for increased rates have higher
wall temperatures and larger radial temperature gradients which can cause product
variations. Accumulation of impurities, such as low MW poly-oils, requires adjustment
of the venting rates on recycle streams.
Polymer fouling of walls can cause cross-linking and high MW gels in the attached
polymer. Poor control of temperature and pressure can ultimately lead to conditions
under which ethylene spontaneously detonates. If localized, detonation can result in
Ethylene Polymerization Processes and Manufacture of Polyethylene
73
fine carbon particles in the finished product. In the extreme, a runaway decomposition can form carbon, hydrogen, and methane, causing a sharp increase in pressure
and temperature. This condition must be relieved immediately by partially or totally
depressurizing the reactor and connected systems. In modern high-pressure process
technologies, quality of control of conditions is such that only one or two of such events
is probable in a five-year period.
3.2.2.4 Current Operation
When low-pressure solution, slurry, and gas-phase process technologies capable of producing LLDPE were introduced in the 1960s through 1980s, it was anticipated that new
capital spending in LDPE plants would rapidly decline because of its higher capital
investment and energy costs and some product deficiencies compared to LLDPE in the
large polymer film market segment [5, 9, 15, 22]. However, the unique polymer structures possible with high-pressure reactors continue to provide advantages in melt processing equipment for film products and in high value products such as high-voltage
cable insulation. Also, the capability to incorporate large percentages of polar compounds such as vinyl acetate into the polymer structure enables production of very high
value polymers not possible with low-pressure processes. The lowering of capital and
operating costs from technology advances and avoidance of expensive coordination or
single-site catalysts, coupled with the product capabilities, has resulted in continued
plant expansions and licensing of technology for construction of new tubular reactors.
Additionally, to meet continuing market demand, old, smaller, and fully amortized
LDPE autoclave and tubular systems continue to operate. For example, in Europe most
autoclave systems are over thirty-five years old and less than 75,000 tonnes a year in
single-train capacity, whereas leading LDPE tubular technology licensors offer singletrain plants up to 450,000 tonnes per year. Additionally, a significant portion of melt
processing equipment is aged and under-powered to process LLDPE resins made by
the low-pressure processes. Thus, demand is retained for high-pressure resins for film
products.
3.2.3 Gas-Phase Fluidized Bed Reactors
Gas-phase reactors in vertical or horizontal configurations with mechanical stirring
have been used extensively for many years for production of a variety of polymers other
than polyolefins. It was not until the 1960s that such reactor systems were successfully developed for polypropylene (PP) by companies like BASF, Amoco, and Chisso.
However, such systems were unable to be used commercially for PE. The exothermic
heat of polymerization of ethylene is substantially higher than that for propylene. Thus,
the small difference between optimal reactor operating temperature and polymer sintering temperature in these reactors would make reactor control difficult. However, in
a fluidized bed reactor with its high level of mixing it is possible to remove the heat of
the highly exothermic reactions and keep temperature variations throughout the bed
below 1 °C.
In 1968, Union Carbide Corporation (UCC) was the first to commercialize a gasphase fluidized bed for HDPE using a chromium-based catalyst like that used in the
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Handbook of Industrial Polyethylene and Technology
slurry reaction systems [23, 24]. A short time later, Naphtachimie, a joint venture of
British Petroleum (BP) and Elf Atochem, commercialized a similar process. By 1977,
UCC had extended the technology to the production of LLDPE at polymer densities as
low as 0.918 g/cm3 using Ziegler-Natta titanium-based coordination catalysts. By the
mid-1980s, fluidized-bed reactor processes had become the leading licensing technology for new plants and were capable of producing both HDPE and LLDPE in the same
reaction system. In 1997, the licensing business for the UCC technology was transferred to Univation Technologies, a joint venture between UCC and ExxonMobil. By
2001, when The Dow Chemical Company acquired UCC and replaced it in the joint
venture, there were 200 fluidized-bed reactors in PE service worldwide, of which 100
were using UCC technology. Univation Technologies became a wholly owned subsidiary of The Dow Chemical Company in 2016.
The Naphtachimie technology was later assigned to BP, who made their own catalyst
and process advances to the technology. BP’s olefins and derivatives businesses were
divested in 2005 to a company called INEOS. The divestiture included the licensing
business of PE process technology for both slurry and gas-phase fluidized-bed processes under the name of Innovene. In the mid-1980s, a number of companies developed technologies with fluidized bed reactors operated in series to produce high-value
products with bimodal molecular weight distributions. These were not commercially
successful and recent developments have focused on hybrid processes utilizing both
a liquid-phase reactor and one or two gas-phase fluidized bed reactors in series (discussed later), or by using multi-active-site catalysts in a single fluidized bed such as
Univation’s Prodigy catalyst system.
3.2.3.1 Process Description
The basic configuration of the fluidized bed reactor is shown in Figure 3.5, and it is
similar among users of the technology [23–28]. Catalyst chemical formulations usually impregnated into inert support particles are fed into the lower section of a bed
of fluidized polymer particles just above a gas distribution plate. In some technologies a liquid-phase pre-polymerizer has been used to create the initial catalyst support.
Reaction starts within the support particle at active sites and the volume of polymer
made causes the support particle to fragment as more polymer is made. It normally,
however, remains as one particle as nascent molten polymer encapsulates the inert support particle fragments. A gas stream containing monomer is recycled through this bed
of polymer, providing a medium for fluidization, a source of reactant, and mechanism
for removal of the heat of reaction. Catalyst reaction activators such as aluminum alkyls,
a chain transfer agent (normally hydrogen), and comonomers are fed to the bed or into
the cycle gas stream. Above the reaction section is a larger-diameter expanded section
which causes most particles to disentrain because of the lower gas velocity. After exiting
the reactor, the cycle gas including fine entrained particles passes through a compressor
and one or more water-cooled heat exchangers to remove the heat of polymerization
before being returned to the bottom of the reactor below the distribution plate. In some
technologies, a cyclone has been used at the reactor outlet to return the fine entrained
particles to the fluidized bed. Reaction temperatures for PE production range from
70 to 120 °C and are dictated by the temperature for optimum catalyst kinetics while
Ethylene Polymerization Processes and Manufacture of Polyethylene
Cycle gas
compressor
Ethylene
Comonomer
Hydrogen
HX
Vent
recovery
Fluidizedbed reactor
Catalyst
CW
Cocatalyst
Purge
bin
Purge
gas
HX
Catalyst
feeder
75
Product
chamber
Additives
Extruder/pelletizer
Product
Product
blow tank
Figure 3.5 Diagram of a gas-phase fluidized bed reaction system licensed by Univation Technologies
[27, 29, 30].
remaining below the polymer sintering temperature. These reactors normally operate
at 20 to 30 bar. Polymer is discharged intermittently through a discharge valve into a
one- or two-stage tank system to reduce pressure and separate gas from solid. The solid
is then conveyed to a purge column to remove absorbed volatiles. Then, additives are
added and the polymer is melt-extruded and pelletized.
Production capacity is a function of the rate at which heat can be removed. The gas
velocity through the bed depends on the particle size distribution in the bed. The velocity is set to maintain turbulent fluidization conditions while not blowing particles out
of the bed [31]. Gas flow rate is therefore not defined by heat removal needs. However,
all of the heat generated in the bed must be removed by the cycle gas. In most gas-phase
systems, cycle gas must be prevented from condensing in the recycle system. Hence,
the reactor capacity is limited to the amount of sensible heat which can be removed
by the gas. This amount can be increased by addition of inert hydrocarbons of higher
heat capacity, such as propane, to the cycle gas. A major breakthrough developed by
Union Carbide allows cycle gas to be partially condensed in the recycle system. It is
achieved by use of condensable comonomers such as 1-hexene and/or by the addition
of inert hydrocarbons such as isopentane and hexane. The condensed droplets exiting
the cooler in the recycle system are carried back into the polymerizing bed by the cycle
gas flow. There, they evaporate rapidly just above the distribution plate. Removal of the
non-condensing constraint significantly increases heat removal capability because of
use of latent cooling on the condensed portion of the recycle stream. This process capability of running with “condensed liquid” in the recycle system can increase capacity
two- or three-fold for a specific size reactor. This technology was first commercialized
by UCC in 1984 [29]. Other companies such as BP developed their own variations [28].
76
Handbook of Industrial Polyethylene and Technology
In 1994 Exxon extended this technology to higher levels, condensing more than 20 wt%
of the recycle stream [32].
3.2.3.2 Product Capabilities
The fluidized-bed reactor process is able to use a large number of different catalyst families with widely different kinetics, including metallocene and other advanced singlesite catalysts [30, 33, 34]. It can thus make the full range of LLDPE and HDPE products.
The process system can operate over a broad range of selected operating temperatures
and gas compositions. Chain-transfer agents, such as hydrogen, and comonomers, such
as 1-butene and 1-hexene, can be used with different catalysts to produce essentially all
PE grades. The conversion per pass of cycle gas through the bed is only about 2%, which
allows for stable control of the required gas compositions for specific product grades.
Metallocene catalysts in particular have specific product and process advantages when
used in fluidized-bed reaction systems [30, 34]. The catalysts can make narrower MWD
polymer products and have a more uniform distribution of comonomer across the
molecular weight range. They also require less hydrogen and comonomer in the gas
phase than Ziegler-Natta catalysts. The more uniform comonomer distribution results
in a narrower melting point range, which increases the polymer sintering temperature,
allowing higher reactor temperatures. The lower comonomer requirement permits the
use of higher levels of inert condensing agents such as isopentane. Thus, reactor production capacity increases and investment and operating cost economics improve.
3.2.3.3 Process Challenges
Although fluidized-bed reactors are an established commercial success, they require
solving a wide range of technical challenges. The following discussion highlights some
of the critical technical challenges.
Catalyst design has been critical to the development of gas-phase technology. The
catalyst has to meet a wide range of criteria: 1) produce the required polymer, 2) be
responsive to chain transfer agents and comonomers for product range capability,
3) possess the desired kinetic behavior in the operating pressure and temperature range,
and 4) have an acceptable cost per weight of polymer made [8, 24, 31]. In addition, for
successful fluidized-bed operation, the catalyst must produce polymer particles that
meet the following fluidization criteria: 1) acceptable average particle size, and 2) have
an acceptable particle size distribution, and 3) acceptable shape. Together, these criteria
enable acceptably high fluidized and settled bulk densities. Achievement of these criteria requires a chemical process to create and load a catalyst species on a support particle
such as silica or alumina. The catalyst may already be active on contact with ethylene
or may require an activation step, usually within the fluidized reactor, to become a fully
active polymerization catalyst. Some technologies have avoided the need for a support
particle by using a pre-polymerization step in which the catalyst is activated and then
reacted with a small amount of ethylene to produce a particle which can then be fed to
the gas-phase fluidized bed. This pre-polymerization is usually performed in the liquid
phase and online such that the exiting catalyst-polymer particle is immediately fed into
a gas-phase fluidized bed. Achievement of an acceptable particle is also dependent on
Ethylene Polymerization Processes and Manufacture of Polyethylene
77
kinetic behavior during the first stages of reaction as it impacts particle growth rate,
support fragmentation, and heat liberation to the gas.
The primary operational concerns in fluidized-bed reaction systems are the formation of polymer agglomerates within the fluidized bed or recycle system, and the formation of thick polymer coatings on reactor walls [30, 31, 33]. This fouling occurs when
attractive forces hold polymer particles together such that they overheat and melt. As
an indication of the challenge when making LLDPE grades, the difference between
reactor operating temperature and the temperature at which sintering of the polymer
begins may only be about 15 °C. Catalyst families with positive kinetic activation energies (reaction rate increases with temperature), make effective mixing even more critical. However, catalyst kinetics has not been a barrier to technical success.
There is a peculiar phenomenon which aggravates fouling issues, namely electrostatics. PE is a dielectric material; accordingly, polymer particles can acquire and maintain an electrical charge upon contact with metal walls and internals. Researchers have
shown that electrostatic attraction can impact the hydrodynamics of fluidization and
adherence of particles to the wall [35, 36]. Companies have been able to address these
challenges using reactor and recycle system design criteria which ensure that the bed is
well mixed and that heat is transferred effectively to the recycle gas. As discussed, catalyst and support design are also critical through their formation of acceptable polymer
particles. In addition, companies have developed techniques to measure electrostatic
effects and to mitigate these surface chemistry phenomena through catalyst chemistry,
additives, and modification of metal surfaces inside the reactor.
Commercial fluidized-bed reactor technologies must be able to manufacture a full
range of PE products. The technology must be capable of operating with a wide range of
catalyst families and over a broad range of reactor gas compositions and temperatures.
Production of individual product grades requires transitions from one grade to another
[33, 37]. Within catalyst families or between chemically compatible catalyst families,
transitions are accomplished by adjustment of reactor conditions to those needed for
the next grade. These changes are called running transitions and are made with little
reduction in production rate or loss of prime polymer material, through a planned
product wheel cycle and computer-driven process control. In the case of incompatible
catalysts, it is necessary to discharge the reactor and then refill it with a compatible, dry
seed bed, with resultant loss of production for a certain period of time. Certain technologies claim the use of a pre-polymerizer or of a liquid-phase reactor in front of the
gas-phase reactor. In the latter case, Basell’s Spherilene technology grows a new polymer bed in an empty fluidized-bed reactor [37]. Even in this case, production cannot
reach design rates until the reactor is filled to the necessary operating height.
3.2.4 Slurry Reactors
3.2.4.1 Slurry Loop Reactors
The slurry loop process, which is primarily used to make HDPE, was the second major
process breakthrough, following the advent of the high-pressure processes by about
16 years. Processes [38] to make HDPE were initially developed by Phillips Petroleum
in the mid-1950s as a means of exploiting the discovery of chromium oxide (chrome)
78
Handbook of Industrial Polyethylene and Technology
catalysts in 1951. The Phillips slurry loop process debuted in 1961 [39]. This process,
including numerous improvements, is the dominant slurry loop process. At least 88
such reactors are operated or licensed worldwide by Chevron Phillips Chemical, the
successor to Phillips Petroleum, far more than any other slurry process [40]. At least
80 of these reactors are licensed [41]. In 2014, the name of the process was changed
to MarTECHTM.1 Specifically, the single-loop and dual-loop processes are now called
MarTECHTM Single Loop Slurry and MarTECHTM Advanced Dual Loop Slurry, respectively. The advanced dual-loop process was developed in collaboration with Total
Petrochemicals.
Numerous slurry loop processes for making HDPE have operated over the past few
decades. Nowadays, such slurry processes for producing HDPE employ high-activity
catalysts that allow catalyst residue to remain in the product. Several of them are similar to earlier designs of the Phillips process. A selected history of licensees that have
used the slurry loop process suggests how the process proliferated [42]. As early as
1959, National Distillers Co., Celanese, and Solvay became licensees of the Phillips
process. Years later, the manufacturing plants of these three licensees became part of
a joint venture between Solvay and BP International. In 2004, BP acquired Solvay’s
share of the joint venture and formed a company, Innovene, to hold its assets. A year
later, INEOS purchased Innovene, which included the interests of the original licensees.
Shortly thereafter, INEOS began licensing the Innovene loop slurry technology to third
parties. Chevron Phillips charged that INEOS was improperly disclosing their technology to third parties and requested that it be stopped. A court in Houston refused
INEOS’ objection to the temporary injunction. Rather than denying the disclosures,
INEOS argued that the Phillips technology was no longer proprietary. INEOS states in
their technology brochure that “INEOS and predecessor companies are the inventor of
HDPE P100 pipe grade using the Innovene S bimodal technology [43].” As of 2014, no
further legal action could be found. INEOS continued its licensing effort and has now
sold 14 licenses for the Innovene S slurry loop process to make HDPE [44], several of
them after the court decision.
Years before Solvay had entered into a joint venture with BP International, it had
developed its own HDPE slurry loop process. That process is now inactive. Other companies who have purchased licenses for the Phillips process and who had previously
licensed their own processes include National Distillers/USI and Showa Denko. The
upshot of this history discussion is that the basic principles of most slurry loop processes are similar. Therefore, description of the slurry loop process will, to avoid repetition, use the Phillips process as the primary example.
A schematic of the Phillips slurry loop process is shown in Figure 3.6. The schematic
shows a single loop, but also incorporates recent advances. Loop reactors behave as
CSTRs. In slurry loop processes, the slurry consists of undissolved polymer in a liquid
mixture of diluent, ethylene, comonomer, and catalyst, and optionally cocatalyst and
hydrogen. The components of the slurry, other than polymer, are continuously fed into
the loop. The slurry is circulated through the loop using axial flow pumps [45]. Cooling
jackets around the vertical sections of the loop remove the heat of polymerization.
Slurry, somewhat concentrated in solids, is discharged from settling legs periodically
1
MarTECH is a trademark of the Chevron Phillips Chemical Company.
Ethylene Polymerization Processes and Manufacture of Polyethylene
Loop
reactor
Deoxo
bed
Line heaters
Coolant
Deactivator
Flash
vessel
Catalyst
mix tank
Axial
pump
Cocatalyst
Purification
Ethylene
1-Hexene
H2
Degassing
Recycle
isobutane and
1-hexene
purification
79
HX
HX
CW
1-Hexene
condensate
Isobutane
condensate
Steam
CW
Purge
column
Heavies
removal
column
N2
Dried
fluff
Lights
removal
column
Isobutane
purification
Isobutane
Figure 3.6 The Phillips (MarTECH) single loop slurry process for production of HDPE. Dual-loop
operation is also employed [40, 46, 47].
or continuously. Slurry exiting the reactor is injected with a catalyst deactivator. Light
components are removed using a flash tank. The resulting polymer fluff is degassed
further in a purge column. The light components are separated in at least two distillation columns to remove the heaviest and the lightest components. The remaining fluid,
mostly diluent and comonomer, is recycled to the reactor.
Diluents used in commercial processes include isobutane, hexane, cyclohexane, and
propane. The diluent must not dissolve the polymer or fouling may occur. Branched
alkanes are preferred because they are poorer solvents [8].
In the Phillips slurry loop process, the diluent is isobutane. Comonomer is usually
1-hexene, but could be 1-butene [45]. Both chrome and Z N catalysts are used in the
process. Metallocene catalyst is used in the dual-loop slurry reactor [48]. Reactor residence time is 30 to 60 minutes [46], but may be as long as 3 hours in some processes [12].
Process conditions are given in Table 3.3 in the rows showing the MarTECH process.
Downstream separation steps in slurry processes for HDPE require considerable
equipment. In all processes using a diluent, the diluent as well as comonomer must
be separated from the slurry exiting the reactor and then recycled. Otherwise, the cost
of use and disposal of these components would be prohibitive. Typically, one or more
flash steps are followed by condensation of the diluent and comonomer. Purification
of these liquids is required before they are recycled, particularly if catalyst in the slurry
exiting the reactor has been deactivated with a poison. The liquid-free polymer generally requires devolatilization to remove dissolved monomers and diluent. In the Phillips
process, this last step is carried out by feeding nitrogen gas up through a countercurrent purge tower in which polymer fluff moves downward. In other slurry processes, a
drying step of an unspecified type is employed [23]. Hottovy and Kreischer of Phillips
Petroleum describe an improved downstream separation system that includes four
flash tanks, two gas compressors, and three heat exchangers [49]. A diagram of their
scheme is shown in Figure 3.7. However, subsequent information indicates that just a
Innovene S
INEOS
Total
Petrochemicals
MarTECHTM
Advanced
Dual Loop
Slurry
Z-N
MarTECHTM
Advanced
Dual Loop
Slurry
Chevron-Phillips
Chemical
and Total
Petrochemicals
Solvay
Cr;
Z-N
MarTECHTM
Single Loop
Slurry
Chevron-Phillips
Chemical
0.948–0.961
0.934–0.955
Z-N; silicasupported
metallocene
0.945–0.962;
0.937;
0.927–
0.935
Z-N Supported HDPE
Z-N
Cr
Catalysts
Process name
Company
Resin density
g/cm3
Co-monomer
1-hexene
uni-modal 1-butene
bimodal
1-hexene
bimodal
1-hexene
uni-modal 1-butene
uni-modal 1-hexene;
bimodal
1-butene
isobutane
n-hexane
isobutane
isobutane
isobutane
isobutane
Diluent
Reactor characteristics
uni-modal 1-hexene
MWD
7–55 (21.6 kg); bimodal
0.23–0.25;
0.9
0.025–0.3
0.1–5;
1–100
Melt Index
g/10 min
2.16 kg
Product properties
Table 3.3 Process conditions and product range for slurry loop processes.
80–110
75–85
80–110
85–105
Temp.
°C
37–45
30
10;
37–45
40
Pressure
bar
150
30–60
Resid.
time
min
80
Handbook of Industrial Polyethylene and Technology
Ethylene Polymerization Processes and Manufacture of Polyethylene
Vapor
Cooler
Vapor
From
reactor
Flash
tank
Flash
tank
C4– and
C2H4
liquid
Diluent
liquid
Cooler
Flash
tank
Flash
tank
Filter
Solids
and
liquid
C4– and
C2H4
vapor
Cooler
O2 removal
(optional)
Filter
81
C5+ and
1-hexene
Vapor
Liquid
Flash
tank
Polymer
product
Figure 3.7 Downstream separation system for the Phillips Process, published in the year 2000 but
subsequently simplified [49].
single-stage flash tank is used [40]. Carrying out the flash at high pressure reduces the
energy needed to recompress the isobutane. And, purification of the recycled isobutane
is no longer necessary.
Design and operation of a slurry loop reactor have evolved considerably. The loop
is basically a pipe. Its diameter represents a tradeoff between reducing pumping power
consumption with larger diameters and increasing cooling surface area by using smaller
diameters. Pipe diameter in commercial reactors ranges from 0.4 to 1 m. Production
rate has increased dramatically by increasing reactor volume by more than two orders
of magnitude over the years [47], reaching 265 m3 [25]. Higher-activity catalysts have
enabled shortening of residence time. Correspondingly, total pipe length has reached
300 m [50] in order to meet cooling needs. For long pipe lengths, multiple pumps ought
to be needed along the loops in order to maintain a reasonably low pressure drop across
each pump. However, Nowlin states that Chevron Phillips has been able to maintain a
single-pump design while maintaining the pipe diameter at 0.6 m [25]. Velocity of the
slurry is also dictated by opposing factors. Higher velocity is required for higher rates
of heat transfer, turbulent mixing, and reduced fouling. However, pressure drop and
pumping cost increase at high velocity. The minimum velocity is 5 m/s. The maximum
velocity is variously reported as 10 [50], 12 [49] or 15 [48] m/s.
Economic considerations in the Phillips process favor operation at a high polymer
solids concentration in the slurry. High solids reduce the cost of downstream separations. Furthermore, it becomes viable to replace settling legs at product discharge
from the loop with continuous discharge. A larger diameter axial pump, inserted in
an enlarged diameter pumping section, provides a higher slurry velocity [51, 52]. This
development contributed to an increase in solids concentration from 25 [1] to 60 wt%
[50, 52]. Higher bulk density of the polymer fluff also contributes to higher weight percent solids. The bulk density depends on the nature of the catalyst, and for supported
catalyst its morphology.
The most recent innovations in slurry reactors are the use of two or more reactors,
usually in series, to make products with bimodal MWD. A patent by Marechal assigned
82
Handbook of Industrial Polyethylene and Technology
to Total Petrochemical [52] discloses that catalyst is fed to the first loop reactor only.
Polymer having high MW is produced in one reactor and low MW polymer is produced
in the other reactor. For optimum product, high MW polymer is made at low polymer
density and low MW polymer is made at high density. The second reactor additionally
serves as a polymer blender. When Z-N catalyst is used, high MW, low density polymer is preferably made in the first reactor because little or no hydrogen is needed. This
sequence avoids the need to remove the hydrogen before the polymer enters the second
reactor. If metallocene catalyst were used, low-MW polymer could be made in the first
reactor because less hydrogen is needed to make it.
Major benefits of slurry loop reactors include: uniform temperature, scale-up to
large capacity, use of most catalyst types, and the ability to make a broad range of
HDPE products. Because the loop is operated at relatively high velocity, fouling can be
virtually eliminated, allowing on-stream time of Phillips process licensees to reach 97
to 98% [40]. The Phillips process produces HDPE products commanding the highest
profit margins. The 30 to 60 minute residence time of the Phillips slurry loop process is
average among PE processes. However, the longer residence time of some other slurry
loop processes, such as the old Solvay process [12], increases off-grade during product
transitions.
A major limitation of the slurry loop process is the inability to make most LLDPE
grades, particularly because of fouling issues. Another major issue, also associated with
other liquid-phase reactors, is that extensive post-reaction separations are required.
Other shortcomings of slurry loop reactors include concentration gradients, unstable
startup, fines, wax accumulation, and fouling [50]. Concentration gradients, which
are most likely to occur at low slurry velocities, can cause broadening of the MWD.
Instability at start-up is characterized or caused by oscillating polymerization rates.
Production of fines can be reduced by pre-polymerization of the supported catalyst
until particle size increases about tenfold. INEOS claims elimination of fouling in their
Innovene S slurry loop reactors [44]. Elimination of wax buildup is claimed in both the
Phillips and Innovene S processes [40, 53].
Companies other than Chevron Phillips that operate or had operated slurry loop
reactors to make HDPE include INEOS, Total Petrochemicals, Solvay, Showa Denko
and USI/National Distillers.
INEOS has sold 14 Innovene S licenses since it began licensing in 2005, most recently
at a proposed cracker site in West Virginia [44]. Its slurry loop reactor and product capabilities appear similar to that of the Phillips process, as shown in Table 3.3. However,
INEOS states that there are no distillation columns. Also, unlike the old Solvay process,
there are no centrifuges [44, 53]. In addition to single-loop technology, dual-loop technology is employed to make bimodal products [43]. Among their products is HDPE for
PE100 pipe grade, which has rigorous specifications. INEOS has worked on the means
to make LLDPE in the Innovene S process. Both Z-N and chrome catalysts are used.
The largest single-train plant built using Innovene S process is 345,000 tonnes/y [54],
and technology is claimed for building a 600,000 tonnes/y plant. An economic analysis
suggests that capital cost for a 400,000 tonnes/y plant is less than that of other bimodal
processes [55].
Solvay had been a licensor of its slurry loop process to make HDPE. The process uses
supported Z-N catalyst. Diluent is variously reported as hexane [12, 56] or isobutane
Ethylene Polymerization Processes and Manufacture of Polyethylene
83
[23]. A 2001 patent [57] mentions use of isobutane as the diluent and 1-hexene as the
comonomer. It also describes two loop reactors in series to make bimodal products.
The patent is now owned by INEOS. A schematic of the process is shown in Figure 3.8.
The process particularly deviates from the Phillips process in that it employs a long residence time of 150 minutes [12]. Also, it uses a centrifuge as well as a stripping device
and dryer to devolatilize the fluffy product.
Showa Denko, an early licensee of the Phillips process [42], perfected its own slurry
loop technology and had licensed it at least as recently as 2001 to BP (later canceled).
It offered Z-N and chrome catalysts [58]. Another early licensor of slurry loop technology was National Distillers/USI. It employed Z-N catalyst and a heavy diluent [12].
National Distillers had also been a prior licensee of the Phillips process.
3.2.4.2 Slurry CSTR Reactors
Production of HDPE in a slurry continuous stirred tank reactor (CSTR) rather than
a loop reactor was another early means of exploiting catalytic polymerization. In fact,
several companies, including Hoechst, Phillips Petroleum, and Mitsui Chemicals, built
CSTRs to make HDPE in the 1950s before the loop process was introduced. During
the past 50 years, many companies built slurry CSTR processes to make HDPE. Today,
it continues to be widely used, though not as much as the loop process. An important
legacy of the process was that it became a springboard to the solution processes for
making LLDPE. Major licensed CSTR processes include the BASF Hostalen Advanced
Cascade Process licensed by LyondellBasell and Mitsui’s CX process.
The main differences between the slurry CSTR and slurry loop processes are in the
means of agitation and cooling of the reactor(s). Slurry CSTR processes use one, two, or
Isobutane
Purification
Diluent
Recovery
Ethylene
Lights
condenser
Recycle
column
Flash
tank
Loop
reactor
Comonomer
Stripper
Catalyst
H2
CW
Axial
pump
Centrifuge
Dryer
to
Pelletizer
Figure 3.8 Solvay slurry loop process with a single loop. Dual-loop technology was added later [23].
84
Handbook of Industrial Polyethylene and Technology
three reactors, usually in series, to produce HDPE. The use of multiple reactors, as with
loop reactors, enables production of polymers having tailored, broad MWD. Diluent,
catalyst, and optionally cocatalyst, are fed into the first reactor. Ethylene, comonomer and hydrogen may be fed into any or all reactors. Agitation of the CSTR typically
employs a rotating shaft with blades and optionally baffles. The agitator design is critical: early designs doomed processes developed by Phillips Petroleum [25] and The Dow
Chemical Company [59] because of fouling. The downstream challenges of separating
diluent and comonomer from the product and then recycling these constituents to the
reaction system have similarities to those for slurry loop reactors. Process conditions,
product range, and capacity are listed for slurry CSTR producers in Table 3.4.
The Hostalen bimodal process for producing HDPE in two slurry CSTRs was commercialized by BASF in 1975. The process uses Z-N catalysts. The process became part
of Basell Polyolefins, which purchased Lyondell in 2007 and became LyondellBasell.
LyondellBasell and its joint ventures operate four [60] or five [61] Hostalen HDPE
plants. In 2009, the process was enhanced to use three CSTRs and was offered for
license as the Hostalen Advanced Cascade Process (ACP) [62]. LyondellBasell has converted their older plants using the bimodal process to ACP. At least one licensee has
been signed up, Shaanxi Yanchang Petroleum Yanan Energy and Chemical Co. [63].
A schematic of the process using three reactors is shown in Figure 3.9. The ACP process
employs cooling jackets around each reactor as well as interstage cooling [60]. It also
uses a centrifuge for devolatilization of the product. The Hostalen bimodal process was
used in 23 operating lines as of 2005, producing 2,900,000 tonnes/year [5]. The newer
ACP process currently produces 1,290,000 tonnes annually [60].
LyondellBasell also licenses the Equistar slurry CSTR process. This process was
developed by the Nissan-Maruzen Polyethylene Company in 1981. It employs two
CSTRs either in series or parallel [65]. Bimodal MWD products can be made in the
series configuration. The process uses 1-butene as the comonomer and hexane as the
diluent [23].
Another slurry CSTR process is the CX process licensed by Mitsui Chemicals.
Worldwide capacity using this process was 3,500,000 tonnes/year in 2005 [5]. It employs
two CSTRs in series to make HDPE and medium density PE (MDPE), and makes both
unimodal and bimodal products [66]. A comparison of the CX process operating conditions in Table 3.4 with that of the Hostalen ACP process indicates many similarities. Both use Z-N catalyst. Operating temperatures and pressures are similar. Both use
1-butene comonomer and hexane as the diluent.
A schematic of the two-reactor CX process that is apparently used by the licensee
TVK in Slovakia, is shown in Figure 3.10 [45]. A more detailed schematic in the reference highlights the multifaceted cooling system needed for slurry CSTRs. Three heat
exchangers employed around each reactor remove the heat of reaction: a reactor jacket,
a slurry cooler on a recirculation stream, and an overhead condenser. This process illustrates the huge role of cooling in reactor design. A super-high-activity Z-N catalyst is
used [66]. The process uses a centrifuge to separate diluent from the polymer. Low
MW polymer is removed from the diluent before it and condensed hexane vapor are
recycled to the first reactor. The polymer proceeds to a final devolatilization step.
Asahi Kasei and LG Chem have also developed slurry CSTR processes for the production of HDPE. They are not currently active licensors. In each case, these companies
CX
Mitsui
Z-N
multiple
Metallocene
Nissan/Maruzen/
Equistar
LyondellBasell
(Equistar)
0.937–
0.964
HDPE
HDPE
Z-N
LG
Hostalen Advanced
Cascade Process
(ACP)
LyondellBasell
HDPE
Z-N
Z-N
Metallocene
Hostalen Bimodal
LyondellBasell
Catalysts
Asahi Kasei
Process name
Company
Resin
density
g/cm3
0.001–20
Melt index
g/10
min
2.16 kg
Product properties
Table 3.4 Process conditions and product range for slurry CSTR processes.
bimodal
bimodal
uni-modal
bimodal
uni-modal
bimodal
uni-modal
bimodal
multi-modal
uni-modal
bimodal
MWD
1-butene
hexane
hexane
hexane
hexane
1-butene
1-butene
hexane
Diluent
1-butene
Co-monomer
Reactor characteristics
70–90
65–85
75–90
Temp.
°C
6–10
10–14
5–10
Pressure
bar
45 per reactor
42–150
Resid. time
min
Ethylene Polymerization Processes and Manufacture of Polyethylene
85
86
Handbook of Industrial Polyethylene and Technology
1-Butene
1-butene and
hexane recovery
Wax
Hexane
Reactor 1
Catalyst
CW
Reactor 2
Receiver
Post
reactor
CW
CW
CW
CW
CW
Centrifuge
Dryer
Hexane
To extruder and
pelletization
Ethylene
1-Butene
H2
Figure 3.9 Hostalen Advanced Cascade Process with three slurry CSTRs [23, 60, 64].
Catalyst
tank
Lights
Hexane
Ethylene
Diluent
recovery
column
Steam
Recycle diluent
1-Butene
H2
Flash
tank
CSTR 2
CSTR 1
CW
Centrifuge
CW
Dryer
CW
CW
Powder to
extruder
Figure 3.10 Mitsui CX Process, a slurry CSTR process for production of HDPE [23, 45, 66].
have adapted their two-CSTR process to use metallocene catalyst [67]. Asahi Kasei had
licensed INSITE constrained geometry catalyst technology around 1997 from The
Dow Chemical Company.2 Dow Chemical itself had developed low pressure HDPE
technology early on. Fouling issues and operating costs were alleviated by development
of high-efficiency Z-N catalysts and by licensing “anhydrous slurry” process technology from Asahi Kasei approximately in the late 1970s [59]. Dow continued to operate a
slurry CSTR plant in Thailand for many years.
The main advantage of slurry multiple-CSTR reactors is the ability to produce HDPE
products with a tailored, broad MWD. Investment cost is probably the highest among
low-pressure processes. Reasons include the use of multiple reactors, and as with slurry
2
INSITE is a trademark of The Dow Chemical Company.
Ethylene Polymerization Processes and Manufacture of Polyethylene
87
loop processes, the use of extensive equipment for downstream separation and recycle
of diluent and comonomer. Because the diluent hexane has a higher boiling point than
isobutane, its separation from the polymer is expected to be more difficult and expensive than most slurry loop processes. The ability to scale up is good. Proven singletrain scale-up to 450,000 tonnes/year for the Hostalen ACP process is almost as high as
the maximum capacity of the Phillips slurry loop process. Many different slurry CSTR
processes operate today. However, most are old and few new licenses have been added
recently.
3.2.5 Solution Reactors
3.2.5.1 The Dow Chemical Company’s DOWLEX Process
The world’s most prolific producer of solution LLDPE is The Dow Chemical Company.
Dow employs a low-pressure, catalytic solution process known as the DOWLEXTM process.3 As early as 1955, Dow built a tubular reactor operating at about 1400 bar and a
temperature of 200 to 250 °C [25]. It used Z-N catalyst to make polymer in solution.
Residence time was a relatively short at 8 minutes. The product had a high molecular
weight and a medium density. Dow’s efforts in the 1960s to produce low-pressure, linear PE included a solution process to make HDPE. The creation of high-activity Z-N
catalysts in the 1970s enabled Dow to adapt that HDPE process to produce LLDPE
in solution using higher alpha-olefin comonomers. Comonomers included 1-butene,
1-hexene, and 1-octene. Today, 1-octene is the most widely used comonomer in the
DOWLEX process.
The DOWLEX process is not licensed. It is only used by The Dow Chemical Company
and its joint ventures. While Dow has patented its PE technology heavily, nearly all of
the patents pertain to catalyst or product polymer. Publicly available process details are
sparse. An operating line typically contains two CSTRs in series. Reactor temperature
is well above the melting point of the polymer. Temperature and pressure are chosen
such that all of the polymer dissolves in the solution. When using Z-N catalyst, the
diluent is a solvent, particularly ISOPARTM E [23, 68], a mixture of C8 and C9 isoparaffins.4 Yet, Soares and McKenna [8] state that normal (unbranched) paraffins should be
better solvents. The addition of lighter alkanes to the solvent enhances the amount of
heat of polymerization that can be removed [68]. Reactor temperature when using a
Z-N catalyst is about 160 to 220 °C [25, 68], but lower temperatures may be used with
metallocene catalysts. Pressure is about 25 bar [23, 25]. Residence time through both
reactors is about 30 minutes [25]. Process conditions and product range of the various
solution processes are summarized in Table 3.5.
Further details of the DOWLEX process are based on the Tau patent [68] filed in
1998 and other sources. Simple schematics for two configurations of the process are
shown in Figures 3.11 and 3.12. Figure 3.11 shows two adiabatic CSTRs in series. There
is no heat removal in the adiabatic reactors; the temperature reached depends on the
rate of polymerization. The weight percent polymer in the solvent increases from one
3
4
DOWLEX is a trademark of The Dow Chemical Company.
ISOPAR is a trademark of ExxonMobil Corporation.
Process
SCLAIRTECH
Advanced
SCLAIRTECH
Dowlex
Stamicarbon
Compact
DuPont Solution
Company
Nova
Chemical
Nova
Chemical
Dow
Chemical
SABIC (was
DSM)
Equistar
VOCl
Z-N
Z-N, constrained
geometry
Z-N
Single-site
VOCl3,
TiCl4,
DEAO
Catalysts
HDPE
0.900–0.967
0.8–100
0.5–25
0.904–0.964
tailored
narrow
mediumbroad;
narrowbroad
0.4–10
0.4–150
0.912–
0.967;
0.915–
0.967
MWD
mediumbroad
Melt Index
g/10
min
2.16 kg
0.2–100
0.920–0.963
Resin density
g/cm3
Product properties
1-octene
1-butene
1-octene
hexane
up to 285
130–200
160–220
ISOPARTM E
C8 and C9
iso-paraffins
1-octene
1-butene
< 200
lighter than
cyclo-hexane
1-octene
75–300
Temp. °C
cyclo-hexane
Solvent
1-butene,
1-octene
Co-monomer
Reactor characteristics
Table 3.5 Process conditions and product range of solution processes for production of LLDPE.
23–27
Supercritical
40–200; up
to 100
Pressure
bar
short
2–30
30
2–3
2–3
Resid.
time
min
88
Handbook of Industrial Polyethylene and Technology
Ethylene Polymerization Processes and Manufacture of Polyethylene
89
Recycled monomers and solvent
Purification
column
Vapor
HX
Pressure reduction
valves
10 wt%
polymer
CSTR 1
15 wt%
polymer
Spray
nozzle
Cooler
Flash
vessel
CSTR 2
30 wt%
polymer
Monomer and
solvent vapor
Devolatilization
unit
HX
Heater
86 wt%
polymer
Former location of heater
Fresh monomers,
catalyst and solvent
Figure 3.11 Adiabatic reactors in series in The Dow Chemical Company’s Dowlex solution process.
Product in solution is heated after being flashed to a lower pressure [68].
HX
CW
Recycle
vapor
Recycle
solvent
Refrigerant
Optional HX
Knockout
drum
CSTR 2
Make-up
comonomer
Polymer dissolved
in liquid
Make-up
ethylene
To downstream separation
Figure 3.12 The Dowlex boiling reactor process. Reactor vapor is cooled, partially condensed, and
recycled to the reactor [68].
reactor to the next. The solution exiting the second reactor in a former configuration was heated before it enters a devolatilization unit. In the improvement shown in
Figure 3.11, heating is delayed until after the solution is flashed to a lower pressure. This
step enables a greater concentration of polymer to be achieved in the second reactor,
providing higher production rates and reduced energy expenditure.
The boiling reactor of Figure 3.12 provides enhanced cooling. A vapor stream consisting of monomers and solvent is continuously withdrawn from the top of the reactor
and cooled, removing more than 15% of the heat of polymerization. The fluid then
enters a knockout pot. Vapor exiting the top of the pot is compressed; liquid from the
90
Handbook of Industrial Polyethylene and Technology
bottom of the pot is pumped, and both are returned to the reactor. Optionally, these two
streams pass through purification columns. Polymer-containing liquid exiting the reactor proceeds to the devolatilization unit. The cool vapor and condensate are returned
to the reactor. Not only is the reactor temperature controlled, but more importantly the
production rate can be much higher. The returning cool stream also provides agitation
to the reactor. Typically two such reactors are in a single train.
There are a number of advantages of the DOWLEX process. The LLDPE products and
ultra-low-density specialty products are highly regarded. The two-reactor system provides a degree of control of MWD. Very-high-activity catalysts can be exploited, allowing for short residence times as long as there is sufficient cooling capacity. Therefore,
transitions to a different product in the same reactor proceed quickly, minimizing offgrade. Operating temperatures can be quite high, typically limited by degradation of
the polymer at temperatures of 230 °C or higher.
Difficulty in cooling large reactors is probably the major limitation of the solution
CSTR process. Reactors are moderately small. An example in the Tau patent employing
a Z-N catalyst suggested a volume of 60 m3 per reactor at an assumed residence time
of 15 minutes. Production rates for a single train would approach 300,000 tonnes per
year. The first grassroots plant designed to use INSITE catalyst, started up in 1998, has
a production rate of 210,000 tonnes per year [69]. While these rates are substantial,
they do not match that of the largest UNIPOL or MarTECH reactors. The DOWLEX
process is not often used to make HDPE. A further limitation, which affects all solution
LLDPE processes, is MW. High MW polymer either does not dissolve in the solvent or
increases solution viscosity [25, 68].
3.2.5.2 The SCLAIRTECH Process
The origin of NOVA Chemical’s Advanced SCLAIRTECHTM process began in the mid1950s when DuPont developed a solution process to make HDPE [70].5 As in most
other processes of that era, Z-N catalysts were used. Efforts by a group of young chemists and chemical engineers at DuPont Canada to convert this process into one that
could also make LLDPE were rewarded when in 1960 the world’s first LLDPE plant
started up. Licensing of the SCLAIRTECH process began in 1980. By 1997, there were
11 lines in operation worldwide. In 1994, the SCLAIR business was sold to NOVACOR
Chemicals, now known as NOVA Chemicals. Advances in the Z-N catalyst and reactor
design in 1998 led to the introduction of the Advanced SCLAIRTECH process [71]. In
2011, NOVA Chemicals became a subsidiary of the International Petroleum Investment
Company (IPIC) of the Emirate of Abu Dhabi.
The two-reactor SCLAIRTECH process is reported to employ a tubular reactor and
a CSTR [71]. The reactors are quite small, having volumes in the range of 2.8 to 9 m3
[25]. When the process is combined with high-activity catalysts, residence time can be
reduced to 2 minutes [23]. Solvent for this solution process is cyclohexane. Catalysts used
are TiCl4 and VOCl3 with diethyl aluminum ethoxide cocatalyst. A vanadium-based
catalyst, VO[OR]2Cl where R is an alkyl group such as decyl, is said to make LLDPE
with a homogeneous (random) distribution of short-chain branches [25]. Comonomer
5
SCLAIRTECH is a trademark of NOVA Chemicals Company.
Ethylene Polymerization Processes and Manufacture of Polyethylene
Ethylene
1-Octene
Solvent
Ethylene
91
Purification
columns
HX
HX
Distillation
CW columns
To solvent recovery
HX
High-intensity
mixers
1-Octene
CW
HX
Deactivator
Stabilizer
HP steam
CW
HX
Solvent
H2 Catalyst 1
Phase
separation
Flash
vessel
Polymer
stripper
H2 Catalyst 2
Extruder and pelletizer
Polymer
product
Figure 3.13 Advanced SCLAIRTECH solution process for producing LLDPE [23, 45, 72, 73].
is predominantly 1-butene. Product capability ranges from LLDPE to HDPE. Products
have a wide range of melt index and medium-broad MWD, as shown in Table 3.5.
In the SCLAIRTECH process, catalyst and monomers are dissolved in solvent before
being introduced to the first reactor. Cooling is convective [46]; heat exchangers are
located before and after the first CSTR [45]. Reactor temperature may approach 300 °C.
A relatively high pressure of 100 to 200 bar is used to facilitate separation of the polymer from the fluid. The catalyst in the solution exiting the reactor is deactivated in situ,
and may be passed through a bed of alumina to remove vanadium residues. Separation
of the polymer from solvent and monomers employs conditions to cause phase separation. Antioxidant and stabilizer are melt-mixed into the polymer before the polymer is
extruded into pellets, and then stripped of residual solvent and monomer. Solvent from
both the phase separator and the stripper is cleaned up in a series of five distillation
columns before being returned to the first reactor. The SCLAIRTECH process has been
licensed worldwide.
A schematic of the Advanced SCLAIRTECH process is shown in Figure 3.13.
Multiple CSTRs are employed in configurations such that polymers having different
properties are made in each reactor. For example, products with bimodal MWD can be
made. Metallocene, single-site and Z-N catalysts are used. The comonomer is 1-octene.
A number of advantages have been incorporated into the advanced process. The solvent
is a lower boiling one than cyclohexane, reducing energy needs for downstream separation. Controlled high-intensity mixing is an important reactor improvement. Reactor
temperatures are lower, allowing the use of steam for heating. The product stripper
section can handle the increased comonomer load of lower density polymer. Removal
of catalyst residue is no longer needed. The process has been shown to operate using
ethylene containing up to 15% ethane, thereby avoiding the need to separate ethane
from ethylene at the cracker.
Because operating temperature and pressure brings the solution to near its critical point, modeling with AspenTech’s PolymersPLUS software is employed to take
advantage of phase separation for isolation of the product.6 However, the models
6
Aspen Plus is a registered trademark of Aspen Technologies, Inc.
92
Handbook of Industrial Polyethylene and Technology
within the software were inadequate to predict conditions of the solvent near the
critical point when the polymer is present [74]. Advanced process control software
is also used.
The Advanced SCLAIRTECH process possesses significant improvements over its
predecessor. A major advantage over other types of processes is the ability to use small
reactors while achieving high production rates up to 425,000 tonnes/year [72]. The
small size allows product transitions to be completed in 30 minutes. Both LLDPE and
HDPE are produced. A wide variety of products can be made using multiple reactors.
However, equipment cost is relatively high. This disadvantage derives from extensive
downstream separations that are endemic to liquid-phase processes, and also from the
high pressure required. Some relief in energy cost is derived from carrying out part of
the downstream separation in the supercritical region. High MW polymers cannot be
made in the solution process.
3.2.5.3 The Stamicarbon Compact Process
The Compact solution process was developed by Dutch State Mines (DSM). In the
1970s it was first used to make HDPE. In the 1980s, it was adapted for production of
LLDPE. The process had been licensed by the Stamicarbon subsidiary of DSM. In 2002,
components of DSM, including PE production, were purchased by SABIC. Currently,
there does not appear to be licensing activity for the Compact process.
The process uses a Z-N catalyst. Comonomer is preferably 1-octene and the solvent
is hexane. The reactor is an adiabatic, liquid-filled CSTR [23]. As in the SCLAIRTECH
process, the residence time is short at about 2 minutes. A schematic of the process
[23, 46] in Figure 3.14 illustrates the complexity of the downstream separations, as in
other solution processes. Polymer concentration in the solvent is relatively low, less
than 10%, inferring a higher cost of downstream separations. The reactor can make 1
tonne of HDPE/hour in 1 m3 of reactor volume. A broad range of LLDPE and HDPE
products are made [75].
1-Octene
1-Octene
recovery
Hexane
HX
CW
Flash 1
Mixer
Flash 2
Steam
H2
Catalyst
Volatiles
Additives
CSTR
Mixing
Ethylene
Hexane
purification
Catalyst
deactivator
HX
Steam
HX
Extruder and
pelletizer
Polymer
product
Figure 3.14 Stamicarbon Compact process of DSM for producing LLDPE in solution [23, 46].
Ethylene Polymerization Processes and Manufacture of Polyethylene
93
3.2.5.4 Equistar Solution Process
An older type of solution process developed by DuPont had been operating until
recently. Equistar operated the Victoria Texas solution plant for the rare purpose of
producing HDPE [76]. Operating parameters, shown in Table 3.5, appear similar to
those of the SCLAIRTECH process. Reactor temperature is high, residence time is
short, and a VOCl3 catalyst is used. The simple schematic shown in Figure 3.15 reveals
three reactors in series, two of which are CSTRs followed by a tubular reactor. Catalyst
residue removal is necessary [77]. The source of this technology is DuPont. DuPont
sold some of its PE business around 1988 to Cain Chemical. It was subsequently transferred to Occidental Petroleum [78]. In 1997, Equistar was created as a subsidiary of
Lyondell when Lyondell acquired part of Occidental. The HDPE products made with
this technology at Victoria continued to use the old DuPont name Alathon .7
3.2.6
Hybrid Processes
There are two primary hybrid technologies, Spherilene from Basell and Borstar from
Borealis. Neither technology has been licensed extensively by unaffiliated companies.
3.2.6.1 Spherilene Hybrid Process
Himont, which was later acquired by Basell, had a commercially successful slurry loop
reactor followed by a gas-phase reactor system called Spheripol for production of PP.
Himont extended the technology to HDPE and it was further developed by Basell
[23, 37, 79]. The system comprises an initial small slurry loop reactor which is used for
catalyst activation and pre-polymerization followed by one or two gas-phase fluidized
beds in series. The catalyst is called Avant Z catalyst by Basell and is an extension of
Himont’s reactor granule technology. The particle morphology and reaction kinetics
are tailored by catalyst and support technology and by the pre-polymerization step.
The formed polymer particle exiting this step maintains its integrity throughout the
following reaction steps and leaves the subsequent reaction steps not as a powder but
CSTR 1
CSTR 2
Tubular reactor
Ethylene
1-Octene
H2
Catalyst
removal
Catalyst
Solvent
Extruder and
pelletizer
Polymer
product
Figure 3.15 Equistar solution process for production of HDPE using former DuPont technology [76].
7
Alathon is now a registered trademark of LyondellBasell.
94
Handbook of Industrial Polyethylene and Technology
Vent
recovery
Vent
recovery
Gas-phase
fluidized-bed
reactor
Liquid-phase
prepolymerizer
reactor
HX
Gas-phase
fluidized-bed
reactor
HX
Product
discharge
tanks
PurgeAdditives
PurgeAdditives
column
Propane
Purge
column
Catalyst +
cocatalyst
Ethylene
Comonomer
Hydrogen
Reagents
Extruder /
Pelletizer
Product
Purification
Figure 3.16 Hybrid process for production of linear PE; Spherilene C configuration shown [23, 80].
as free-flowing beads of 0.5 to 3.0 mm diameter. All catalyst is fed to the initial slurry
reactor and it retains its polymerization activity through all reaction steps. For startups and transitions, there is no need for a “seed bed” in the gas-phase reactors as the
polymer particles made in the initial liquid phase reactor are used to fill those reactors.
The gas-phase reactors do not operate in condensing mode. They use an inert diluent,
propane, instead of nitrogen. The higher heat capacity of propane combined with the
controlled morphology and polymerization reaction kinetics of the Avant Z catalysts
results in a process which is not susceptible to polymer sintering or wall fouling. The
process can also use the Avant C chromium-based catalyst. When only one gas-phase
reactor is installed it is called Spherilene S. When a second gas-phase reactor is installed
in series it is called Spherilene C, as shown in Figure 3.16. This mode adds capability
for broad bimodal MWD products by independent control of gas compositions in each
gas-phase reactor. Very precise process control of density and molecular weight in each
gas-phase reactor is required. The split of production between each reactor produces
these high-value bimodal products.
3.2.6.2 Borstar Hybrid Process
The Borealis Borstar process is designed to produce PE products with a broad range
of properties not possible in single-reactor systems [81–83]. The primary feature that
enables these properties is a dual reactor system in which the first reactor is a liquidphase loop reactor and the second is an in-series gas-phase fluidized bed reactor. A
diagram of the Borstar process is shown in Figure 3.17. This process can produce film
extrusion products with a wide range of densities and MWDs and bimodal MWD
HDPE, MDPE, and LLDPE products. The process utilizes a proprietary catalyst which
is fed to a small liquid-phase loop reactor acting as a pre-polymerizer. A morphologically advantaged catalyst-polymer particle is created for subsequent feeding to a
large, liquid-phase loop reactor. No further catalyst is fed throughout the process. The
loop reactor uses propane as the reaction diluent and is operated at conditions under
Ethylene Polymerization Processes and Manufacture of Polyethylene
Catalyst Prepolymerization
reactor
Loop
reactor
95
Gas
Flash
tank
Gas-phase
fluidized
bed reactor
HX
Coolant
pipe
CW
Polymer
Ethylene
Comonomer
Hydrogen
Diluent
To polymer
degassing and
pelletizing
Purification
Cycle gas compressor
Figure 3.17 Borstar hybrid process capable of producing linear PE [84].
which propane is in a supercritical state. These conditions reduce polymer solubility
in the diluent. Coupled with the tailored catalyst-polymer particle, these conditions
reduce polymer sintering and wall fouling. The polymer-diluent stream leaving the
loop reactor is depressurized to vaporize the propane diluent, which is recycled back
to the loop reactor while the polymer with active catalyst is conveyed to the in-series
gas-phase reactor. Effective control of distinctly different conditions in each reactor is critical in enabling the production of high-value bimodal products, as was the
case in the Spherilene process. The low MW fraction is made in the loop reactor and
the higher MW fraction in the gas-phase reactor. The comonomer content in each
segment is regulated as is the production split between each reactor for optimum
properties.
3.3 Resin Property and Reactor Control in
Catalytic Polymerization Reactors
Reactor conditions needed to make a given product are primarily a function of the
catalyst. Therefore, a basic guide to resin property control will be discussed here for
each of the three major catalyst families: Ziegler-Natta, Phillips chromium oxide, and
metallocene. Keep in mind that within a given family, there may be several versions,
such as different levels of aluminum alkyl reduction or different catalyst calcination temperatures. These versions will also affect resin properties. Note that both the
effects of supporting the catalyst and of the phase at which the reaction is carried out
will be large.
3.3.1
Production Rate
Catalyst feed rate is the primary driver of production rate. The ability to remove the heat
of reaction likely limits the production rate and thus the rate of the catalyst addition.
96
Handbook of Industrial Polyethylene and Technology
3.3.2 Catalyst Productivity
Ethylene partial pressure is usually the main driver of catalyst productivity (kg polymer/kg catalyst). Cocatalyst feed rate, used with Z-N and metallocene and other singlesite catalysts, will affect catalyst productivity. The feed ratio of cocatalyst to catalyst is
often adjusted to the lowest level at which productivity is maximized. However, this
ratio can also affect resin properties such as MW. In instances where production rate is
to be maximized, the cocatalyst/catalyst ratio is adjusted accordingly.
3.3.3 Reactor Pressure
The desired ethylene partial pressure is a major driver of reactor pressure in reactors involving a gas phase. In these reactors, total pressure needs to be high enough
to fit the ethylene and other components into the reactor at their required partial
pressures. Another consideration is minimization of reactor vent flow rate. In solution reactors, higher reactor pressure may reduce the energy needed for post-reactor
separations.
3.3.4 Crystallinity
3.3.4.1 Ziegler-Natta and Metallocene Catalysts
For Ziegler-Natta and metallocene catalysts, the dominant driver increasing crystallinity and polymer density is comonomer concentration. The comonomer concentration is often expressed as comonomer/ethylene molar ratio. The concentration of
comonomer required in the reactor to make a given product is typically substantially
lower for metallocene catalysts. Further, the amount of comonomer incorporated in
the polymer at the same polymer density will be a little less for metallocene catalysts.
Hydrogen concentration is known to affect comonomer response of the catalyst. The
dominant driver of crystallinity with chromium oxide catalysts is also comonomer/
ethylene molar ratio.
3.3.5 Molecular Weight
Molecular weight is inferred from measurement of melt index. A high melt index
means low MW. For Ziegler-Natta catalyst, the primary means of controlling melt index
employs hydrogen as a chain transfer agent. To increase the melt index, increase hydrogen/ethylene molar ratio. In order to make a polymer with high melt index with some
Z-N catalysts, this ratio can exceed 1 mol/mol. For chromium oxide catalyst, the melt
index is increased by increasing the reactor temperature. Hydrogen also increases melt
index, but the effect is weak. Higher ethylene partial pressure may decrease melt index
[45]. For metallocene catalyst, hydrogen is used to increase melt index. The response
of MW to hydrogen is substantially greater for metallocene catalysts than for other
catalysts [85–87]. For dicyclopentadienyl zirconium catalysts, and perhaps other metallocenes too, increasing the concentration of the cocatalyst methylaluminoxane (MAO)
increases melt index [88].
Ethylene Polymerization Processes and Manufacture of Polyethylene
3.4
97
Economics
Licensors’ claims about being low cost are often biased. A few independent sources,
such as Nexant’s CHEMSYSTEMS Process Evaluation/Research Planning Program, IHS
Chemical Economics Handbook (formerly SRI Consulting), and Intratec Technology
Economics Report, periodically assess the costs of various polymerization processes.
These assessments are available only to subscribers and not to the public.
The following economics comparison [18, 45] has been published by TVK in Europe.
Two high-pressure processes and five low-pressure processes for production of PE were
considered. The high-pressure autoclave process required three trains; the others used
one train. The basis for the economic assessment is production of 400,000 tonnes/year
of PE at the US Gulf Coast in 2008. Total investment and operating cost are shown in
Tables 3.6 and 3.7, respectively.
Investment and operating costs for new LDPE facilities clearly show the advantage
of large-capacity, single-train tubular reactors over in-parallel, multi-train autoclave
reactors [18, 89]. Projected total investment costs for a single-train reactor autoclave
system of 135,000 tonnes a year is about $215 million and for three reaction trains of
135,000 tonnes per year each (405,000 tonnes per year total capacity) is $535 million. A
single-train tubular reactor system of 405,000 tonnes per year investment is projected
at $375 million or about 70% of a three-reactor autoclave system of identical capacity.
In terms of operating costs, the large-scale single tubular system has only a 4%
advantage over a three-reactor autoclave system. The obvious reason is that over 90%
of costs are due to monomer costs which are similar regardless of type or number of
reactors required. Power costs are higher for tubular systems but they have an advantage in capturing the heat of polymerization and generating steam for use in other areas
of the process.
Table 3.6 Investment costs for major processes to manufacture PE at 400,000 tonnes/year.
Technology:
High-pres High-pres
Gas
Gas
Slurry Slurry
autoclave
tubular Solution
phase
phase
loop
CSTR
Product:
LDPE
LDPE
LLDPE
LLDPE
HDPE
HDPE HDPE
Cost, millions of US dollars
Inside battery
limits
231.6
130.1
126.4
93.8
94.7
114.1
126.8
Outside battery
limits
142.3
118.4
115.3
102.6
102.7
117.7
116.4
Other project
costs
160.5
125.1
131.3
116.5
114.2
123.9
127.3
Total investment 534.4
373.6
373
313
312
352
371
Specific investment, USD/
tonne poly
934
933
782
779
879
926
1320
98
Handbook of Industrial Polyethylene and Technology
Table 3.7 Operating costs for major processes to manufacture PE.
Technology:
Product:
High-pres High-pres
Gas
Gas
Slurry Slurry
autoclave
tubular Solution
phase
phase
loop
CSTR
LDPE
LDPE
LLDPE
LLDPE
HDPE
HDPE
HDPE
Cost, US dollars/tonne
Ethylene
1354
1347
1223
1222
1326
1324
1334
Comonomer
0
0
176
159
25
24
25
Catalysts &
chemicals
8
5
47
14
27
42
22
Additives
14
14
13
15
18
18
18
Total raw
materials
1376
1366
1459
1410
1396
1408
1399
Power
51
58
15
22
25
30
26
Steam
0
34
16
2
2
6
13
Other utilities
9
7
8
10
6
11
8
Total utilities
60
31
39
34
33
47
47
Operating cost
1436
1397
1498
1444
1429
1455
1446
Direct cash cost
22
14
14
12
12
14
14
Allocated cash
cost
27
18
17
15
15
16
17
Total fixed cost
49
32
31
27
27
30
31
Total cash cost
1485
1429
1529
1471
1456
1485
1477
Low-pressure gas-phase processes have inherent advantages over other processes
because the process is simple, with an absence of many of the unit operations that are
required in liquid-phase processes. This simplicity results in lower investment and operating cost advantages. The advantage is larger when comparing single fluidized-bed systems to multiple reactor systems such as those used in low-pressure hybrid technologies
or staged liquid-phase systems. Also, single gas-phase fluidized bed systems can be built
to high production capacities (in the order of 600,000 tonnes per year) for economies of
scale. However, other systems can have product advantages over a single gas-phase reactor. Their products may have less residue resulting from removal of low MW species by
diluents. Or they may produce bimodal or broader MWD products from the use of two
or more reactors. Operating costs are very dependent on monomer efficiency, which
varies little among the advanced technologies. Use of high-cost advanced catalysts
for product or process advantages can be a factor in overall operating costs. However,
because of fewer process operations, lower energy needs and large single-reactor plant
capabilities, gas-phase fluidized beds should have the most competitive operating costs.
Ethylene Polymerization Processes and Manufacture of Polyethylene
99
Table 3.8 Single-train capacity and worldwide capacity of PE processes.
Company
Single-train
capacity
Wordwide product
capacity
thousands of
tonnes/y
thousands of
tonnes/y
year
Autoclave
150
8000
2015
Tubular
450
13000
2015
600
18000
2007
2450
2009
Process
High-Pressure LDPE
Gas-Phase
Univation
Fluidized Bed
LLDPE
Nova Chemical
SCLAIRTECH
350
Nova Chemical
Advanced SCLAIRTECH
425
Dow Chemical
Dowlex
300
SABIC
Stamicarbon Compact
75
650
2002
9000: single &
dual loop
2011
HDPE
Chevron-Phillips
Chemical
MarTECHTM Single-Loop
Slurry
550
Chevron-Phillips
Chemical and Total
Petrochem
MarTECHTM Advanced
Dual-Loop Slurry
550
INEOS
Innovene S
345
2200
2011
LyondellBasell
Hostalen Bimodal
400
7000
2009
LyondellBasell
Hostalen Advanced
Cascade Process
450
1290
2014
Mitsui
CX
3600
2005
The gas phase process is lowest in investment. Only the high-pressure tubular process has lower operating expense. The high-pressure autoclave process incurs by far the
highest investment; construction of new autoclave plants is rare, but many old facilities
continue to be operated. Among low-pressure processes, the solution process is highest
in both investment and operating cost. Among the slurry HDPE processes, the loop
process has a substantial advantage in investment cost, but a modestly higher operating
cost than the CSTR process.
Investment cost is highly dependent on the capacity of a single-train unit. The
maximum capacities of several processes, as well as worldwide capacity, are compared in Table 3.8. The gas-phase fluidized-bed technology of Univation has the lead
100
Handbook of Industrial Polyethylene and Technology
in single-train capacity at 600,000 tonnes per year. In the last two decades, ChevronPhillips has made great leaps in single-train capacity, placing them close behind
Univation. Capacities of the tubular high-pressure technology, the Hostalen Advanced
Cascade Process, and the Advanced SCLAIRTECH process are about 100,000 tonnes
per year smaller.
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53. Alarcon, F., Differentiated Dual PE Technologies Platform, presented at: Maack PEPP
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58. BP Licenses Showa Denko HDPE Process, Chemical Week, p. 21, May 2, 2001.
59. Dow Chemical polyolefins expert bio, www1.intota.com/expert-consultant.asp?bioID=
774846, 2007.
60. LyondellBasell, Licensed Polyolefin Technologies and Services. Hostalen, www.lyondellbasell.com/technology, undated.
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php?s=43&item=708&printable, March 23, 2009.
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Company, at www.lyondellbasell.mediaroom.com/index.php?s=43&item=898&printable,
January 25, 2012.
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Phase, Ph.D. Thesis, University of Twente, The Netherlands, 2009.
65. LyondellBasell, Matagorda Slurry Technology, www.lyondellbasell.com/techlit/techlit/
Tech_Topics/Plant_Profiles/Matagorda.pdf/, undated.
66. Mitsui Chemicals, Mitsui CX Process: High and Medium Density Polyethylene Technology,
http://www.mitsuichem.com/technol/license/pdf/1.CX_PROCESS.pdf, undated.
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page_9.html, 2011.
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Content in an Olefin Solution Polymerization Process, US Patent 6420516, assigned to The
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70. Dyer, G., Manuel, L., and Robinson, D., The SCLAIR Story, Canadian Chemical News,
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4
Types and Basics of Polyethylene
Rajen M. Patel
The Dow Chemical Company, Freeport, Texas, USA
Contents
4.1 Introduction ............................................................................................................................... 106
4.2 Low Density Polyethylene (LDPE) ......................................................................................... 109
4.3 Ethylene Vinyl Acetate (EVA) Copolymer............................................................................. 112
4.4 Acrylate Copolymers ................................................................................................................ 115
4.5 Acid Copolymers....................................................................................................................... 116
4.6 Ionomers..................................................................................................................................... 117
4.7 High Density Polyethylene (HDPE) ....................................................................................... 118
4.8 Ultra-High Molecular Weight HDPE (UHMW-HDPE) ..................................................... 120
4.9 Linear Low Density Polyethylene (LLDPE) ........................................................................... 121
4.10 Very Low Density Polyethylene (VLDPE) ............................................................................. 123
4.11 Single-Site Catalyzed Polyethylenes........................................................................................ 124
4.12 Olefin Block Copolymers (OBC) ............................................................................................ 133
4.13 Concluding Remarks ................................................................................................................ 136
Acknowledgments .............................................................................................................................. 136
References ............................................................................................................................................ 137
Abstract
Polyethylene (PE) has a long and rich history of product, process and fabrication innovations
to meet growing market needs over the last 75 years. Many different types of ethylene homopolymer and copolymer resins have been developed with a broad range of performance to meet
the requirements of a variety of applications, and as a result PE is the highest volume plastic available today. This chapter will cover many different types of ethylene homopolymer and
copolymer resins widely used today, their molecular characterization, and structure-properties
relationships. The chapter will also cover fundamentals and a variety of applications of various
types of ethylene homopolymer and copolymer resins.
Keywords: Polyethylene, LDPE, LLDPE, VLDPE, HDPE, HMW HDPE, UHMW HDPE, EVA,
acid copolymers, acrylate copolymers, ionomers, ECO copolymers, plastomers, elastomers,
metallocene LLDPE, EPDM, olefin block copolymers
Corresponding author: rmpatel@dow.com
Mark A. Spalding and Ananda M. Chatterjee (eds.) Handbook of Industrial Polyethylene and Technology, (105–138)
© 2018 Scrivener Publishing LLC
105
106
Handbook of Industrial Polyethylene and Technology
4.1 Introduction
Polymers are typically classified as thermoplastic or thermoset polymers, among several classification systems. Thermoplastic polymers can be melted and reprocessed
multiple times at high temperatures whereas thermoset polymers cannot be melted
and reprocessed due to the presence of cross-links between molecules. This classification of polymers is based on the processing aspects of polymers. However, polymers
can also be classified as either amorphous or semicrystalline based on structural order.
Amorphous polymers have zero degree of crystallinity and molecules are packed in a
random manner (no order). Semicrystalline polymers typically have 1 wt% to about 80
wt% crystallinity, with the rest being an amorphous phase (assuming two-phase model).
The crystalline phase consists of crystallites which are typically in a chain folded lamellar form. Note that polymers cannot be 100 wt% crystalline due to their high molecular
weight and therefore a high degree of entanglements. Amorphous and semicrystalline
polymers can be further classified based on whether their glass transition temperature
is above or below room temperature (~25 °C). The glass transition temperature (Tg) is a
key property of the amorphous phase of polymers. It is the temperature at which amorphous phase transitions from glassy (rigid) to a rubbery behavior. Below Tg, molecules
in the amorphous phase are not mobile, hence the polymer is rigid (glassy). Above Tg,
molecules in the amorphous phase become mobile, leading to a rubbery behavior.
Semicrystalline polymers having glass transition temperatures (Tg) below room temperature include PE (ethylene-based polymers with >50 mol% units derived from ethylene) and polypropylene (PP) (propylene-based polymers with >50 mol% units derived
from propylene) resins; both referred to as part of the polyolefin family. Since the
amorphous phase of polyolefins is in a rubbery state, they exhibit much lower modulus
compared to amorphous (e.g., polystyrene) or semicrystalline polymers (e.g., polyethylene terephthalate) having the amorphous phase in a glassy state due to Tg above room
temperature. Modulus, tensile strength, elasticity, and polarity of PE can be dramatically changed by changing the degree of crystallinity of the resins via copolymerization
with a wide range of comonomers. As a result, PE is a very versatile polymer exhibiting
a very broad range of properties, ranging from stiff/brittle to ductile/tough to elastomeric, depending upon the degree of crystallinity. This tremendous flexibility in tuning
mechanical properties has led to a wide range of applications for PE. Therefore, PE is
commercially very important and is the highest volume thermoplastic produced, having a rich history of major innovations in products, processes, and breadth of applications [1, 2]. Global demand for PE in 2014 was about 85 million metric tons (187 billion
pounds). Global demand is expected to increase to about 105 million tons (230 billion
pounds) by 2019.
Polyethylene is composed of only carbon and hydrogen (with some notable exceptions) which can be combined in a number of ways to make different types of PEs.
There are various molecular architectures that have been commercialized over the last
70 years to make different types of ethylene homopolymers and copolymers. These can
be generally grouped into eleven major types:
Low density PE (LDPE);
Ethylene vinyl acetate copolymer (EVA);
Types and Basics of Polyethylene 107
Acrylate copolymers such as ethylene methyl acrylate (EMA), ethylene
ethyl acrylate (EEA) and ethylene n-butyl acrylate (EnBA);
Acid copolymers such as ethylene acrylic acid (EAA) or ethylene methacrylic acid (EMAA) copolymers;
Ionomers;
High density PE (HDPE);
Ultra high molecular weight high density PE (UHMWPE);
Linear low density PE (LLDPE);
Very low or ultra low density PE (VLDPE or ULDPE);
Homogeneous PE produced via single-site catalysts (including metallocene) (polyolefin plastomers, polyolefin elastomers, mLLDPE,
mVLDPE) [3];
Olefin block copolymer (e.g., INFUSE OBC) [4].1
There are still other minor types of ethylene copolymers including:
Chlorinated PE (CPE) made via chlorination of HDPE powder;
Ethylene/carbon monoxide (ECO) copolymers, which are photodegradable and made using a high pressure free-radical polymerization process,
and are used to make six-pack loop carriers for beverage;
Ethylene/vinyl trimethoxy silane (VTMOS) copolymers, used in the wire
and cable industry for moisture curing;
Ethylene/maleic anhydride (MAH) copolymers made using the high
pressure process and used mainly as a compatiblizer;
Specialty ter-polymers made in a high pressure process such as ethylene/acrylic acid/acrylate; ethylene/butyl acrylate/carbon monoxide; ethylene/vinyl acetate/carbon monoxide; ethylene/butyl acrylate/glycidyl
methacrylate, etc.;
Cyclic olefin copolymers (COC), made from ethylene and norbornene
comonomer using a single-site catalyst;
MAH-grafted-PE resins, used in tie-layers, tying ethylene vinyl alcohol
(EVOH) or polyamide to a PE layer in multilayer co-extruded barrier
films. Such resins typically contain about 0.07% to 0.15% of MAH in
ready-to-use formulations to provide good adhesion to polyamide and
EVOH.
These eleven major types of ethylene homopolymers and copolymers can
be grouped into three categories: high pressure process resins, heterogeneous PEs catalyzed by Ziegler-Natta (Z-N)/Chromium Oxide catalysts,
and homogenous PE catalyzed by single-site (including metallocene)
catalysts. This is shown in Figure 4.1.
Most of the ethylene polymer types listed above consist of a polymer backbone composed of carbon and hydrogen with different types of branches from the backbone
[5]. The branches range from alkyl in LDPE, LLDPE, and VLDPE to acetoxy and
1
INFUSE is a trademark of The Dow Chemical Company (“Dow”) or an affiliated company of Dow.
108
Handbook of Industrial Polyethylene and Technology
High pressure
Z-N and chromium
oxide catalyzed
LDPE
EVA
Acrylate copolymers
acid copolymers
ionomers
HDPE
UHMWPE
LLDPE
ULDPE/VLDPE
Plastomer/elastomers &
mLLDPE/mVLDPE
Olefin block copolymer
(OBC)
Single site
catalyzed
Figure 4.1 Grouping of eleven types of ethylene homopolymers and copolymers into three buckets.
alkyl branches in EVA to carboxyl branches in EAA or EMAA. The degree and type
of branches control the degree of crystallinity of PE by introducing defects in regular
chain architecture, thereby affecting the solid-state properties of PEs. Molecules in the
crystalline phase typically are packed tightly together in a regular 3-dimensional order.
Hence, the crystalline phase typically has a higher density than the amorphous phase.
As a result, the density of PE is directly related to the degree of crystallinity and, in fact,
nonpolar PEs (LDPE, LLDPE and HDPE resins) are classified and specified by resin
density. The density of PE can be used to estimate the wt% and vol% crystallinity in PE
as follows:
Wt% Crystalinity =
c
a
c
a
Vol% Crystalinity =
c
100
(4.1)
100
(4.2)
a
a
where is the nonpolar PE density, a = 0.855 g/cm3 for nonpolar PE, and c = 1.00 g/cm3;
is the amorphous phase density and c is the crystal density.
a
The ability to control types and degree of branching via the incorporation of different comonomers allows one to make PE from nonpolar to polar, from stiff and rigid
plastics to soft and flexible elastomers, from having a high melting point to a low melting point, etc. As a result, a wide range of properties are obtained from the abovementioned ethylene homopolymer and copolymer types, making ethylene polymer one
of the most versatile plastics.
Even though there are eleven major types of ethylene homopolymers and copolymers, generally PE resins are classified into three broad categories; LDPE, LLDPE
and HDPE. Global demand for these three broad categories of PE in 2014 is shown
in Figure 4.2. HDPE represents the largest fraction of the PE produced, followed by
Types and Basics of Polyethylene 109
2019 Global demand
2014 Global demand
LDPE
LDPE
HDPE
HDPE
LLDPE
LLDPE
Growth forecast
24%
Total global demand = 84.8
Million metric tons
Total global demand = 105.4
Million metric tons
Figure 4.2 Global demand for the three broad categories of PE. (Source: IHS, Polyolefins conference, 2015).
Figure 4.3 Molecular structures, properties and general applications of three broad categories of PE resins.
LLDPE and LDPE. Global demand for LLDPE resin is expected to increase as a percentage, mainly at the expense of LDPE.
Molecular structures, properties, and general applications of these three broad categories of PE resins are shown in Figure 4.3. The demand for PE can also be classified
by the type of conversion process used to make the final end-use part. This is shown in
Figure 4.4.
The eleven types of ethylene homopolymers and copolymers will be described in
more detail in the following sections.
4.2 Low Density Polyethylene (LDPE)
Low density polyethylene was accidentally discovered at Imperial Chemical Industries
(ICI) by Fawcett and Gibson as a waxy solid powdery substance found while
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Handbook of Industrial Polyethylene and Technology
Rotomolding
2%
Blow molding
12%
Wire & cable
2%
Fiber
1%
Raffia
1%
Other
7%
Film & sheet
52%
Extrusion coating
3%
Pipe & profile
7%
Injection molding
13%
Figure 4.4 PE demand by conversion processes (2013). (Source: The Dow Chemical Company)
investigating high-pressure and high-temperature reactions of ethylene and benzaldehyde. Researchers at ICI realized after months of experimentation that oxygen had to
be present to initiate the conversion of ethylene to PE under high pressure. The presence of oxygen leads to the formation of free-radical initiators (peroxides) that subsequently decompose to start the free-radical addition polymerization of LDPE. ICI filed
a series of patents to claim the inventions, claiming continuous production of LDPE
under pressures greater than 500 atmospheres and temperatures greater than 100 °C,
preferably in the presence of 0.01 wt% to 5 wt% oxygen [6]. Polymerization of ethylene
to PE is a highly exothermic reaction, releasing approximately 3370 J/g of ethylene.
Removal of the exothermic heat of polymerization determines the percent conversion
of ethylene to PE per pass.
Low density polyethylene is made using either an autoclave or a tubular process,
using a free radical polymerization process initiated by organic peroxides. The autoclave reactor is a continuous stirred-tank reactor (CSTR) with an agitator designed
to promote good mixing. Autoclave reactors are essentially adiabatic and rely on back
mixing of the hot reactants with the cold incoming ethylene feed stream. The heat of
polymerization is removed with effluent from the reactor to keep the reaction stable.
Autoclave reactors typically operate at 1200 to 2400 bar. Tubular reactors are non-adiabatic reactors that allow additional heat removal via a coolant in the reactor jacket
(containing high pressure hot water) in addition to heat removal via cold ethylene feed.
High pressure hot water (120 °C to 180 °C) in the reactor jacket prevents fouling of the
reactor wall by preventing the crystallization of LDPE. Tubular reactors typically operate at 2000 to 3200 bar. Conversion (ethylene to PE) per pass is typically about 14 to
20 wt% for the autoclave process and about 21 to 38 wt% for the tubular process.
Types and Basics of Polyethylene 111
The free radical polymerization process leads to both long-chain branching (via
chain transfer to another molecule) and short-chain branching (via back-biting mechanism) in LDPE resins. LDPE has ratio of about ten short-chain branches to every longchain branch. LDPE density (crystallinity) is controlled by the reactor temperature and
pressure. A small reduction in density can also be achieved by adding a small amount
of comonomer like propylene or 1-butene (both of which also act as a chain transfer
agent reducing molecular weight). Higher reactor pressure (via increasing propagation rates) and lower reactor temperature (via decreasing back-biting rates) increases
LDPE density, and conversely, higher reactor temperature and lower reactor pressure
decreases LDPE density. Due to the limitations in the reactor pressure and temperature
range, LDPE resins can be commercially produced only in the density range of about
0.915 to 0.935 g/cm3 (about 45 to 59 wt% crystallinity).
A key molecular feature of LDPE resins is the presence of a high level of “tree-like”
long-chain branching (LCB), typically about 2 to 3 LCB/1000 carbon atoms. It has been
proposed that the length of LCB is greater than the critical entanglement molecular
weight [7]. For PE, this equates to 270 carbon atoms or greater [5]. As a result of a
high level of LCB, LDPE resins exhibit excellent extrusion processability (due to a high
degree of shear thinning) and melt strength properties but inferior impact, tear, abrasion, and environmental stress-cracking resistance (ESCR) properties. Though crystalline domains control low deformation (strain) properties of PE, such as modulus and
yield stress, it has been clearly established by many researchers that large deformation
(strain) properties, such as stretchability, impact, tear, failure processes, etc., are also
controlled by the amorphous region, particularly by tie chains , the amorphous chains
that bridge adjacent lamellae [8, 9]. Inferior mechanical properties of LDPE are caused
by the compact/smaller size of LDPE molecular coil (lower radius of gyration) in the
molten state due to the presence of a high degree of LCB, resulting in reduced tie-chain
concentration [10]. LDPE resins are widely used as a blend component with LLDPE
resin to improve bubble stability, gauge uniformity, optics, and rates in the blown film
process [11]. LDPE resins having fractional melt indices are also widely used as a blend
component in collation shrink film (shrink bundling) applications because of the need
for cross-direction (CD) film shrinkage (~10–25%). Due to their very high melt strength
as a result of the high level of long-chain branching in the resin, LDPE resins are especially suitable for extrusion coating and foam applications. The high melt strength of
LDPE resins, especially autoclave LDPE resins, allows lower neck-in (change in the
width of the molten web) and high line speed during the extrusion coating operation.
Autoclave-produced LDPE has broader molecular weight distribution than tubular
LDPE [12, 13]. This is shown in Figure 4.5. Molecular weights were measured using a
low angle laser light scattering detector to accurately measure the absolute molecular
weight of the LDPE resins containing high levels of long-chain branching. Autoclave
LDPE also has a higher level of long-chain branching and, hence, a higher melt strength
than tubular LDPE at similar melt indices. Melt strength of an autoclave resin is compared with a tubular grade in Figure 4.6. Melt strength was measured using a Goettfert
Rheotens apparatus. Autoclave LDPE is generally the preferred resin for extrusion coating applications due to its lower neck-in performance.
Due to its high melt strength and, in particular, strain hardening in elongational flow
(increase in elongational viscosity with deformation/strain) during foaming process,
Handbook of Industrial Polyethylene and Technology
dWf / {dLog molecular weight (by LALLS)}
112
1
0.8
Autoclave
Tubular
Tubular (1.9 MI)
0.6
0.4
Autoclave (2.3 MI)
0.2
0
2
3
4
5
6
Log molecular weight (by LALLS)
7
8
Figure 4.5 Molecular weight distributions of autoclave and tubular LDPE resins at similar melt indices.
20
190 °C
18
Autoclave (2.3 MI)
Melt strength (cN)
16
17.0 cN
Autoclave
Tubular
14
12
10
Tubular (1.9 MI)
8
7.6 cN
6
4
2
0
0
50
100
150
200
Velocity (mm/s)
250
300
350
Figure 4.6 Melt strengths of autoclave and tubular LDPE resins at similar melt indices.
LDPE resins allow the production of foams with nearly 100% closed cells,. Blown films
made from LDPE resins with a relatively narrow molecular weight distribution and low
melt strength, typically made via the tubular process, exhibit very good optics and are
used in high clarity applications such as in soft bread packaging. Other applications of
LDPE include trash bags, food storage bags, food wrap film, and injection molded articles. LDPE, as well as silane copolymers or silane-grafted LDPE, is widely used in wire
and cable insulation applications due to being a “clean” resin having no metal/catalyst
residues, yielding excellent insulation properties (high electrical breakdown potential).
4.3
Ethylene Vinyl Acetate (EVA) Copolymer
Ethylene vinyl acetate is a random copolymer of ethylene and vinyl acetate monomers
made using a high pressure free-radical polymerization process. Vinyl acetate (VA)
wt% crystallinity
Types and Basics of Polyethylene 113
50
45
40
35
30
25
20
15
10
5
0
y = –1.034x + 47.7
R2 = 0.9936
0
5
10
15
20
25
30
wt % VA
Figure 4.7 Wt% crystallinity of EVA resins as a function of wt% VA content.
(Courtesy of Dr. Ken Anderson, Celanese)
content has two fundamental effects that influence the properties of EVA copolymers.
The first effect is to disrupt the crystallinity formed by long ethylene sequences of the
copolymer. The second key effect of VA content is the polar nature resulting from acetoxy side chains. EVA resins have a lower degree of crystallinity and a lower melting
point compared to high pressure LDPE resin; both properties are controlled by the
VA content. Weight percent vinyl acetate in the EVA copolymer typically ranges from
about 4 to 40%. EVA polymers with VA content higher than 40 weight percent create
pellet handling issues; e.g., pellet blocking and stickiness. At about 45% VA content
by weight, depending on molecular weight and molecular weight distribution, EVA
becomes completely amorphous (non-crystalline). A plot of the crystallinity of EVA
resins as a function of VA content is shown in Figure 4.7.
EVA copolymers have been commercially available. EVA is the largest volume polar
ethylene copolymer used today. Major suppliers of EVA include DuPont (ELVAX ),
ExxonMobil (Escorene
and Nexxstar ), Celanese (Ateva ), LyondellBasell
(Ultrathene ), Arkema (EVATANE ), etc.2 As the vinyl acetate level is increased, the
degree of crystallinity (and hence, stiffness) decreases but density increases. This is due
to bulky acetoxy side groups in EVA which increases the amorphous phase density
and therefore overall density even though crystallinity is lower. As the vinyl acetate
level increases, the EVA copolymer becomes more polar. Melt indices of EVA resins
typically range from 0.35 to 2500 dg/min (190 °C, 2.16 kg). EVAs have a high degree
of long-chain branching which make them easy to extrude and melt process. However,
they are thermally less stable compared to other ethylene copolymers due to presence
of weaker carbon-oxygen bonds which can lead to the generation of acetic acid (vinegar) and the formation of cross-linked gels during high temperature melt processing.
Therefore, EVA resins are generally melt processed below 230 °C. Also, the high degree
of long-chain branching in EVA resins leads to inferior puncture, dart, and tear properties compared to linear PE resins (e.g., ATTANE from The Dow Chemical Company)
at similar crystallinity.
2
ELVAX is a trademark of E.I. du Pont de Nemours and Company or its affiliates; Escorene and Nexxstar
are trademarks of ExxonMobil Corporation; Ateva is a trademark of Celanese Corporation; Ultrathene is
a trademark of LyondellBasell; EVATANE is a trademark of Arkema.
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Handbook of Industrial Polyethylene and Technology
EVAs are widely used in film applications that require a lower crystallinity and/or a
lower melting point; e.g., sealants, barrier shrink bags, high clarity films, applications
requiring low temperature toughness such as ice bags, stretch hood films, and thermal
lamination films. EVA resins having wt% VA levels from 7 to 18% VA are widely used
in the core layer of stretch hood film formulations, depending upon on-pallet stretch
level requirements. Both EVA and LDPE resins exhibit very low hot-tack strength as a
sealant. This is due to a very high degree of long-chain branching which slows down
diffusion across the molten sealant interface significantly. EVA resins are also used as
a tie-layer resin for tying polyvinylidene chloride (PVDC) to a PE layer in multilayer
co-extruded barrier films. EVA resins having 18 to 20% VA content and 15 to 20 melt
index are extrusion coated (8 to 15 micron thick) onto a biaxially oriented polypropylene (BOPP) film and such films are widely used as thermal lamination films for
laminating onto glossy printed paper for use as magazine covers. EVA resins are also
widely used in wire and cable applications such as heat shrinkable insulation, semiconductive insulation jackets, and flame retardant formulations [14]. EVA resins are
used to make cross-linked foams used as mid-soles in footwear applications as well
as injection molded articles. Other uses for EVAs include medical applications such
as IV bags and cryo-containers. Ease of sterilization of EVAs using gamma radiation
makes them very suitable in such medical applications. High melt index EVA resins
(melt index typically greater than 50) having 18 to 33 wt% VA levels are widely used for
hot melt adhesive (HMA) applications. EVA-based hot melt adhesives are formulated
with waxes and tackifiers and are used in applications such as case and carton closing,
labeling, lamination, book binding, furniture and wood working. EVA-based hot melt
adhesives exhibit poor thermal and oxidative stability leading to shorter pot life due to a
rapid rise in viscosity, yellowing, and gel/char formation (which can clog the applicator
nozzle). EVA-based hot melt adhesives also exhibit inferior low temperature adhesion
due to higher Tg of EVA resins (Tg of about −30 °C).
Figure 4.8 compares the glass transition temperature of EVA versus single-site catalyzed ethylene/1-octene polyolefin plastomer (POP) and polyolefin elastomers (POE)
ENGAGE /AFFINITY resins from The Dow Chemical Company.3 Weight % crystallinity at 23 °C was obtained by dropping perpendicular at 23 °C in the differential
scanning calorimeter (DSC) melting endotherm during the heat of fusion integration
(area under the curve). EVA resins exhibit a higher Tg compared to POP/POE resins at
a given crystallinity. The Tg of POP/POE resins increase with the increase in resin density/crystallinity. Note that above 0.91 g/cm3 density, Tg cannot be observed as a step
change in heat capacity in a DSC melting endotherm.
EVA copolymers at 28 to 33 wt% VA content and high melt indices (40 to 50 dg/min)
are widely used to make an encapsulant film (0.3 to 0.8 mm thick) for photovoltaic
(solar) cells. Such soft thick cross-linked encapsulant films exhibit excellent light transmission and help to protect fragile silicon wafer solar cells. However, under exposure to
atmospheric water and/or ultraviolet radiation, EVA will decompose to produce acetic
acid, lowering the pH and increasing the surface corrosion rates of embedded devices
leading to deterioration of module performance. Also, the glass transition of EVA,
3
ENGAGE and AFFINITY are trademarks of The Dow Chemical Company (“Dow”) or an affiliated
company of Dow.
Types and Basics of Polyethylene 115
0
0
10
20
30
40
50
–10
Tg (ºC) by DSC
–20
POP & POE
EVA
28% VA
–30
25% VA 18% VA 15% VA
0.909 g/cc
0.897 g/cc
–40
y = 0.7316x – 62.746
–50
–60
0.876 g/cc
0.863 g/cc
0.879 g/cc
0.857 g/cc
–70
wt% Crystallinity at 23 ºC
Figure 4.8 Glass transition temperature (Tg) of EVA and ethylene/octene POP/POE (AFFINITY/
ENGAGE from The Dow Chemical Company) resins as a function of crystallinity at 23 °C, measured
using differential scanning calorimetry (DSC). %VA levels are shown for the EVA resins and measured
densities are shown for the 1-octene-based POP/POE resins.
measured using dynamic mechanical analysis, begins at about −15 °C. Temperatures
lower than this can be reached for extended periods of time in some climates. Because of
increased modulus below the glass transition temperature, a module may be more vulnerable to damage if a mechanical load is applied by snow or wind at low temperatures.
The Dow Chemical Company has developed ethylene/alpha-olefin copolymer-based
encapsulant film under the trade name of ENLIGHT . The ENLIGHT encapsulant
film provides improved moisture and UV resistance, high volume resistivity and low
leakage current compared to traditionally used EVA films, enabling long-lasting protection to significantly extend solar module service life.
EVA copolymers also provide an alternative to polyvinyl chloride (PVC) in medical
applications because of migratory plasticizer and incineration issues of the PVC resins.
EVA copolymers are also used as a carrier resin to make compounds and masterbatches
as well as in automotive applications (e.g., sound dampening materials), often with calcium carbonate fillers
4.4
Acrylate Copolymers
Acrylate copolymers are made using a high pressure free-radical polymerization process typically using autoclave reactors. There are three types of acrylate copolymers
depending upon the type of acrylate comonomer used; ethylene methyl acrylate
(EMA), ethylene ethyl acrylate (EEA), and ethylene butyl acrylate (EBA). The glass
transition temperature of acrylate copolymers is in the order of EMA > EEA > EBA at
a given crystallinity. Acrylate comonomers add polar functionality to ethylene polymers, decreasing the crystallinity and melting point. Acrylate copolymers made using a
tubular reactor results in a more heterogeneous molecular structure and a higher crystalline melting temperature compared to those made with autoclave reactors. Acrylate
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Handbook of Industrial Polyethylene and Technology
copolymers containing up to 35 wt% comonomer level are commercially available. Due
to their low glass transition temperature, EBA resins exhibit superior low temperature
toughness and impact resistance. Acrylate copolymers have a high degree of long-chain
branching (and exhibit high melt strength), similar to all PE resins made using a high
pressure free-radical polymerization process. A key feature of acrylate copolymers is
that they have very good thermal stability (similar to LDPE) and, hence, can be melt
processed at very high temperatures, up to 325 °C. This is in contrast to EVA resins
which are thermally not as stable and are melt processed below 230 °C. Acrylate copolymers are also non-corrosive to equipment and are not hygroscopic. Acrylate copolymers, especially EMA, are used in extrusion coating and lamination applications as
extrusion coated tie layers for adhesion to printed oriented polypropylene (OPP), oriented polyethylene terephthalate (OPET), and PVDC coated substrates. Acrylate copolymers are also used as sealant layer resins, in flexible hoses and tubing, as well as carrier
resins for making masterbatchs and compounds that allow for an increased loading of
fillers, pigments, and additives. Acrylate copolymers, especially EEA and EBA, are also
used in wire and cable applications such as semiconductive layers and insulation layers. They have also found applications in hot melt adhesive formulations and asphalt
modification. Suppliers of acrylate copolymers include Arkema, DuPont, ExxonMobil
Chemical, The Dow Chemical Company, Westlake Chemical, and Mitsubishi.
4.5 Acid Copolymers
Copolymers of ethylene and acrylic acid (EAA) or methacrylic acid (EMAA) are commonly referred to as “acid copolymers.” Acid copolymers are made using a high pressure free-radical polymerization process using autoclave reactors. They have a high level
of long-chain branching and therefore exhibit high melt strength. EAA and EMAA
are nearly equivalent in performance at the same mole percent acid comonomer level.
For equivalent polar functionality (mole % comonomer), EMAA copolymers have a
slightly higher comonomer (acid) weight % than EAA copolymers. Major suppliers of
EAA include The Dow Chemical Company (PRIMACOR ), DuPont (Nucrel ), and
ExxonMobil chemical (Escor ).4 Major suppliers of EMAA include DuPont (Nucrel).
The acid content generally ranges from 3% to 20 wt% and melt index ranges from 1 to
1300 dg/min (190 °C, 2.16 kg). Increasing the acid level increases adhesion to polar
substrates and decreases adhesion to nonpolar substrates. Increasing acid level also
improves resistance to grease and oil. Acid functionality enables excellent adhesion
to aluminum foil, metalized substrates, and paper, and hence acid copolymers are
widely used as tie layers in many flexible packaging applications. These applications
include laminated tubes, aluminum foil packaging, and metalized aluminum packaging. Laminated tubes are used in oral care (e.g., toothpaste tube), food, and pharmaceutical applications. Packaging applications utilizing acid copolymers include condiment
pouches and sachets, aseptic packaging, and processed meat and cheese packaging.
4
PRIMACOR is a trademark of The Dow Chemical Company (“Dow”) or an affiliated company of Dow;
Nucrel is a trademark of E.I. du Pont de Nemours and Company; Escor is a trademark of ExxonMobil
Corporation.
Types and Basics of Polyethylene 117
Polar Ethylene Co-polymers
Ethylene vinyl acetate (EVA)
H2C
– 0.4 to 2500 MI
– 3-20% AA, 4-12% MAA
– Tie-layers (adhesion to metal)
– Partially neutralized to make
Ionomers
CH3
CH
CH2 CH2
O
H
H
H
H
H H
C
C
C
C
C C
H
H
H
C O
x
y
z
H C O
O H
O– Na+
CH2 CH2
CH2
– EEA (Ethylene Ethyl Acrylate)
– EnBA (Ethylene n-Butyl Acrylate)
H
CH2 C
O
H3C
Acrylate copolymers
– EMA (Ethylene Methyl Acrylate)
n
C
O
– High clarity film, sealants & hot
melts adhesives
Ethylene acrylic acid (EAA),
ethylene methacrylic acid
(EMAA) & Ionomers
EVA
Vinyl acetate (VA)
– 5-40% wt VA
C
m
O
x = ethylene
y = acrylic acid
z = sodium neutralized acrylic acid
x, y, z are randomly distributed
H
C
C
RO
O
– Wire & cable, tie-layers
Figure 4.9 Chemistry and key attributes of polar ethylene copolymers made using high pressure process.
Non-packaging applications of acid copolymers include protective metal coating for
corrosion protection, powder coating (e.g., football face masks, playground equipment,
and outdoor furniture), and wire and cable. Ethylene methacrylic acid is also used as an
intermediate for producing ionomers.
Ethylene acrylic acid copolymers having high acid levels (~20 wt%) are also converted into water-based dispersions. Primary applications of the dispersions include
priming (coating weight of 0.5 to 1.2 g/m2) and lamination, paper and textile sizing,
paper and paperboard coating, metal coating, and heat seal coating.
Chemistry and key attributes of PE polar copolymers are summarized in Figure 4.9.
Note that all polar copolymers of ethylene are made using the high pressure free-radical polymerization process. Such polar copolymers cannot be made using Ziegler-Natta
or current single-site catalysts (coordination catalysis) because the polar comonomers
are catalyst poisons. Catalysts that allow the copolymerization of ethylene with polar
comonomers using coordination catalysis are still of academic and commercial interest.
4.6 Ionomers
Ionomers are copolymers of ethylene and methacrylic or acrylic acid that are ionically
cross-linked with zinc, sodium, or lithium cations [15, 16]. Ionomers are prepared from
acid copolymer via the partial neutralization of the acid groups using a base such as
sodium hydroxide (to form sodium ionomers) or zinc oxide (to form zinc ionomers).
Since ionomers are partially neutralized acid copolymers, they have some acid functionality enabling adhesion to foil, polyamide, and paper. Ionic functionality/crosslinking imparts unique properties such as improved optics, toughness, tensile strength,
abrasion, and oil/grease resistance. Ionomers also exhibit higher stiffness/hardness and
lower peak crystallization temperature (freeze point) at a given crystallinity, higher melt
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Handbook of Industrial Polyethylene and Technology
strength, and hot-tack strength. DuPont is the major supplier of ionomer resins under
the trade name of Surlyn.5 Surlyn ionomers are derived from ethylene methacrylic acid
copolymers, having about 4 to 12 wt% acid level. The ionic nature of ionomers makes
them very sensitive to high humidity environments. They must be packaged in foillined bags or boxes to prevent moisture pick-up.
Ionomers are widely used as extrusion coatings or co-extruded sealants in food
packaging with an emphasis on processed meat, medical device packaging, and skin
packaging. Ionomers used as an inner sealant layer exhibit very good adhesion to processed meat, helping to minimize migration of liquids in the packaging for a more
appetizing appearance. Due to its low adhesion to HDPE, ionomers are used in delamination peel seal applications such as cereal liners. They are also used in durable applications such as golf ball skins, bowling pin covers, athletic footwear, ski boots, perfume
bottle stoppers, and the impact modification of polyamide. Ionomers have excellent
melt strength at thermoforming temperatures, desired for deep-draw thermoforming
applications. Ionomers are also used as interlayers for laminated glass in safety and
security applications (e.g., hurricane-proof glass). Such laminated glass is capable of
stopping projectiles from penetrating glass windows during a hurricane.
4.7 High Density Polyethylene (HDPE)
Polyethylene resins having a density greater than 0.940 g/cm3 are defined as HDPE.
HDPE is the highest volume type of PE used today. The need of pressures in excess
of 2000 bars (30,000 psi) for the manufacturing of LDPE resins requires thick-walled
autoclave and tubular reactors, and large compressors leading to high capital cost, high
maintenance and energy cost. As a result, a large research effort was undertaken to
enable polymerization of ethylene at lower pressures to improve the process economics. HDPE was first synthesized by Prof. Karl Ziegler of the Max Planck institute in
Germany using titanium and zirconium halides with aluminum alkyls at much milder
process conditions (much lower pressure and temperature) compared to LDPE. Prof.
Ziegler shared the Nobel Prize for his discovery with Prof. Giulio Natta, who discovered
that polypropylene can also be produced using the same catalyst. These catalysts, in
general, are referred to as Ziegler-Natta (Z-N) catalysts. Around the same time, Hogan
and Banks at Phillips Petroleum synthesized HDPE using silica/alumina supported
chromium oxide catalyst at relatively low pressures. Phillips Petroleum commercialized
chrome catalyzed HDPE resins using a loop slurry process in late 1950s [17]. Hoechst
commercialized Z-N catalyzed HDPE resins using a continuous stirred tank reactor
(CSTR) slurry process in the late 1950s. Subsequently Union Carbide developed a gas
phase process to make HDPE resins also in the late 1960s.
Completely linear HDPE resins (no comonomer) are quite brittle and prone to environmental stress cracking. To overcome this, small amounts of alpha-olefin comonomer (1-butene or 1-hexene or 1-octene) are incorporated into the PE backbone. This
decreases the density/crystallinity and increases environmental stress crack resistance
(ESCR) due to an increase in tie-chain concentration. Use of a higher alpha-olefin
5
Surlyn is a trademark of E.I. du Pont de Nemours and Company or its affiliates.
Types and Basics of Polyethylene 119
Weight percent
High MW copolymer
fraction for tie chains
for improved ESCR,
impact, burst, slow
crack growth
Low MW homopolymer
fraction for stiffness
and processability
Comonomer
content (wt%)
Mol. Wt.
Figure 4.10 Molecular weight distribution of Z-N catalyzed bimodal HDPE resin [1].
comonomer (1-hexene or 1-octene) leads to a significant improvement in ESCR compared to using 1-butene at the same final density. Subsequent advances led to the development of bi-modal HDPE resins, having improved mechanical properties using Z-N
catalysts in a dual reactor slurry or gas phase process. Such bi-modal HDPE resins have
a very high molecular copolymer fraction (using 1-hexene or 1-butene as a comonomer) and a very low molecular weight homopolymer fraction (little or no comonomer).
The very high molecular weight copolymer fraction provides improved environmental
stress crack (ESCR), impact, and slow crack growth resistance due to a high tie-chain
concentration. The very low molecular weight homopolymer fraction provides stiffness
and processability, as shown in Figure 4.10. Bimodal HDPE resins exhibit superior balance of stiffness/impact toughness/ESCR balance compared to unimodal chromium
oxide catalyzed HDPE resins. Attempts have been undertaken to produce such bimodal
resins in a single reactor, using bimodal catalyst systems (e.g., a dual-site catalyst or
bimetallic catalyst) or catalyst blends [18]. However, product design freedom is limited
compared to the use of multiple reactor systems.
Due to a high degree of crystallinity, HDPE resins exhibit very high modulus (stiffness) and excellent chemical resistance. The first very big application of HDPE was
the “hula hoop” made by a California-based toy manufacturer, the Wham-O company.
Subsequently, blow molded bottles were developed from chromium oxide catalyzed high
molecular weight high density PE (HMW-HDPE). Chromium oxide catalyzed HMWHDPE unimodal resins typically exhibit very high weight average molecular weight (MI
< 0.1 dg/min), a broad molecular weight distribution (MWD), and often a very low level
of long-chain branching. Chromium oxide catalyzed HMW-HDPE unimodal grades
for blow molding exhibit the desired parison swell, top load, and adequate environmental stress crack resistance (ESCR) and toughness. Blow molded bottles made from unimodal and bimodal HDPE resins are used in household and industrial chemical (HIC)
bottle applications such as detergent, bleach, fabric softener, and agricultural chemicals
such as pesticides and herbicides. Such bottles are also used in dairy, water, and juice
packaging applications as well as pharmaceutical, medical, and cosmetics applications.
Both unimodal and bimodal HMW HDPE resins are also widely used to make pipes for
gas and water distribution, blow molded fuel tanks, motor oil containers, in automotive
applications, and also large canisters, tanks, and drums, requiring specific approvals for
the storage and transportation of dangerous filling goods (UN certification). Bimodal
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Handbook of Industrial Polyethylene and Technology
HDPE resins are used to make T-shirt bags used for grocery shopping applications,
and oriented tapes for woven raffia applications. A unique application of HDPE is high
moisture barrier liners for dry food packaging such as cereals, crackers, and cookies.
Such very high moisture barrier grades are broad MWD bimodal HDPE resins with a
density typically greater than 0.960 g/cm3. Other HDPE applications include injection
molded drink cups, crates, pails, small tanks, and containers. HDPE resins are used
to make monocomponent and bi-component fibers (HDPE sheath and polypropylene
or polyethylene terephthalate core) to make spunbond fabric as well as bicomponent
staple fibers for hygiene applications [19]. The Dow Chemical Company has commercialized ASPUN fiber grade resins since 1986 for spunbond fabric and staple fiber
applications.6
4.8 Ultra-High Molecular Weight HDPE (UHMW-HDPE)
Ultra-high molecular weight HDPE (UHMW-HDPE) resins have a molecular weight
greater than about three million Daltons. UHMW-HDPE has many outstanding properties such as high abrasion resistance, impact strength, chemical resistance, environmental stress-crack resistance, and a low coefficient of surface friction. However, such
polymer cannot be processed using conventional melt processing techniques in plastics
fabrication due to its extremely high melt viscosity. UHMW-HDPE is produced in a
commercial gas or slurry phase process as a powder. Suppliers of UHMW-HDPE resin
include Celanese corporation (trade name is GUR ), Mitsui Chemicals (trade name
is Mipelon), etc.7 The powder is directly compression molded, sintered and machined
into the final shape of products for some biomedical applications such as hip, knee, and
spine implants, or applications requiring low friction surfaces such as alpine skis, bearings, and gliding parts. UHMW-HDPE powder is also used as an additive to enhance
the lubricity and abrasion for compounding of rubbers and engineering plastics.
UHMW-HDPE is also converted into very high strength fibers (e.g., Dyneema from
DSM, Spectra from Honeywell, and Tekmilon from Mitsui Chemicals) using a gel
spinning process [20–23].8 A gel consisting of 0.5–10 wt% UHMW-HDPE in decalin or
xylene is processed by an extruder through a spinneret. The extrudate is drawn through
the air and then cooled in a water bath. The resulting fiber has a very high degree of
molecular orientation, and therefore exceptional tensile strength. These fibers are used
in body armor, fishing line, bow strings, and high performance sailing. UHMW-HDPE
is also used to make microporous battery separator membranes using solution processing techniques.
6
ASPUN is a trademark of The Dow Chemical Company (“Dow”) or an affiliated company of Dow.
GUR is a trademark of Celanese Corporation; Mipelon is a trademark of Mitsui Chemicals.
8
Dyneema is a trademark of DSM; Spectra is a trademark of Honeywell International; Tekmilon is a
trademark of Mitsui Chemicals.
7
Types and Basics of Polyethylene 121
4.9 Linear Low Density Polyethylene (LLDPE)
Since Z-N and chromium oxide catalyzed HDPE resins (density > 0.940 g/cm3) exhibit
much higher stiffness and inferior dart impact, tear, and optics compared to LDPE
resins (typical density of ~0.920 g/cm3) they are not suitable for many flexible packaging applications. Hence, there were strong commercial incentives to manufacture Z-N/
Chromium oxide catalyzed PE resins using low pressure processes (low capital cost) at
lower densities (0.915 to 0.940 g/cm3), by incorporating higher levels of alpha-olefin
comonomer. These resins were termed as “linear” low density PE (LLDPE), reflecting
the lack of long-chain branching in the resins. Note that sometimes the density range
of 0.915 to 0.940 g/cm3 is sub-classified into LLDPE (0.915 to 0.930g/cm3) and MDPE
(0.930 to 0.940 g/cm3) resins. However, for the purpose of this chapter, the entire density range of 0.915 to 0.940 g/cm3 is termed as LLDPE resins.
DuPont in Canada was first to commercialize Z-N catalyzed LLDPE resins made in
a solution process in the early 1960s. However, the commercial potential of the Z-N
LLDPE resins was not realized for the next 15 years due to the poor extrusion processability and poor bubble stability of LLDPE resins compared to LDPE resins as a
result of the lack of long-chain branching. Union Carbide adapted their HDPE gas
phase process to make LLDPE resins in the late 1970s. The UNIPOL LLDPE gas phase
process was made available for worldwide licensing in 1977 to make PE resins above
0.915 g/cm3, which accelerated the commercialization of LLDPE resins [24].9 Currently,
UNIPOL gas phase process is licensed by Univation technologies. British Petroleum
announced its own gas-phase process for LLDPE in the early 1980s. This gas phase process is currently licensed by INEOS technologies under the trade name of Innovene
G.10 Initially, gas-phase LLDPE resins were made with 1-butene as the comonomer.
In 1978, The Dow Chemical Company commercialized LLDPE resins based on a
1-octene comonomer using their proprietary solution process, under the trade name
of DOWLEX .11 Octene-based Z-N LLDPE resins exhibited superior film mechanical
properties such as dart impact and tear compared to the 1-butene-based gas-phase Z-N
LLDPE resins. This led gas phase resin suppliers to develop their own LLDPE resins
using 1-hexene as a comonomer. For the chromium oxide catalyst-based slurry process,
the lowest density achievable is about 0.920 g/cm3.
Z-N LLDPE resins exhibit a relatively broad molecular weight distribution (MWD
from about 3.5 to 4.5) and a broad short-chain branching distribution (comonomer
distribution). This is due to the multi-site nature of the Z-N catalyst with differences
in each site’s ability to incorporate alpha-olefin comonomer. Catalyst sites that readily
incorporate alpha-olefin comonomer produce lower molecular weight chains as the
incorporation of alpha-olefin slows down the rate of polymerization. Catalyst sites that
do not readily incorporate alpha-olefin produce higher molecular weight chains. Thus,
Z-N catalyzed LLDPE resins are molecular blends of high molecular weight lightly
short-chain branched molecules, lower molecular weight highly short-chain branched
molecules and everything in between. Hence, LLDPE (and VLDPE/ULDPE) resins are
9
UNIPOL is a trademark of The Dow Chemical Company (“Dow”) or an affiliated company of Dow.
Innovene is a trademark of INEOS Europe Limited.
11
DOWLEX is a trademark of The Dow Chemical Company (“Dow”) or an affiliated company of Dow.
10
122
Handbook of Industrial Polyethylene and Technology
Melt strength of Z-N LLDPE (DOWLEX) and LDPE resins
Goettfert Rheotens Data @ 190 C
14
DOWLEX 2045
LDPE A
12
LDPE A
LDPE B
Force (cN)
10
8
LDPE B
6
Dowlex 2045
4
2
0
0
50
100
150
200
250
Velocity (mm/s)
DOWLEXTM 2045, MI = 1.0, 0.920 g/cc
300
350
LDPE A, MI = 1.0, 0.923 g/cc
LDPE B, MI = 2.0, 0.925 g/cc
Figure 4.11 Rheotens melt strength of Z-N LLDPE and LDPE resins.
classified as heterogeneous PEs. Note that HDPE and LLDPE resins produced with
chromium oxide catalysts exhibit much broader molecular weight distribution (MWD
~10 to 20) compared to Z-N catalyzed HDPE and LLDPE resins. Chromium oxide
catalyzed LLDPE/HDPE resins exhibit very good melt strength during film blowing
due to their broad MWD [25]. These resins are used in thick gauge film applications
such as geomembranes as well as a blend component in film formulations to improve
melt strength/bubble stability.
Due to lack of long-chain branching, Z-N LLDPE resins exhibit poor extruder processability (high torque, amps and back pressure) and melt strength/bubble stability
compared to LDPE resins. Melt strength of Z-N LLDPE and LDPE resins are compared
in Figure 4.11. LDPE exhibits about three times the melt strength compared to a Z-N
LLDPE resin at the same melt index (MI=1 dg/min). A 2 dg/min melt index LDPE resin
exhibits about twice the melt strength of a 1 dg/min melt index Z-N LLDPE resin.
The processing deficiencies of LLDPE resins have been systematically addressed
using lower (L/D) screws, barrier-flighted screws, improved die designs and cooling air-ring designs, and the use of wider die gaps for extrusion and film processes.
Blown film properties of LLDPE resins strongly depend upon fabrication conditions
such as blow up ratio (BUR), die gap, melt temperature, and frost line height [26].
Z-N LLDPE resins exhibit substantially improved toughness (ESCR, tensile strength,
dart impact, puncture and tear) properties compared to high pressure LDPE resins.
This has resulted in the growing use of LLDPE resins in flexible packaging applications since their introduction in the late 1970s, mainly at the expense of LDPE resins.
LLDPE/LDPE blends are also widely used in blown film processes to achieve a balance
of toughness and processability and bubble stability, to allow down-gauging and higher
rates. Many of the LLDPE/LDPE blends show a synergy in melt strength over the entire
composition range [27]. LLDPE and LLDPE/LDPE blends are used in applications such
Types and Basics of Polyethylene 123
as trash bags, heavy duty shipping sacks, food storage bags, food and frozen-food packaging, bag-in-box films, lamination films, stretch wrap, stretch hood, greenhouse films,
mulch films, silage films, breathable diaper back sheets, and garment bags. One of the
key and high volume applications of LLDPE resins is stretch films for wrapping pallets to impart pallet stability. LLDPE blown film does not exhibit cross-direction (CD)
shrinkage and cannot be used as a collation shrink film. LDPE is blended in LLDPE to
achieve the desired CD shrinkage (10–25%) and to obtain the desired “bullseye” in the
package for some of the collation shrink bundling applications. Z-N LLDPE resins are
prone to shark skin melt fracture for resins with fractional melt indices due high shear
stress levels at the die lip wall. Melt fracture and die lip build up can be mitigated with
the addition of fluoropolymer processing aids. Z-N LLDPE resins are widely used to
make double bubble oriented shrink film for retail high clarity shrink wrapping and
bundling applications. Such films exhibit balanced shrinkage in the MD and CD directions. Double bubble oriented shrink film process can be described generically by the
following steps: Extrusion-quenching-reheating-biaxial stretching-cooling. The film is
reheated after quenching to a temperature below the melting point of the LLDPE resin
and subsequently biaxially stretched followed by cooling to lock in the orientation.
LLDPE resins are also used in rotomolding applications for making toys, water tanks,
and kayaks. High melt index LLDPE resins are used to make injection molded articles
such as lids and tubs. The higher toughness and ESCR coupled with a faster molding cycle due to the higher crystallization temperatures have made LLDPE resins very
attractive for injection molding applications, displacing LDPE resins. LLDPE resins are
also used to make pipes for raised temperature applications (PERT) such as underfloor
heating applications.
Suppliers of LLDPE resins include The Dow Chemical Company under the trade
name of DOWLEX and TUFLIN , Chevron Phillips Chemical under the trade
name of MARLEX and MARFLEX , ExxonMobil Chemical, Nova Chemicals under
the trade name of SCLAIR , Westlake Chemicals under the trade name of HIFOR and
HIFOR XTREME, LyondellBasell under the trade name of Petrothene , and Borealis
under the trade name of Borstar .12
4.10 Very Low Density Polyethylene (VLDPE)
Very low density polyethylene (also known as ultra low density PE, ULDPE) resins
are Z-N catalyzed resins having densities in the range of 0.885 to 0.915 g/cm3. A lower
density/crystallinity is achieved by incorporating even higher levels of alpha-olefin
comonomer into the copolymer. A key limitation of Z-N catalysts is the inability to incorporate very high levels of alpha-olefin comonomer (1-butene, 1-hexene, or 1-octene)
to produce resins with a density less than 0.885 g/cm3. VLDPE resins exhibit improved
puncture, dart impact, tear, ESCR, low temperature toughness, heat seal, and hot-tack
12
TUFLIN is a trademark of The Dow Chemical Company (“Dow”) or an affiliated company of Dow;
MARLEX and MARFLEX are trademarks of Chevron Phillips Chemical Company; SCLAIR is a trademark of Nova Chemicals Corporation; HIFOR is a trademark of Westlake Chemical; Petrothene is a trademark of LyondellBasell; Borstar is a trademark of Borealis AG.
124
Handbook of Industrial Polyethylene and Technology
properties compared to EVA and LLDPE resins. Applications of VLDPE resins include
flexible tubing, bag-in-box, sealants, low-temperature packaging (e.g., ice bags), barrier
shrink bags for packaging of primal and subprimal meat cuts, medical packaging, flexible hoses, and as a cling layer in stretch films. Suppliers of Z-N catalyzed VLDPE resins include The Dow Chemical Company (ATTANE and FLEXOMER ),13 Westlake
Chemical (MXSTEN ),14 and Nova Chemicals (SCLAIR ).
4.11 Single-Site Catalyzed Polyethylenes
The most recent commercially significant advancement in PE technology is the development and commercialization of ethylene homopolymers and copolymers produced
using single-site catalyst (SSC) technology (e.g., metallocene). The discovery of methylaluminoxane (MAO) as a cocatalyst to activate dicyclopendadienyl (bisCp)-based
metallocene catalysts for highly improved efficiency by Sinn and Kaminsky in 1980
marked the most significant breakthrough in the field of metallocene catalysis [28].
Three major families of high efficiency single-site catalysts (SSC) have been commercially used for the preparation of PE copolymers. These are bis-cyclopentadienyl
single-site metallocene catalyst (also known as a Kaminsky catalyst), a half sandwich, constrained geometry mono-cyclopentadienyl single-site catalyst (known as a
Constrained Geometry Catalyst, CGC, under the trademark of INSITE technology
by The Dow Chemical Company), and post-metallocene catalysts. “Post-metallocene”
refers to a class of single-site catalysts, which are not metallocenes. The use of these
catalyst technologies has allowed a very rapid development of olefin copolymers with
a wide range of structures and related properties [29–31]. Since the early 1990s, SSC
technology has initiated a major revolution in the polyolefin industry.
A key feature of single-site catalyzed PE is the narrow composition (intermolecular) and molecular weight distribution [32]. Such resins are classified as homogeneous
PE. Single-site catalysts have enabled the commercial production of homogeneous PE
resins over a broad density range of 0.857 to 0.965 g/cm3, thus overcoming the key
limitations of Z-N catalysts. These homogeneous PE resins exhibit a broad range of
morphology (from lamellar morphology at high crystallinity to granular morphology
at low crystallinity) and solid-state properties (from necking and cold drawing at high
crystallinity to uniform drawing and high elastic recovery at low crystallinity) [33]. The
morphology of 0.920 g/cm3 and 0.87 g/cm3 ethylene/octene copolymers made by constrained geometry catalyst (CGC) technology is shown in Figure 4.12.
The narrow composition (comonomer) distribution of single-site catalyzed versus
Z-N catalyzed LLDPE is readily observed in crystallization elution fractionation (CEF)
profiles as shown in Figure 4.13. The narrow molecular weight distribution of singlesite catalyzed versus Z-N catalyzed LLDPE is shown in Figure 4.14.
Homogeneous PE random copolymer resins produced by a single-site catalyst
(e.g., metallocene or constrained geometry catalyst) with a density range of 0.885 to
13
ATTANE, FLEXOMER, and INSITE are trademarks of The Dow Chemical Company (“Dow”) or an
affiliated company of Dow.
14
MXSTEN is a trademark of Westlake Chemical.
Types and Basics of Polyethylene 125
(a)
(b)
Figure 4.12 TEM micrographs of 0.920 g/cm3 (A) and 0.87 g/cm3 (B) ethylene/octene homogeneous PE
made using constrained geometry catalyst (CGC) technology [3].
9
8
AFFINITY PL1880G
DOWLEX 2045G
Mass (dWf/dT)
7
6
5
4
3
2
1
0
20
30
40
50
60
70
80
90
100
110
120
–1
Temperature (°C)
Figure 4.13 Crystallization elution fractionation (CEF) of SSC (AFFINITY PL1880G – 1 dg/min MI,
0.902 g/cm3) vs. Z-N catalyzed LLDPE (DOWLEX 2045 – 1 dg/min MI, 0.920 g/cm3) resins.
0.910 g/cm3 are known as polyolefin plastomers (POP), and copolymers with density
below 0.885 g/cm3 are known as polyolefin elastomers (POE). Several families of SSC
technology-based PE random copolymers have been commercialized since the 1990s.
These include polyolefin elastomers (e.g., ENGAGE from The Dow Chemical Company,
TAFMER from Mitsui Chemicals),15 polyolefin plastomers (e.g., AFFINITY from
The Dow Chemical Company; EXACT from ExxonMobil Chemicals, Queo from
Borealis, KERNEL from Japan Polychem Corporation),16 EPDM (NORDEL IP
from The Dow Chemical Company),17 enhanced PEs such as ELITE and ELITE
AT from The Dow Chemical Company, SURPASS single-site catalyzed resins from
15
TAFMER and Evolue are trademarks of Mitsui Chemicals.
EXACT, Exceed and Enable are trademarks of ExxonMobil Corporation; Queo is a trademark of
Borealis AG; KERNEL is a trademark of Japan Polychem Corporation.
17
NORDEL is a trademark of The Dow Chemical Company (“Dow”) or an affiliated company of Dow.
16
126
Handbook of Industrial Polyethylene and Technology
1.2
Mn
1
Mw
AFFINITYPL 1880G 38,594 86,598
DOWLEX2045G
MWD
2.24
25,981 111,612 4.30
dWf / {dLogM}
0.8
AFFINITY PL 1880G
DOWLEX 2045G
0.6
0.4
0.2
0
2.00
3.00
4.00
5.00
Log (molecular weight)
6.00
7.00
Figure 4.14 Gel Permeation Chromatography (GPC) of SSC POP (AFFINITY PL1880G – 1 dg/min MI,
0.902 g/cm3) vs. Z-N catalyzed LLDPE (DOWLEX 2045G – 1 dg/min MI, 0.920 g/cm3) resins.
Nova Chemicals,18 NEXLENE single-site catalyzed resins from SK Chemicals;19 gas
phase metallocene LLDPE/VLDPE (EXCEED and ENABLE from ExxonMobil
Chemicals, EVOLUE from Mitsui Chemical, HARMOREX from Japan Polychem
Corporation); slurry metallocene LLDPE/VLDPE (mPACT from Chevron Phillips
Chemical Company, Lumicene from Total Petrochemicals)20.
Some of the POP/POE (e.g., AFFINITY, ENGAGE, Queo) and mLLDPE resins
have a small amount of long-chain branching and are referred to as substantially linear
homogeneous PEs. The polymer with a vinyl chain-end (formed due to chain termination via a beta-hydride elimination mechanism) gets incorporated into a growing
polymer chain (in addition to ethylene and a-olefin), leading to the formation of longchain branching. Substantially linear homogeneous PE resins exhibit improved shear
thinning (e.g., as measured using I10/I2 or I21/I2 ratio) and extruder processability due to
the presence of a small amount of long-chain branching [34].
The narrow composition distribution (short-chain branching distribution) of singlesite catalyzed homogeneous PE leads to a lower melting point compared to heterogeneous Z-N catalyzed LLDPE/VLDPE resins at a similar density. This is illustrated
in Figure 4.15. Homogeneous octene-based POP resin (AFFINITY PL1880) exhibit a
sharp/single melting point (with a shoulder at lower temperature) versus a broad melting distribution of Z-N catalyzed octene-based VLDPE resin (ATTANE 4203). Note that
the heat of fusion (area under the melting curve) of both resins is similar due to similar density. Single-site catalyzed POP resins exhibit narrow crystallite sizes (thickness)
18
SURPASS is a trademark of Nova Chemicals.
NEXLENE is a trademark of SK Chemicals.
20
mPACT is a trademark of Chevron Phillips Chemical Company; Lumicene is a trademark of Total
Petrochemicals.
19
Types and Basics of Polyethylene 127
1
AFFINITY PL 1880
0.902 g/cc
Heat flow (watts/gram)
0.8
ATTANE 4203
0.905 g/cc
Tm = 98 °C
Tm = 123 °C
0.6
0.4
0.2
0
30
40
50
60
70 80 90 100 110 120 130 140
Temperature (°C)
Figure 4.15 DSC comparison of ethylene/1-octene POP (AFFINITY PL1880) and Z-N catalyzed
VLDPE (ATTANE 4203) resin at similar density.
1.6
POP, 0.909 g/cc
Z-N LLDPE
0.922 g/cc
Heat flow (Watts/gm)
1.4
POP, 0.902 g/cc
1.2
1
POP, 0.896 g/cc
0.8
0.6
Z-N VLDPE
0.912 g/cc
POE, 0.885 g/cc
POE, 0.870 g/cc
0.4
0.2
–20 –10 0
10 20 30 40 50 60 70 80 90 100 110 120 130 140
Temperature (°C)
Figure 4.16 DSC melting curves of SSC (POP and POE) and Z-N catalyzed LLDPE /VLDPE ethylene/
octene copolymers at various densities.
distribution (due to narrow composition distribution) and hence a lower melting point.
Z-N catalyzed LLDPE/VLDPE resins exhibit a broad crystallite size (thickness) distribution due to the presence of a chain fraction having a higher density (low comonomer
incorporated) resulting in thicker crystallites exhibiting a high melting point. Note that
the melting point of a lamella is related to the thickness of the lamella.
DSC melting curves of various octene-based POP and POE resins are compared
with Z-N catalyzed octene-based LLDPE/VLDPE resins in Figure 4.16 (based on
128
Handbook of Industrial Polyethylene and Technology
140
SSC
SSC PE, LLDPE/VLDPE, LDPE & EVA RESINS
SSC
Regression
Melting peak temp. (°C)
120
LLDPE/VLDPE
100
Heterogeneous
LDPE
9% VA
12% VA
18% VA
EVA
80
Homogeneous
60
40
0.86
0.88
0.9
0.92
Density (g/cc)
28% VA
0.94
0.96
Figure 4.17 Melting point of 1-octene-based SSC, Z-N LLDPE/VLDPE, LDPE, and EVA resins as a
function of density [3].
ASTM D-3417 – cooling and heating rate of 10 °C/min). The melting peak temperature
of POP/POE resins decreases with density whereas the melting peak temperature of
Z-N catalyzed LLDPE/VLDPE resins is insensitive to density (due to broad composition distribution).
A plot of melting points of octene-based SSC resins (POP/POE), octene-based Z-N
LLDPE/VLDPE (except resin at 0.885 g/cm3, which is a butene-based Z-N FLEXOMER
VLDPE made in a gas phase process), LDPE, and EVA resins as a function of density is
shown in Figure 4.17.
All samples were cooled and heated at 10 °C/min using a Perkin-Elmer DSC-7. The
LDPE resins have slightly lower melting peak temperature compared to octene-based
SSC resins (made using CGC technology) at a given density. The density of EVA resins increases as vinyl acetate content increases, even though the degree of crystallinity
and melting peak temperature decreases. This is due to the bulky nature of the acetoxy
group increasing the amorphous phase density, as mentioned earlier. Therefore, the
density of EVA resins cannot be compared with the density of ethylene/alpha-olefin
copolymer resins. A plot of melting points of LDPE, LLDPE/VLDPE and SSC ethylene/octene copolymers, and EVA resins as a function of resin crystallinity is shown
in Figure 4.18. Melting points of EVA resins are very close to that of SSC PE resins at
the same crystallinity. The optical properties of LLDPE resins (both Z-N and singlesite catalyzed) depends upon molecular weight and composition distribution and can
be significantly affected by LDPE addition, especially for single-site catalyzed PE. For
a wide variety of PE blown films, haze shows a complex parabolic relationship with
the logarithm of the recoverable shear strain parameter (melt elasticity) [35]. The large
majority of the contribution to the total haze in blown and cast films is observed to
come from the surface roughness of the films, with the bulk (internal) contribution
being relatively minor [11, 36, 37].
Types and Basics of Polyethylene 129
140
SSC
SSC PE, LLDPE/VLDPE, LDPE & EVA RESINS
0.922 g/cc 0.928 g/cc 0.937 g/cc
0.907 g/cc
Regression
Melting peak temp. (°C)
120
LLDPE/VLDPE
0.914 g/cc
0.885 g/cc
LDPE
4% VA
100
9% VA
12% VA
EVA
18% VA
80
SSC
28% VA
60
40
0
10
20
30
40
50
60
70
80
Wt% crystallinity
Wt% crystallinity = heat of fusion/292 (J/gm) * 100
Figure 4.18 Melting point of octene-based SSC (POP/POE), Z-N LLDPE/VLDPE, LDPE, and EVA
resins as a function of DSC crystallinity [3].
Compression molded plaques
Young's modulus (Psi)
105
POP
10
4
POE
Single-Site
LLDPE/VLDPE
Single-Site
103
Regression
102
0.84
0.86
0.88
0.90
Density (g/cc)
0.92
0.94
Log(Modulus) = -130.0787+270.267*den-134.832*den*den, r
0.96
2=0.99
Figure 4.19 Plot of Young’s modulus of nonpolar PE resins as a function of density (single-site and
Ziegler-Natta catalyzed resins). Filled circles are constrained geometry catalyst-based single-site resins.
Filled triangles are bis-cyclopentadienyl metallocene catalyst-based single-site resins.
Figure 4.19 shows the plot of Young’s modulus of compression molded plaques of
single-site and Ziegler-Natta catalyzed resins as a function of density. It can be seen
that Young’s modulus of nonpolar PE resins is primarily a function of resin density
[26]. Mechanical properties such as impact and tear strength of single-site (and Z-N)
catalyzed resins having 1-octene as a comonomer are far superior compared to the
resins having 1-butene as a comonomer [38, 39]. The mechanical properties of resins
having 1-hexene as a comonomer are closer (but still lower) to resins having 1-octene
as a comonomer [39].
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Handbook of Industrial Polyethylene and Technology
Heat sealing is a key functionality in packaging applications [40, 41]. The lower melting point of homogeneous PE (POP) at the same density is advantageous in key application functionalities such as heat sealing and heat shrinkage. Before the 1990s, EVAs and
ionomers were the primary choice for heat sealant layers due to their low melting point
and low heat-seal-initiation temperatures. However, EVA resins exhibit very low hot-tack
strength due to the presence of high levels of long-chain branching inhibiting molecular
diffusion across the molten interface. They can also have taste and odor issues due to
poor thermal stability during extrusion. Ionomers exhibit higher hot-tack strength compared to EVA resins due to the presence of ionic domains and were predominantly used
where higher hot-tack strength was needed such as in vertical form fill and seal (VFFS)
machines. However, ionomers are more expensive resins than EVA. Introduction of
single-site catalyzed POPs in the 1990s exhibiting lower heat-seal and hot-tack initiation temperatures (due to lower melting points), and lower levels of extractibles to meet
FDA direct food contact regulations, has allowed the packaging industry to extensively
use POPs in the sealant layers of various packaging applications. The lower heat-seal
and hot-tack initiation temperature allows faster packaging line speeds and therefore
improved productivity. Single-site catalyzed POP resins also exhibit significantly higher
hot-tack strength compared to EVA. This is illustrated in Figure 4.20 for AFFINITY
POPs of various melt index and densities. Single-site catalyzed POP resins also exhibit
low heat seal temperatures comparable to EVA resins, as shown in Figure 4.21. As a
result, these POPs are widely used for sealant applications (e.g., in VFFS machines).
These resins are also used in fresh product packaging applications (e.g., fresh cut salad)
due to their high oxygen transmission rate (OTR). High melt index (MI 8 to 20 dg/min)
POPs are blended with LDPE for use in extrusion coated sealants for packaging applications such as sachets. POPs are used for the impact modification of polypropylene for
clear, tough polypropylene-based containers (impact needed at typically > 0 °C). The
rapid growth of multilayer packaging applications in the last 20 years has led to significant growth in the usage of single-site catalyzed POP resins as sealants.
Single-site-catalyzed (e.g., metallocene) LLDPE/VLDPE resins exhibit improved
dart impact and puncture, and improved optics (especially when blended with LDPE),
compared to standard Z-N LLDPE/VLDPE resins. These resins are widely used in flexible packaging applications such as stretch wrap film, heavy duty shipping sacks, food
and frozen-food packaging, lamination films, bag-in-box films, stretch hood, greenhouse film, mulch film, and silage film, as well as hygiene applications such as breathable diaper back sheets.
Another key feature of single-site catalysts is their ability to incorporate a very high
level of alpha-olefin comonomers to make polyolefin elastomers (POE). Such low density POE resins exhibit very low modulus and low Shore A hardness for enhanced flexibility and soft touch, very low glass-transition temperature, and high elastic recovery.
Before the 1990s, ethylene-propylene rubber (EPR) and ethylene-propylene-diene
monomer (EPDM) resins based upon Ziegler vanadium catalysts were predominantly
used in applications requiring high UV and oxidation resistance as well as compatibility
with polypropylene. Thus EPDM materials were used as impact modifiers for polypropylene resins. Since the 1990s, metallocene POE resins with densities from 0.857 to
0.87 g/cm3 have found wide commercial acceptance as impact modifiers for polypropylene to make thermoplastic olefins (TPO) primarily used in automotive applications.
Types and Basics of Polyethylene 131
16
AFFINITY POP
1 MI, 0.908 g/cc
Hot-tack strength (N/inch)
14
AFFINITY POP
1 MI, 0.902 g/cc
12
AFFINITY POP
1.6 MI, 0.896 g/cc
10
EVA
0.4 MI, 12% VA EVA
8
EVA
0.8 MI, 18% VA EVA
6
4
2
0
60
70
80
90
100
110
Temperature (°C)
120
130
140
Figure 4.20 Hot-tack curves of ethylene/octene AFFINITY POP vs. EVA resins. Polyamide/EAA/Sealant
(1/1/1.5 mil) blown co-extruded film structure was used [3].
14
AFFINITY POP
0.908 g/cc
Seal strength (lb/inch)
12
AFFINITY POP
0.902 g/cc
10
AFFINITY POP
0.896 g/cc
8
EVA
0.4 MI, 12% VA EVA
6
EVA
0.8 MI, 18% VA EVA
4
2
0
60
70
80
90
100
Temperature (°C)
110
120
130
Figure 4.21 Heat seal curves of ethylene/octene AFFINITY POP vs. EVA resins. Polyamide/EAA/
Sealant (1/1/1.5 mil) blown co-extruded film structure was used [3].
Both 1-octene – and 1-butene-based POE resins are used in these applications. POEs
made using 1-octene as a comonomer exhibit the lowest glass transition temperature and
highest melting point (e.g., DSC Tg ~ −55 °C for 0.87 g/cm3 ethylene/octene POE) and
best compatibility with polypropylene resins compared to either ethylene/butene copolymers (EB) or EPR and EPDM resins, leading to improved low-temperature toughness
and stiffness balance in TPO resins. Availability of POE resins in pellet form vs. bales
for EPDM and EPR grades is also advantageous for compounding with polypropylene.
132
Handbook of Industrial Polyethylene and Technology
Table 4.1 Features, benefits, and applications of ethylene-based plastomers (POPs) and elastomers (POEs).
Features
Benefits
Applications
Narrow Composition
Distribution (CD) and
Lower Tm at a given density
Excellent optics
Low heat seal initiation
temperature (HSIT)
Excellent hot-tack
strength
Sealants
Packaging
Shrink film
Controlled Long-Chain
Branching (LCB)
(Substantially Linear
homogeneous PE)
Improved processability
Extrusion
Ability to make very low
density ~0.86 g/cc
Soft and Flexible due to
lower modulus
High elastic recovery
Impact modification of
PP (TPO)
Elastic films, Fibers
Soft/Flexible goods
With the rapid growth of TPO in the last 15 years, especially in automotive applications such as bumper fascia and instrument panels, the use of POE resins, especially
those with 1-octene as a comonomer, for the impact modification of polypropylene has
increased significantly. POE resins have also found commercial acceptance in soft and
flexible goods, footwear (e.g., cross-linked foam for midsoles), adhesives, cling layers in
stretch films, and elastic films and fibers. Features, benefits and applications of ethylenebased plastomers (POPs) and elastomers (POEs) are summarized in Table 4.1.
The Dow Chemical Company has developed high melt index single-site catalyzed
resins (melt index from 500 to 1200 dg/min, density less than 0.875 g/cm3), under the
trade name of AFFINITY GA, for hot melt adhesive applications as an alternative to
EVA. These resins exhibit much improved thermal and oxidative stability compared
to EVA. Such metallocene-based HMAs give clean and char-free operation leading to
reduced maintenance cost such as time and money to change filters and nozzles, excellent viscosity stability, improved color, and virtually no smoke/odor; all of these lead
to reduced total system cost. AFFINITY GA-based HMA also exhibits aggressive bond
strength over a wide range of service temperatures versus EVA resins due to the lower
Tg of 1-octene-based AFFINITY GA resins.
Single-site catalysts are also used along with modern manufacturing techniques to
yield very high purity, high consistency EPDM elastomers. In addition to ethylene and
propylene incorporation, EPDM elastomers also incorporate dienes as termonomers,
having a pendant unsaturated olefin group to enable sulfur curing. The most commonly used diene is ethylidene norbornene (ENB). Due to the allylic hydrogens associated with ENB, it reacts readily with sulfur and sulfur donors. Less commonly used
are dicyclopentadiene (DCPD) and vinyl norbornene (VNB). VNB exhibits better cure
response to peroxide than ENB, but sulfur cure response is poorer. Figure 4.22 illustrates commercially used dienes for EPDM. Incorporation of the diene termonomer
Types and Basics of Polyethylene 133
Ethylidene
norbornene
Dicyclopentadiene
Vinyl
norbornene
Figure 4.22 Commercially used polymerizable dienes as termonomer for EPDM elastomer.
enables EPDM to exhibit a high degree of cure reactivity for traditional thermoset elastomer applications. The alpha-olefin incorporating capability of metallocene catalysts is
employed to incorporate from 30 to 60 wt% propylene to create semicrystalline to fully
amorphous elastomers.
4.12 Olefin Block Copolymers (OBC)
In the case of the single-site catalyzed random copolymer of ethylene and alpha-olefin
copolymers, incorporating more comonomer along the polymer backbone reduces density, decreases the melting point, and decreases the crystallization peak temperature. As
a result, the heat resistance decreases and cycle times in injection molding increases for
POE. These deficiencies have limited the use of POEs in applications where heat resistance, high temperature compression set, and faster cycle times are desired.
The most recent advancement in PE technology is by The Dow Chemical Company,
overcoming most of the deficiencies of POEs via the introduction of olefin block copolymers (OBCs), commercialized under the trade name of INFUSE . OBCs are block
copolymers of ethylene and alpha-olefin comonomers arranged into alternating “soft”
and “hard” blocks [4]. Soft blocks contain a high level of alpha-olefin comonomer (e.g.,
1-octene) and have a low density/crystallinity and a low melting point. Hard blocks
contain almost no or a very low level of alpha-olefin comonomer and have a high density and crystallinity, and high melting and crystallization temperatures. The soft blocks
deliver flexibility/softness and the hard blocks deliver improved heat resistance and
compression set at 70 °C and faster cycle time via a higher crystallization temperature. Therefore, OBCs combine flexibility and softness properties of POEs but with an
improved heat resistance, elastic recovery, compression set at 70 °C, abrasion resistance,
and faster cycle/setup times.
In order to shuttle or transfer growing chains between two distinct catalysts having
different comonomer (alpha-olefin) selectivity, OBCs are made using a chain-shuttling
agent (CSA). This is shown in Figure 4.23. OBCs are produced in a continuous solution
polymerization process.
The DSC melting curves and peak melting temperatures of OBCs are compared with
those of POE random copolymers in Figure 4.24. It can be seen that at the same density, OBCs exhibit a much higher melting point compared to POP and POE random
copolymers, resulting in an improved heat resistance. The melting point of OBC resins
is almost independent of density due to the presence of high density blocks which melt
at around 120 °C.
Handbook of Industrial Polyethylene and Technology
134
Figure 4.23 Catalytic block technology used to make OBCs [1].
130
OBC (1MI, 0.877 g/cc)
120
OBC (5 MI, 0.865 g/cc)
0.7
Random (3 MI, 0.875 g/cc)
0.6
Random (0.5MI, 0.863 g/cc)
110
Melting temperature (C)
Heat Flow, W/g
0.8
Density versus peak melting temperature
DSC melting curves
0.9
0.5
0.4
0.3
0.2
90
80
70
60
0.1
0
–60
100
50
–10
50
Temperature, C
90
140
40
0.865
OBCs
Random copolymers
0.875
0.885
Density (g/cc)
0.895
Figure 4.24 DSC melting curves and peak melting temperature of OBCs and POE/POP as a function of
density [1].
The DSC crystallization curves and crystallization peak temperatures of OBCs are
compared with POE random copolymers in Figure 4.25. Again, it can be seen that, at
the same density, OBCs exhibit a much higher crystallization peak temperature compared to POP and POE random copolymers. This results in much faster cycle time for
OBCs than for POEs in injection molding applications and much faster setup time in
profile extrusion applications. The OBCs also exhibit improved elastic recovery (lower
permanent set) compared to POEs, as shown in Figure 4.26. OBCs at 0.865 g/cm3 density exhibit elastic recovery similar to styrenic block copolymers such as Kraton G1657
(from Kraton Polymers). Applications of OBC resins include elastic film and laminate,
Types and Basics of Polyethylene 135
Density versus peak crystallization temperature
110
DSC cooling curves (10C/min)
0
100
Crystallization temperature (C)
–0.2
Heat flow, W/g
–0.4
–0.6
–0.8
–1
–1.2
OBC (1MI, 0.877 g/cc)
OBC (5MI, 0.865 g/cc)
–1.4
90
80
70
60
50
Random copolymers
40
Random (3MI, 0.875 g/cc)
OBCs
Random (0.5 MI, 0.863 g/cc)
–1.6
–50 –30 –10 10 30 50 70
Temperature, C
90 110
30
0.865
0.875
0.885
Density (g/cc)
0.895
Figure 4.25 DSC crystallization curves and peak crystallization temperature of OBCs and POE/POP as a
function of density [1].
300
Random copolymer: 0.870 g/cc, 5 MI
Permenant set (%)
250
SEBS
OBC: 0.865 g/cc, 0.8 MI
200
150
100
50
0
0
100
200
300
400
Applied strain (%)
500
600
Figure 4.26 Permanent set of OBC, ethylene/octene random copolymer (POE) and SEBS resins as a
function of applied elongation.
profile extrusions, soft and flexible goods, soft touch over-moldings, and adhesives
applications. OBCs have expanded the competitive space for polyolefin elastomers
against a range of flexible materials and compete with high-value styrenic block elastomers (SBS, SIS, and SEBS).
In 2013, Dow introduced a new addition to the OBC family. These new products are
marketed as propylene-based OBCs and have the INTUNE Olefin Block Copolymers
trade name.21 INTUNE Olefin Block Copolymers contain segments of both isotactic
PP and PE, bridging the gap between the two largest volume thermoplastic polymers
produced today. Because PP and PE are immiscible, blends of these key thermoplastics
21
INTUNE is a trademark of The Dow Chemical Company (“Dow”) or an affiliated company of Dow.
136
Handbook of Industrial Polyethylene and Technology
Table 4.2 Summary of density, melting point, degree of crystallinity, and year of commercial
introduction of major types of PE resins.
Polyethylene technology
Melting point
(°C)
% Crystallinity
Year
0.915–0.931
106–120
45–56
1939
EVA
0.93–0.95*
40–105
5–40
1955
HDPE
0.94–0.967
125–134
62–80
1955
LLDPE
0.915–0.94
120–125
45–62
1975
VLDPE
0.885–0.915
118–122
23–45
1983
SSC PE
0.857–0.967
40–134
2–80
1991
Type
Density (g/cm3)
LDPE
*Density of EVA is higher due to bulky vinyl acetate comonomer. Hence, EVA density cannot be compared
to nonpolar polyethylenes.
typically have poor physical properties, and engineered multilayered structures with PP
and PE show interlayer adhesion failures due to incompatibility. The use of INTUNE
OBCs expands the range of useful blends and resolves interlayer adhesion issues in
multilayered structures, facilitating unique combinations of stable properties in fabricated articles and engineered multilayered systems.
4.13 Concluding Remarks
Ethylene homopolymers and copolymers have a long and rich history of product,
process, and fabrication innovations that has met the growing market needs over the
last 75 years. Many different types of ethylene-based resins have been developed with
broad performance ranges, to meet the requirements of a variety of applications, and
as a result, PE is the highest volume plastic sold today. Ethylene homopolymers and
copolymers constitute a broad and diverse family of large-volume commodity and differentiated resins, exhibiting a very wide range of properties, covering rigid plastics to
elastomers and nonpolar to polar copolymers. Table 4.2 provides a brief summary of
some of the key product categories and lists density, melting point, degree of crystallinity, and year of commercial introduction. This very wide range of properties is accomplished by molecular design, primarily enabled by catalysts, and has led to high growth
rates in applications and the usage of these polymers, especially in flexible and rigid
packaging. Our industrial world would indeed be much poorer without them.
Acknowledgments
The author would like to thank his colleagues Herbert Bongartz, Shaun Parkinson,
Hrishikesh Munj, Teresa Karjala and Sanjib Biswas for their critical review and suggestions.
Types and Basics of Polyethylene 137
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15. Lantman, C.W., MacKnight, W.J., and Lundberg, R.D., Structural Properties of Ionomers,
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5
Molecular Structural Characterization
of Polyethylene
A. Willem deGroot,* David Gillespie, Rongjuan Cong, Zhe Zhou and Rajesh Paradkar
The Dow Chemical Company, Freeport, Texas, USA
Contents
5.1 Introduction ............................................................................................................................... 140
5.2 Molecular Weight – High Temperature GPC ........................................................................ 140
5.3 Comonomer Distribution Measurement Techniques .......................................................... 147
5.3.1 Temperature Rising Elution Fractionation (TREF) .................................................. 149
5.3.2 Crystallization Analysis Fractionation (CRYSTAF) ................................................. 152
5.3.3 Crystallization Elution Fractionation (CEF) ............................................................. 153
5.3.4 High-Temperature Liquid Chromatography (HT-LC) ............................................ 157
5.3.5 Thermal Gradient Interaction Chromatography (TGIC) ........................................ 159
5.3.6 Statistical Parameters .................................................................................................... 161
5.4 PE Characterization with NMR .............................................................................................. 162
5.5 Polymer Analysis Using Vibrational Spectroscopy............................................................... 173
5.5.1 Basic Theory of Infrared and Raman Spectroscopy.................................................. 174
5.5.2 General Applicability of Infrared and Raman Spectroscopy to Polymers
and Related Materials.................................................................................................... 174
5.5.3 Qualitative Identification Using Infrared and Raman Spectroscopy...................... 175
5.5.3.1 PE Unsaturation and Short-Chain Branching ............................................ 176
5.5.3.2 Ethylene Copolymers ..................................................................................... 176
5.5.3.3 Analysis of Multilayer Films .......................................................................... 177
5.5.3.4 Analysis of Polymer Gels ............................................................................... 179
5.5.4 Quantitative Analysis Using Infrared and Raman Spectroscopy ............................ 179
5.5.5 PE Morphology.............................................................................................................. 180
5.5.5.1 Crystallinity ..................................................................................................... 180
5.5.5.2 Orientation ...................................................................................................... 181
5.5.5.3 In-Situ Monitoring .......................................................................................... 184
5.6 Emerging Techniques ............................................................................................................... 186
Acknowledgments .............................................................................................................................. 187
References ............................................................................................................................................ 187
*Corresponding author: AdeGroot@dow.com
Mark A. Spalding and Ananda M. Chatterjee (eds.) Handbook of Industrial Polyethylene and Technology, (139–216)
© 2018 Scrivener Publishing LLC
139
140
Handbook of Industrial Polyethylene and Technology
Abstract
Although polyethylene (PE) is composed of only carbon and hydrogen atoms, the structure of
the molecules are very complex with an almost infinite number of molecular structures. It is
these complex structures that provide the unique properties of the different classes and grades of
PE. This chapter presents some of the many methods for characterizing the molecular structures
of PE resins.
Keywords: ATREF, CEF, TGIC, NMR, IR, GPC, characterization, branching, crystallinity,
structure
5.1 Introduction
Polyethylene (PE) can be thought of as being a schizophrenic polymer. One of its personalities is reflected by its simplicity. It is composed of only carbon and hydrogen.
While the other personality is reflected by the diversity of ways these simple atoms can
be combined to produce unique mixtures of macromolecules. It is this same complexity along with its semicrystalline nature that make it so difficult to characterize for the
modern polymer chemist. The nonpolar nature of the backbone makes it very challenging to dissolve in common solvents, and it usually requires high temperatures and
extended dissolution times to achieve complete dissolution. This coupled with the fact
that most detectors aren’t particularly sensitive to carbon and hydrogen, combine to
make characterizing these polymers surprisingly difficult. Despite these obstacles tremendous progress has been made over the last couple of decades in developing better
instruments and techniques for measuring the molecular structure of these complex
polymers.
The purpose of this chapter is to provide an update to a polymer technology expert
in the basic methodology available today to measure the molecular structure of PE.
The first section will cover developments in the area of high temperature gel permeation
chromatography (GPC) for measuring the molecular weight and molecular weight
distributions of PE. The next section will cover methods for measuring the comonomer distribution using techniques such as analytical temperature rising fractionation
(ATREF), crystallization elution fractionation (CEF) and thermal gradient elution
chromatography (TGIC). The third section will cover the application of nuclear magnetic resonance (NMR) to measure the comonomer content of PE and the final section
will cover the application of vibrational spectroscopy, both Fourier transform infrared
spectroscopy and Raman spectroscopy. All four of these sections bring the reader up to
date on the most recent and important developments in these techniques. It is not the
purpose of this chapter to train an inexperienced reader in all there is to know about
these methods; however, there are ample references given throughout all the sections
which would allow any reader the ability to learn more on any of these topics.
5.2 Molecular Weight – High Temperature GPC
Molecular weight (MW) and molecular weight distributions (MWD) are typically measured using high temperature gel permeation chromatography (GPC). High temperature
Molecular Structural Characterization of Polyethylene 141
GPC characterization of polyolefins was enabled by Waters Chromatography’s development of a high temperature refractive index detector and chromatographic system
[1, 2] shortly after Moore’s revolutionary discovery [3] on polymer separation [4] using
porous beads of styrene divinyl benzene at the bequest of The Dow Chemical Company.
Because of the immediate needs to understand the structure of the popular low density
PE (LDPE) [5, 6] made in the high pressure process along with cross-linked PE [7],
high temperature GPC helped spark immediate innovation in multi-detector technology, on-line and at-line infrared [8], viscometry [9, 10] and light scattering [11–14]
devices. Because of the breadth of polyolefin molecular weight distributions, immediate challenges for column technology were seen and challenges were put forth for
increased pore volume. Secondary challenges included increases in robustness and
resolution [15, 16], especially with the advent of the Ziegler-Natta formulations and
metallocene chemistries. Moreover, calculation methodologies [17] for PE needed to
be developed. A rather complete literature summary of early GPC work, including that
dedicated to polyolefins, was done by Johnson and Porter [18].
Due to the high polydispersity ranges among polyolefin resins, along with the
potential for varied microstructure, including comonomer and long-chain branching
(LCB) incorporation, high temperature GPC systems and their corresponding analysis
[19, 20] conditions have considerable variation, especially column separation technologies, polymer detection [21, 22], and calculation methodologies [23–25]. Because
of this, it is clear that there is still considerable diversity in the techniques practiced
and their results. For instance, an IUPAC study in 2010 [26] monitoring users with
Waters 150C [27, 28], Waters 2000 [29], Polymer Laboratories [30] 210 and 220 systems showed poor reproducibility for LDPE, particularly towards the higher moments
of the distribution (Mz). Much of the variation [31] can be explained through hardware, signal to noise, and analysis conditions, but surprisingly variations in quantities could very much be due to the numerical algorithms [32], extrapolations [33–37],
and baseline choices by the individual examiners [38–43]. Brusseau [44] and Sokolov
[45] elaborate on controlling root causes of general GPC discrepancies. Two classic
samples often used to compare systems are SRM 1475 [46] linear PE homopolymer and
SRM 1476 LDPE [47]. Moreover, there remains a dispersity in the techniques themselves [48], with some laboratories generating “conventional GPC results” based on a
homopolymer PE-equivalent calibration, some relying on relative polystyrene measurements, and some laboratories equipped with Universal Calibration [49–52] or light
scattering analysis [53–57] for “absolute” results. Even among laboratories doing conventional GPC, there are variations in the method between converting polystyrene to
PE equivalent molecular weight, generally using either Mark-Houwink or “Q-factor”
ratioing methods [58, 59]. Other studies [60] have been done and generally show better
results for linear as opposed to long-chain branched polyolefins. With ultra-high MW
polyolefins [61], GPC separations are certainly challenged [62]. Certainly concentration effects [63] in polyolefins need to be considered. For low molecular weight polyolefins, mass spectrometry (MALDI-TOF and proton NMR – end-group determination)
[64] has also been a comparator to GPC, particularly in regards to the number average MW (Mn) moment of the distribution and more recently high-temperature, highperformance liquid chromatography (HT-HPLC) techniques have been employed [65].
Low molecular weight GPC separations extend towards polyolefin waxes [66].
142
Handbook of Industrial Polyethylene and Technology
Because of the need to compare molecular weight distributions between various
GPC techniques [67], systems, and analysis, much research has been accomplished to
understand how detectors and separations are convoluted by axial dispersion [68, 69]
as well as sensitivity and drift differences, detector volume, and offset differences. Many
variants [70] of GPC band-broadening techniques to correct the moments of the distribution, or the entire distribution, are nicely summarized in a book authored by Yau,
Kirkland, and Bly [71]. A review of column selection for application and background
on competing separation modes such as HDC can be found summarized well in compilations by Wu [72, 73] along with Mori and Barth [74]. Some newer research can be
found in the work [75] by Hunt and Holding. Barth, Boyes, and Jackson give an overview of related techniques, including field flow fractionation, TREF, detector choice,
and their use in an excellent review article [76].
Williams and Ward [77] proposed a method for detector alignment whereby conventional GPC, Universal Calibration, and light scattering results could be reconciled
by correcting for band-broadening differences between the techniques. This strategy
artificially increases (or decreases) the geometric volume offset between the detectors
in order to allow all of the techniques to be viewed with the same amount of bandbroadening. Since this offset (a constant value) is indicative of the absolute column
retention slope, it will generate some bias on very heavily branched polymer systems;
however, it is in these systems that the correction is most necessary. Other techniques in
commercial software [78, 79] rely on convoluting a small amount of broadening (usually in the form of a convolution of an exponential and normal distribution function
“exponentially modified Gaussian – EMG”) into the “narrower” and earlier detectors to
allow a more precise chromatographic division, although this has caused some internal
disagreement in the GPC community on how narrow polydispersity should be for high
polymers made with typical “living” polymerizations [80, 81].
Predicting the MWD is necessary to understand rheology [82–87], including swell
[88], reaction control engineering [89], and end-use properties. Moreover, understanding the synthesis [90–92] and structure-property relationships [93] of branching
architecture of model systems in GPC [94–103] including hyper-branched polymers
through a chain-walking mechanism [104] has been well described and is of significant
importance since LCB strongly effects the processing balance of polyolefins [105, 106]
and has led to coupling with multi-angle light scattering (MALS) [107–109], Universal
Calibration, and triple detector systems [110]. Since comonomer addition in linear
low density PE (LLDPE) and high density PE (HDPE), along with naturally forming
short-chain branch (SCB) incorporation [111] in LDPE, dictates melt properties and
influences strength, GPC systems are now frequently coupled to IR and Fourier transform IR (FTIR) systems for quantification and observational trends of comonomer
dependence with molecular weight. For other complex polymers, often a combination
of techniques is used for interpretation such as nuclear magnetic resonance (NMR)
and GPC [112–113]. Considering, blends, and unique architectures, it is often necessary to understand the full microstructure and the two-dimensional MWD-SCBD
dependence, giving rise to multi-detector as well as [114] coupled cross-fractionation
techniques for complex polypropylene (PP) copolymers [115] as well as TREF-GPC
[116, 117] for the analysis of polyolefins, 2D high-temperature, high-performance liquid chromatography (HPLC) methods [118, 119], and high temperature GPC-coupled
Molecular Structural Characterization of Polyethylene 143
NMR [120]. Certainly MW as determined by GPC and tacticity both [121] play an
important role in ethylene-propylene (EP) copolymers [122–124], including block
copolymers [125], impact PP copolymers [126], and blends [127]. And correlations of
molecular weight properties from GPC to standard melt flow measurements are quite
prevalent [128]. Moreover, elution transition between GPC and HPLC behavior for PE
and PP has been demonstrated using porous graphite [129]. Ethylene-propylene-diene
monomer (EPDM) polymers and their compositional distribution and long-chain
branching distributions have been studied through multiple detection coupled with
GPC [130–133] and its relationship to rheological properties [134]. For two-dimensional and copolymer systems, high temperature coupling of evaporative light scattering detectors [135] for concentration measurement is also common, although care
must be taken to linearize the detector for concentration response (particularly with
respect to a “knee” at low concentrations). Finally, solution viscosity effects of shortchain branches are certainly measurable and have been demonstrated on systems such
as polyoctene [136].
Infrared (IR) spectrum detection is becoming a staple mode of detection for high
temperature GPC of polyolefins since being tried successfully [137] with a Perkin-Elmer
Model 12C spectrophotometer at 3.4 m. A popular detector several decades ago (particularly for TREF but also for GPC) was the Wilkes/Miran/Foxboro IR technology [138]
that coupled an infrared cell with a variable wavelength detection mechanism. Although
this detector did not, in general, replace the incumbent standard differential refractometer
designs contained in high temperature GPC chromatographs, it demonstrated the utility
of high temperature infrared for polyolefin analysis. An inherent advantage of the refractometer, is that dn/dc is relatively constant for PE and incorporated alpha-olefins, including PP, whereas the IR absorbance may vary depending upon the filter wavelengths. The
disadvantage of refractometers are their detection of air, water, and solvent background of
stagnant reference with typical additive (BHT) discoloration over time, although vendors
have facilitated flowing references to help mitigate this discrepancy (such as in Polymer
Laboratories, TOSOH, and Malvern units). Moreover, online IR detection technology
is limited in solvent choice to solvents having a suitably low IR background, including
1,2,4 trichlorobenzene [139], ortho-dichlorobenze [140], and tetrachloroethylene [141].
More recently with Polymer Char’s introduction of the IR4 and IR5 fixed multiple filter
wavelength units, sensitivity has been increased by being able to ratio filter bands directly
with the convenience of a single-channel averaged spectrum with continuous dark current and energy referencing. This technology has increased the sensitivity and popularity
of using IR analysis by reducing the complexity of setup and the necessity of using several
software programs congruently which is required for standard FTIR-coupled systems.
For a more rigorous analysis involving the entire FTIR spectrum, with solvent
removed, liquid chromatography (LC) transform [142, 143] and more recently,
DiscovIR [144, 145], systems deposit material on a rotating plate timed with GPC
retention volume. This is extremely useful in the deformulation of compounds, because
the solvent is removed, allowing a “blank background” and yielding full spectrum
interrogation. Also, detection sensitivity is very high since the plate can be rescanned
after the chromatography has taken place. However, in using this type of apparatus,
quantitation can be more difficult because of thickness variations about the deposition
and potential to spread across the plate as an inverse of viscosity. Users will often use
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Handbook of Industrial Polyethylene and Technology
the best practice of establishing a reference band for absolute ratioing within the data.
Functional polyolefin copolymers [146, 147], including those with significant carbonyl
stretching [148], have been studied at length with GPC-coupled FTIR [149] analysis or
a combined interpretation [150] from GPC and IR.
Coupling of a high temperature flow cell to an FTIR unit gives the advantage of constant path length along with full spectrum (as allowable by solvent choice) analysis and
modern mercury-cadmium-telluride (MCT) detection. DesLauriers et al. [151, 152]
successfully coupled this technique [153] with partial least squares (PLS) models to
improve the sensitivity to quantify the very low levels of comonomer commonly used
for HDPE pipe and blow molding applications. Applications also exist for unusual
blocky structures [154]. An approximation to this is the Polymer Char IR5 [155] device
which uses filter wheel technology to give constant ratioing of several fixed wavelength
bands (whereby both broadband and narrow band filters are available). The advantage
of this device is the simplicity that accompanies the ratioing technique, as each filter is
essentially a fully integratable chromatogram.
Of particular importance is the determination of comonomer level (and type – FTIR)
by compositional GPC. Also GPC and NMR have been successfully coupled [156] at
high temperatures. For heterogeneous catalysts, multi-reactor offerings, and blends
such as bimodal resins, the placement of comonomer can produce copolymers that
are advantaged in stiffness and toughness, and thus the determination of comonomer
distribution is of primary concern. For HDPE pipe materials, a description of “tie molecule” relative concentration can be ascertained [157, 158]. Where the polydispersity is
moderate to low (2–8), often CEF [159], ATREF [160], or TGIC [161] has an inherent
advantage in characterization. For cases of very high comonomer content polymers
(elastomers region) often TGIC, NMR, or compositional GPC are the best means of
determining content (as the polymer eluting in ATREF and CEF may reside in the “solubles” [162] region). For very broad molecular weight distributions of medium to high
density PE resins, the MWD can affect elution order of the aforementioned techniques,
and therefore discrimination of the components is advantaged by IR-coupled GPC
because end-group effects can be properly subtracted (as long as an adequate description of the polymer initiation and termination mechanisms are easily understood). It
should be mentioned here that often the objective is to strive for [163, 164] the creation
of “tie molecules” [165] which are a function of both molecular weight and comonomer
incorporation [166–168], and certainly the effect of comonomer will be related to density [169, 170] and thus of interest also to the melting temperature [171] of the resins.
LDPE represents a specific challenge as the back-biting mechanisms [172] produce a
variety of short-chain branches (and for proper determination of LCB—whether by
triple detection or MALS—any SCB must be quantified and subtracted first). For example, ATR-FTIR, rheology and GPC were combined to probe tubular LDPE [173]. IR is
also shown to be a suitable mechanism for looking at exchange reactions with MW for
PE [174]. Obtaining a proper linear reference, shifted by comonomer amount, or the
shifting of a Mark-Houwink plot (triple detection) or the conformation plot (MALS) is
the subject of several articles [175, 176].
For controlled long-chain branching such as that made by Dow Chemical’s
Constrained Geometry Catalysts [177–179] and metallocene catalysts from Exxon
[180], predictive models [181] can be made from the MWD. These models demonstrate
Molecular Structural Characterization of Polyethylene 145
that the bulk of the long-chain branches exist on the higher molecular weight molecules, and multi-detector [182] GPC has been used extensively to study them. As such,
for the determination of number-average molecular weight (Mn), often the effects
of LCB can be ignored; however, it may still be necessary to take into account the
effects of SCB for Mn. Typical Ziegler-Natta (Z/N) catalysts, for instance, have heterogeneous distributions of comonomer (larger amounts of comonomer on the lower
molecular weight molecules [183]). The effect of SCB is directly evident whether the
polyolefin resin is analyzed by MALS or by triple detector GPC, but is not detected
by “conventional” GPC, which employs only a concentration detector. A correction
for the comonomer by IR could be made to determine an absolute molecular weight
“linear-equivalent” backbone. GPC is used as a primary determination of molecular
weight across catalyst types, including metallocene [184] and Ziegler-Natta [185–187].
Additionally, dynamic light scattering has been performed [188] on polyolefins in high
temperature GPC.
Multi-angle light scattering remains the most direct means of measuring branching for polyolefins as it employs Zimm-Stockmayer equations directly. However, when
applied to LDPE or very broad molecular weight distributions sometimes associated
with HDPE polymers, there are potential problems: 1) the LDPE molecules may be
elongated in the flow stream, thus distorting the radius of gyration, Rg, measurements,
and 2) the equations are applied to either a completely polydisperse system (nonseparating) or a monodispersed slice (neither of which are completely true for LDPE
volumetric fractions on a GPC column). Local polydispersity [189] has been shown
to be detectable for polyolefins, such that the Mn, Mw, and Mz of a single slice can be
calculated. Since MALS extrapolations (and subsequent Rg measurements) are highly
influenced at each slice by any higher Rg and Mw constituents, readily matching of Rg
versus Mw of the slice can be distorted, as the slices are neither monodispersed or truly
polydisperse, MALS measurements of LCB frequency may be inaccurate [190] across
these highly branched structures. Careful work has to be done to understand the influence of side-branches [191]. For Universal Calibration of polyolefins [192], a value of
“b” must be assumed (the factor which relates hydrodynamic volume to Rg [193]) and
it has been demonstrated that this value [194] may depend on the branched structure
itself. Moreover, comonomer addition in LLDPE has been noted to affect coil dimensions as measured by MALS and triple detection [195]. Proposals for two-angle light
scattering (TALS) [196, 197] have also been made [198] and iterative corrections for
simultaneous light scattering (higher angle) and viscosity [199] have been attempted
that employ relations from Flory and Fox [200] and Ptitsyn and Eizner [201], including
single-angle observations. Nonetheless, for LDPE there have been numerous attempts
at relating end-use properties to the observed GPC-MALS measurements [202], as well
as Universal Calibration and triple detection. When using a multi-detector GPC for
ratioing, decisions have to be made for data handling at low molecular weights [203]
for viscosity and light scattering detectors.
Thus many of the LCB determinations in PE are comparative rather than absolute
(with method descriptions in patents being fairly elaborate so that others can repeat the
work specifically [204] or simply focusing flow behavior with g or g’ ratios [205–209]).
An abbreviated method for the determination of LCB in PE has been suggested by
Yau [210], which reportedly increases the repeatability of the measurements and is
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Handbook of Industrial Polyethylene and Technology
described as the GPC BR method. GPC is also used for determining novel branching
architecture in PP [211].
Modern high-temperature light scattering measurements coupled with GPC were
first accomplished using an LCD/Milton Roy KMX low-angle light scattering device
[212, 213]. Popular light scattering devices for the analysis of polyolefins include the
Wyatt (Santa Barbara, CA) MALS system with up to 18 angles, a two-angle (TALS)
detection system from Precision Detectors (now under Agilent – Santa Clara, CA), and
a high temperature two-angle device from Malvern (Malvern, UK).
The GPC work on ultra-broad and ultra-high molecular weight distributions [214]
has led to further studies in high temperature asymmetrical flow field-flow fractionation (HT-AFFFF) [215–220] and hydrodynamic [221, 222] chromatography, along
with possibilities [223] including molecular topology [224, 225] fractionation.
Because of column requirements for both low shear [226] and broad distribution
analysis, the polyolefins approach to high throughput GPC has been more conservative, using several analytical columns, often of large particle size. But in recent years,
higher throughput applications in catalyst discovery have been well documented [227–
234]. High throughput [235] GPC applications, are often accompanied by robust data
systems that allow practical data mining capabilities (Symyx/Accelrys Renaissance
Library Studio /BIOVIA Discovery Studio ) [236], and up-front design of experiments
(DOE) design (Freeslate – Lab Execution and Analysis software (LEA) combined with
their “core module” approach [237]). Polymer Laboratories PL XT-20 Rapid HTGPC
Polymer Screening System [238] is coupled to a robotic sample preparation station specifically for polyolefins analysis. Often, combinatorial methods [239, 240] are used to
determine the relationship of MWD to application specific properties.
In general, GPC and chromatography have been used to determine PE degradation in melt [241–244] processing [245–251] and reprocessing [252, 253], weathering
[254–258] including UV exposure and compatibility [259], g rays [260], and potential mechanisms for biodegradation [261, 262], along with tolerance thereof [263], as
well as degradation due to chemical exposure [264, 265], and in accelerated exposure
testing [266, 267] and their design-resistance [268]. GPC helps validate the theoretical
mechanisms [269] for degradation. GPC–IR–MW detection of coupling of polyolefins using direct azide-coupling [270] and ene-reactions [271] for PP as well as diene
incorporation [272] for PE have been documented along with GPC analysis [273] of
functionalization mechanisms via reactive extrusion [274]. GPC is also one of the many
tools used to analyze chlorinated PE (CPE) resins [275].
Challenges exist in the preparation of GPC samples, including potential thermal
degradation of samples, handling of inert fillers such as carbon black, and dissolution
itself for ultra-high molecular weight PE. The Waters 150C chromatographs have a
novel air-pressured ram filter, Polymer Laboratories provides an independent filtering station for samples, and Malvern and Polymer Char, for example, have integrated
self-flushing sample filters. Polymer Char offers a station for carbon black removal by
passing samples through a cartridge-based system. Degradation of PP within GPC systems has been subject to much scrutiny, and padding with nitrogen, dissolving within
the autosampler, and handling of warm zones in an autosampler have all been rigorously explored. Microwave dissolution [276] and subsequent GPC analysis of polyolefin additives have also been explored.
Molecular Structural Characterization of Polyethylene 147
Column technology has extended to 30 m analysis for ultra-broad molecular
weights. Batch light scattering measurements and high temperature field flow fractionation [277–279] methods have proven successful for descriptions of ultra-high
molecular weight content which may not be separated appropriately [280–282] by size
exclusion chromatography (SEC) techniques, or push the limits of detectability by SEC
[283], although the extent to which molecules at these high molecular weights can fully
swell is a subject of debate (on whether the molecules are indeed branched or they cannot fully unentangle even in solution), and modern flow channels have significantly
improved resolution (PostNova Analytics – Landsberg am Lech, Germany) although
there are still remaining issues with low molecular weight content [284].
Alternative high-temperature chromatographic solvents [285–287] are being
explored for multi-detector GPC, such as the possibility of replacing 1,2,4-trichlorobenzene (TCB) for the analysis by GPC of LDPE with dibutoxymethane (DBM, butylal), a
halogen free and less toxic solvent.
Preparative GPC fractionation is quite possible with many manufacturers supplying
columns, with the advantage of obtaining fractions available for subsequent characterization [288, 289], for example, by NMR, ATREF, light scattering (LS), or differential
scanning calorimetry (DSC) analysis, and the accompanying high temperature systems
are often home-built [290]. Solvent gradient separations [291–293] have also delivered
narrow molecular weight “cuts” up to several hundred thousand MW and there are
commercial units available [294–296]. Conversely, analysis of preparative TREF fractions and subsequent GPC analysis is popular [297–299] and a recent interesting study
[300] compares preparative and analytical cross-fractionation modes.
Handling of antioxidants at the low-MW end of the distribution has also been a
topic of much concern, generally with precision concerns and attending to detail in
regards to the integration analysis [301] being used to describe the methodology.
Additional work has been done in describing the nature of PE waxes and their subsequent blends [302].
For chemical composition distribution analysis, it is often relevant to plot the distribution by a variety of techniques as a function of molecular weight, and certainly GPC
serves as one of the most-popular molecular weight separation basis for this endeavor
[303–310]. As detection technologies multiply and coupling of systems becomes more
common, GPC systems are evolving [311].
5.3 Comonomer Distribution Measurement Techniques
Polyethylene is a polymer containing repeating units of ethylene, C2H4. This polymer is
sometimes named as homopolymer PE. PE, especially in commercial use, can also be
a class of polyolefins made from ethylene and various comonomers having the general
formula of CnH2n. Common comonomers of strong commercial interest are propylene,
butylene, hexene and octene. One key microstructure parameter that material scientists
use to control PE performance, property, processability, and application, is comonomer
content (CC) and comonomer content distribution (CCD). CCD is often referred to as
the short-chain branching distribution (SCBD). Often SCBD and CCD are used interchangeably. Compared with the PE backbone, the comonomer side chains are rather
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Handbook of Industrial Polyethylene and Technology
short. The side chains found in high pressure LDPE are generated during the manufacturing process, without the addition of comonomers.
Comonomer content is closely related to polymer density (g/cm3), a most important
characteristic used in PE product specification. 13C nuclear magnetic resonance (NMR)
and Fourier transform infrared (FTIR) are powerful spectroscopic techniques to quantify comonomer content, but not the comonomer distribution among the molecules.
The distribution of comonomer content along PE chains is affected by many factors.
For example, during free-radical polymerization, statistics predict that the comonomer content of PE chains is distributed around a certain average value. Catalysts with
multiple sites, such as heterogeneous Ziegler-Natta catalysts, make PE with a broad
CCD. Process parameters significantly affect the CCD, as well. Using multiple reactors
in series [312], varying comonomer feed rates and injection locations into the reactor
and/or reactors [313]; and employing a chain shuttling agent in one reactor [314] or in
two reactors [315], adds extra dimensions for tuning CCD. Therefore, it is very important to precisely and accurately quantify the CCD. Additionally, the CCD result can be
used to adjust and optimize the production conditions, leading to further improved
products.
Differential scanning calorimetry (DSC) is a popular technique widely used in industrial quality control laboratories. The melting temperature of PE is correlated with average comonomer content for random copolymers. DSC measures the heat associated
with chain crystallization or melting. The percentage of crystallinity can be calculated
from the heat of fusion [316]. The thermogram of the DSC provides semiquantitative
information about the CCD. A review of the DSC of semicrystalline polymers can be
found elsewhere [317]. The only way to obtain a quantitative measurement of CCD is
to use fractionation techniques.
In the past half century, various separation techniques have been developed to determine the CCD. They are categorized into crystallization-based and interaction-based
techniques. All of them share similar fundamentals. Side chains of comonomers, or
branches, behave as chain defects, and interrupt the chain regularity. The reduction in
chain regularity lowers the ability of the chains to crystallize from a dilute solution. This
property, crystallizability, is related to the Gibbs energy available in the system, which
is influenced by both enthalpy and entropy. The defects also hinder the adsorption and
the interaction of the chains with certain substrates in dilute solutions for interactionbased techniques.
Crystallization-based techniques include temperature rising elution fractionation
(TREF) [318–322], crystallization analysis fractionation (CRYSTAF) [323, 324] and
crystallization elution fractionation (CEF) [325, 326]. High temperature liquid chromatography (HT-LC) of polyolefins was developed recently by using fully porous
spherical graphite particles and solvent gradients [327–330]; or through thermal gradients in an isocratic solvent, a technique known as high temperature thermal gradient interaction chromatography (HT-TGIC) [331–333]. The use of multiple detectors
in series provides detailed information. The use of mass-sensitive detectors, such as
differential refractive index (DRI), infrared absorbance, or evaporative light scattering
(ELSD) detectors, provides concentration as a function of either temperature or solvent
strength. The addition of a light scattering (LS) detector and a viscometer allows one to
quantify not only CCD, but also yields molecular weight, and intrinsic viscosity across
Molecular Structural Characterization of Polyethylene 149
the CCD. The combination of a concentration detector (DRI, IR, or ELSD), with a light
scattering detector and a viscometer is often referred to as a triple-detector system:
triple-detector TREF, triple-detector HT-TGIC, etc.
Both crystallization-based and interaction-based techniques can be scaled up as preparative techniques to provide fractions of specific interest for further characterization,
such as by NMR or density. Combining CCD techniques with other techniques, such as
GPC, leads to two-dimensional methodologies, capable of providing a comprehensive
microstructure map of the CCD as a function of a second property, such as molecular
weight, or comonomer content map versus molecular weight for polyolefins, in addition to comonomer content and distribution [332–336].
These powerful techniques are routinely used in academia, industrial research, and
quality control laboratories to quantify CCD of PE products. In the following section,
an overview of each technique, the key parameters affecting the result of each analysis,
and the pros and cons of the techniques are discussed.
5.3.1 Temperature Rising Elution Fractionation (TREF)
As previously discussed, temperature rising elution fractionation (TREF) is a technique
for the characterization of the chemical composition distribution (CCD) in polyolefins.
According to the Flory-Huggins equation [337], the melting temperature of a random
copolymer in a dilute solution at thermodynamic equilibrium can be expressed by:
1
Tm
1
Tmo
R
ln(P )
Hu
(5.1)
where Tm and Tmo are the melting temperatures of the polymer in a dilute solution and
pure polymer, respectively. Hu is the heat of fusion for pure polymer, R is a constant
and P is the mole fraction of crystalline units in the random copolymer. In this case, P
is the mole fraction of ethylene units. This equation can be simplified when the mole
fraction of the comonomer content (CC) is low:
Tm
Tmo
R (Tmo )2
CC
Hu
(5.2)
At the thermodynamic equilibrium, comonomer content (CC) can be directly
correlated with the melting temperature of each PE fraction (not including olefinic
block copolymers). These two equations provide the fundamental basis for the TREF
technique.
The term “temperature rising elution fractionation” debuted when the technique was
used to fractionate LLDPE in 1965 by Shirayama et al. [318]. The technique was actually introduced early in 1950 [318]. In 1982, Wild and his coworkers further developed
the TREF technique as an analytical temperature rising elution fractionation (ATREF),
making it an indispensable analytical tool for measuring CCD for LLDPE [320].
The TREF technique comprises a liquid chromatography system using a column
packed with an inert substrate [338]. The fractionation is achieved by controlling the
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Handbook of Industrial Polyethylene and Technology
temperature of the column. Experimentally, it consists of the following steps: sample
dissolution in solvent at a high temperature; introduction of the polymer solution onto
a column; crystallization of the polymer chains in the column by decreasing the column
temperature without intentionally introducing flow through the column; and elution
of polymer fractions to a concentration detector or a series of detectors by increasing the column temperature with a moderate solvent flow. During the cooling step,
PE chains without any comonomer (homopolymer PE) are first to crystallize onto the
inert substrate. Further decreasing of the column temperature leads to the crystallization of the chains with increasing comonomer units. Thus, fractions elute in a reverse
order in which they were deposited onto the substrate during cooling. Chains with
high comonomer content, in general above 8 mol%, and/or the chains of low molecular
weight, remain soluble at the final temperature of the cooling step. This uncrystallized
fraction is referred to as the soluble fraction (SF), or purge fraction. Figure 5.1 shows
a CCD profile, with triple detection, of a LLDPE resin. The first y-axis indicates mass,
the second y-axis represents log (MW) in red and log (IV-intrinsic viscosity) in black.
Various solvents can be used in TREF. Typical solvents are chlorinated solvents
such as 1,2,4-trichlorobenzene (TCB), ortho-dichlorobenzene (ODCB), and tetrachloroethylene (TCE). These solvents are sufficiently IR transparent in the region where
non-aromatic hydrocarbons absorb, enabling the use of infrared detectors to monitor
the concentration of polymer during elution. Other non-chlorinated solvents, such as
xylene [318, 339], and green solvents, such as dibutoxymethane [340], have also been
used. The choice of solvent affects elution temperature; the poorer the solvation power,
the higher the elution temperature will be [341]. This effect may be used to “fine-tune”
the separation when one fraction is of special interest. The choice of solvents also
depends on personal preferences regarding green chemistry, solvent toxicity, and highest elution temperature required, and this affects which detectors are to be used. The
differential refractive index (DRI) detector, a commonly used concentration detector
for molecular weight distribution measurement, may be used in TREF analysis with
care [319, 340].
6
Purge fraction
14
5
12
4
10
3
8
2
6
1
4
0
2
–1
0
20
30
40
50
80
60
70
Temperature (C)
90
100
110
Log Mw/log IV
Mass (dW/dT)
16
–2
120
Increase in octene content
Reduce in density
Figure 5.1 CCD of a LLDPE obtained by triple-detector TREF. The fraction eluting at each elution
temperature is measured for its mass, the molecular weight by online light scattering detector, and
intrinsic viscosity by a viscometer.
Molecular Structural Characterization of Polyethylene 151
The elution temperature of the x-axis in Figure 5.1 can be converted into comonomer
content or CH3/1000 carbons by using a calibration curve according to Equation 5.1
or 5.2. The calibration is constructed by using reference materials with narrow CCDs
and known comonomer content, or by using the fractions obtained from preparative
TREF [342]. The fractions should be as narrow (monodisperse with respect to CCD) as
possible, and must be characterized off-line for comonomer content, generally by 13C
NMR. In general, a linear correlation of mole % comonomer versus temperature, or
log comonomer molar fraction versus 1/temperature (kelvin) are used. According to
Equations 5.1 and 5.2, the elution temperature in TREF is only a function of comonomer content, not the comonomer type when crystallization and subsequent melting are
performed at thermodynamic equilibrium. Otherwise, many factors related to polymer
crystallization will also influence the TREF profile. For example, increasing the cooling
rate leads to an elution temperature shift to low temperature [318] and a broadened
TREF profile [343], along with a reduction in measurement accuracy [318].
Wild and his coworkers reported that a slow cooling rate of 1.5 °C/h was necessary
to allow the TREF experiment to be performed close to equilibrium conditions. For
clarification purposes, the method developed by Wild et al. [320] is herein referred
to as “Wild TREF.” Although, not all the experimental parameters of the Wild TREF
method were explicitly released in the reference, a working method based upon the
Wild paper [320] might include: sample being loaded at 125 °C, cooling at 0.125 °C/
min from 125 °C to 100 °C, followed by 0.025 °C/min from 100 °C to room temperature; elution at 20 °C per hour from room temperature until all of the crystallized
polymer elutes where the elution flow rate is 4 mL/min. The volume of sample loaded
to the column is 20 mL, and the sample concentration 0.5 wt% (w/w). This version of
the Wild TREF method takes more than three days to analyze one sample. In 1990, an
automated analytical TREF (ATREF) was developed with improved ease of operation
and reproducibility. The ATREF was capable of analyzing eight samples in 24 hours by
using five TREF columns, each in a separate oven [319]. It was common for users to
construct their own TREF instruments and devise their own methodology in the early
years of TREF development. Although Wild used Chromosorb-PAW, other chemically
inert materials, such as sand, crushed firebrick, glass beads, glass fibers, silylated silica
gel, stainless steel shot, and Al2O3, have been used to pack TREF columns. Other experimental parameters deviated from user to user. Therefore, care must be observed when
comparing TREF results obtained by different instruments with different experimental
parameters at different locations. Surprisingly, there has been no standard reference
material commercially available with a known CCD and few interlaboratory uniformity studies on TREF have been documented in the literature. Because of the possible
differences in TREF data obtained with different instruments, different solvents, and
different test conditions, the Wild TREF method has served as a reference point for
CCD analysis in the polyolefin industry for many years.
Theoretically, the type of comonomer does not affect the elution temperature of TREF
based on Equations 5.1 and 5.2 when the measurement is performed at or close to thermodynamic equilibrium. In other words, in theory, a universal TREF calibration, suitable for different types of common comonomers can be constructed by a set of reference
materials (random copolymer having a narrow CCD) with a known CH3/1000 carbon
value. There is some data showing that octene and hexene fall onto the same calibration
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Handbook of Industrial Polyethylene and Technology
line while 1-octadecane has a different calibration [344]. Ethylene-propylene copolymers may not fit well with this universal TREF calibration, especially at high propylene
contents, because propylene units have different crystallization properties than ethylene
units [345]. Narrowly modulated temperature steps during cooling to improve accuracy
of CCD analysis were developed and were successfully used in PP [346]. A comprehensive review of TREF for polyolefins can be found in the literature [347].
5.3.2 Crystallization Analysis Fractionation (CRYSTAF)
Crystallization analysis fractionation (CRYSTAF) was developed to decrease CCD
analysis time by Monrabal in 1991 [325, 326]. CRYSTAF relies only upon the cooling step for fractionation and it takes 6 to 8 hours per sample in typical applications.
The CRYSTAF instrument, manufactured by Polymer Char in Spain, has five sample
reactors (vessels) that allow up to 5 samples to be analyzed simultaneously by one
sample in one reactor, thus increasing overall throughput; however, to accommodate
this an IR detector is shared amongst all reactors making the sampling process discontinuous. Each reactor is a stirred stainless vessel equipped with an internal ceramic
filter. Polymer sample is added into the vessel and dissolved at high temperature in a
solvent (usually TCB or ODCB) at a concentration generally ranging from 1 to 4 mg/
mL. A small aliquot of the sample solution is drawn from the vessel through the filter,
and the temperature at the time of sampling is recorded. The concentration of the filtrate (X) is measured by an infrared detector. A cumulative mass profile as a function
of temperature, T, is constructed for each vessel. Each vessel is sampled in sequence
while the oven temperature is decreased. The first derivative (dX/dT) of the cumulative profile gives the CCD profile, wherein typically 40 to 50 data points are collected
per sample. A mathematical function may be used in data processing to smooth the
data, but this can broaden the apparent CCD. Resolution can be improved by using a
narrow temperature cooling range, short temperature intervals, and/or a slow cooling
rate, but the number of data points is ultimately limited by the volume of the reactor.
For example, in pipe grades, and blends of PE, separations can be tailored to provide
resolution enhancement at higher temperatures, or enhanced component separation,
respectively. An IR detector is used as the concentration detector. It is difficult to use
a DRI detector due to flow instability; however, it is possible to add a light scattering
detector to simultaneously provide a molecular weight measurement for each fraction.
Materials with high comonomer content are challenging for CRYSTAF, most likely
due to poorly formed crystals and the subsequent possibility of penetration through
the filter. The highest comonomer content suitable for CRYSTAF without sub-ambient
cooling is around 6 to 7 mol% comonomer when the samples have a moderate molecular weight (>~50,000 Daltons). If PE with a high comonomer content is of interest, it
is necessary to cool the vessel to sub-ambient temperature to facilitate crystal formation. It is impossible to analyze materials with more than 12 mol% comonomer by
CRYSTAF, due to the freezing points of the commonly used solvents, ODCB or TCB. It
may be possible to use some other combination of solvent and detector(s); but presently
it isn’t commonly practiced. Similar methodology to that used in TREF can be used
in CRYSTAF to construct the calibration curve of crystallization temperature versus
comonomer content.
Molecular Structural Characterization of Polyethylene 153
The CCD profiles obtained by CRYSTAF and by TREF are similar and both CCD
profiles provide information regarding how comonomer units distribute along polymer chains. However, a direct comparison of CCD obtained by CRYSTAF to that from
TREF is not practical, as the data is gathered under non-equilibrium thermodynamic
conditions. Peak temperatures in TREF are higher than those from CRYSTAF, and are
roughly offset by the temperature of crystallization of the individual components being
compared.
Several experimental parameters affect CRYSTAF results. Crystallization temperature increases linearly with a decrease in cooling rate in the range of 2 °C/min to
0.003 °C/min [348]. The use of slower cooling improves crystallization kinetics, thus
increasing the accuracy of the analysis. Extremely low cooling rates are impractical in
CRYSTAF for routine analysis due to the length of the experiment. Accuracy is also
constrained by the cooling efficiency of the reactor (vessel), although Polymer Char has
supplied heat exchange fins on the reactors to help equilibrate them closer to the GC
oven temperature under ramping conditions. Typical literature cooling rates opt to balance analysis time and temperature generally from 0.2 to 0.5 °C/min.
As in TREF, CRYSTAF profiles are influenced by the molecular weight (MW) of the
PE chains [349, 350], wherein increasing molecular weight narrows the CRYSTAF profile, as predicted by Stockmayer’s equation [351]. Peak temperature initially increases
with MW, and gradually reaches a plateau when MW reaches a certain critical number,
and this critical molecular weight is reported as a function of comonomer content.
For HDPE, the number average molecular weight (Mn) at which the peak temperature
is independent of MW is about 10,000 Daltons [349]. For PE-co-hexene having 1.27
mol% hexene the Mn is reported to be around 28,000 Daltons [350], which is suitable
for polymers with general commercial applications. The MW effect can be partially
corrected by taking terminal methyl groups into account (1/MW). Aside from MW,
the study of homopolymer PE shows that the LCB level may have a minor effect on the
CRYSTAF profile [349].
The effect of comonomer length on CRYSTAF crystallization temperature varies.
Studies show that CRYSTAF peak temperature was independent of comonomer type
for comonomers with the structure CnH2n where n < 6 [348, 350, 352]. Conversely, at
the same comonomer mole fraction, the peak temperature has been shown to increase
with the number of carbons in the structure CnH2n, where n > 6 [353], which was
speculated to be caused from a reduced possibility of branch inclusion in the crystal
formation [348]. A universal calibration curve is therefore not applicable to the most
commonly used comonomers, mainly butene, hexene and octene. This limitation posts
a great challenge to CRYSTAF, and also to all of the CCD techniques discussed here,
when an unknown blend is analyzed. Using multiple characterization techniques in
combination with CCD is highly recommended.
5.3.3 Crystallization Elution Fractionation (CEF)
Crystallization elution fractionation (CEF) was developed by Monrabal in 2007 with
dynamic cooling in an attempt to improve resolution in TREF and CRYSTAF [325, 326].
Dynamic cooling refers to the addition of a small flow of solvent during the crystallization cycle to allow material to be crystallized on fresh inert surface material, helping to
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Handbook of Industrial Polyethylene and Technology
reduce the chance of co-crystallization. Other than dynamic cooling, the CEF experiment is similar to the TREF experiment: sample loading to the column, fractionation
on the column during cooling, elution and detection during heating with an elution
flow rate. A model blend can be used to explain the similarity and difference between
CEF and TREF. This model blend consists of a material ‘A’ with no comonomer, a material ‘B’ with medium comonomer content, and a material ‘C’ with high comonomer
content. Materials A, B, and C have their own CCD without overlapping, as shown in
Figure 5.2. Without the dynamic cooling process in TREF, the materials A, B, and C
stay at the same location on the column during the sample loading and cooling steps;
upon heating, the materials A, B, and C elute in a reverse order according to comonomer content: C followed by B followed by A. In CEF, the blend solution is loaded at the
beginning of the CEF column. At the end of the dynamic cooling process in CEF, the
material A is located at the beginning of the column, the material B is at the middle
of the column, and the material C is at the end of the column. These three materials
are physically separated. Upon heating in CEF, the material A, B, and C elute in the
same order as TREF, C followed by B followed by A, but the material B would have to
travel through a larger volume than material C to reach the detector, and the material
A would travel through even more volume than the material B to reach the detector.
The higher the flow rate during dynamic cooling, the larger the separation between the
three materials. This model blend clearly shows a much improved separation in CEF
compared with TREF for this case [345].
It is important that all of the fractions, including the soluble, or purge, fraction
remain in the column at the end of the cooling step [326]. In other words, the maximum flow rate (Fc) is calculated according to Equation 5.3.
Vc CR
TF TI
Fc
(5.3)
where Vc is the column void volume, TF and TI are the final temperature and initial
temperature of the cooling step, respectively; and CR is the cooling thermal rate. Recent
Crystallization
Elution
TI
TF
TI
TF
FE
(a)
TREF
TF
TI
TI
TF
FC
(b)
FE
Dynamic crystallization
Crystallization elution fractionation
Figure 5.2 Schematic depicting fractionation by crystallizability; (a) the TREF separation process,
and (b) dynamic crystallization and CEF. TI is the initial crystallization temperature. TF is the final
crystallization temperature. FE is elution flow rate. FC is flow rate during dynamic cooling [326].
Molecular Structural Characterization of Polyethylene 155
commercial CEF instrumentation allows multiple cooling steps, and each step can have
its cooling rate CR(i), dynamic flow rate Fc(i), and cooling temperature range T(i).
This feature makes optimization possible to achieve enhanced resolution for specific
species. Equation 5.4 states that all of the fractions must remain within the column
before the heating step starts.
n
Vc
T (i) Fc(i)/CR(i)
(5.4)
1
where Ti is the crystallization temperature range for cooling step i, and CRi is the cooling thermal rate during the cooling step i. Care is needed when multiple cooling and
multiple heating steps are selected. A very stringent elution temperature calibration is
required if long-term data uniformity is of interest.
It has been proven that CEF has an excellent reproducibility for consecutive analyses
[326]. Compared with ATREF and other commercial TREF instruments with multiple
columns and multiple ovens, CEF is equipped with one column and one oven, and an
autosampler. All the consecutive analyses are performed at the same sample concentration, same dissolution temperature and time, and on the same column, which greatly
contributes to the excellent reproducibility of CEF. CEF has been successfully proven
to provide a more rapid CCD measurement than TREF, and is two orders of magnitude faster than Wild TREF, which requires approximately 100 hours per sample. This
ability is highly welcomed by industrial laboratories and is beneficial to high throughput research activities.
The CEF technique is crystallization based. The crystallization process is controlled
by factors such as nucleation rate, chain movement, and crystal growth rate. Dynamic
cooling in CEF offers the following advantages: 1) a fresh surface for each polymer
chain to crystallize, contributing to an improved repeatability when compared with
CRYSTAF and TREF, 2) improved temperature homogeneity in the cross-section of the
column, facilitating a higher crystallization cooling rate, and 3) forced-liquid flow helps
minimize the difference in the diffusion constants of different chains.
The CEF tehnique is not performed at or close to thermodynamic equilibrium. The
change in cooling rate, while maintaining all other experimental conditions constant,
affects the elution temperature. In general, the experimental parameters affecting TREF
have similar effects on CEF. Furthermore, dynamic cooling adds an extra complexity to
CEF. Since the flow rate during the cooling step of CEF can significantly affect the actual
elution temperatures of polymer fractions as shown in Figure 5.2, it is very important
to have a robust temperature calibration method to calibrate the elution temperature
in CEF. One approach is to use a blend of homopolymer PE and a fully soluble monodispersed standard such as eicosane. The elution thermal rate can then be adjusted
so that eicosane and homopolymer PE elute at a specific reference temperature such
as 30 °C and 101 °C, respectively [354]. Daily temperature calibration also minimizes
the effects of pump flow variations due to environmental changes on the liquid chromatography system. Implementation of the same temperature calibration method is a
necessity for interlaboratory uniformity studies. Temperature calibration allows a better understanding of how the experimental parameters, such as dynamic flow rate and
elution flow rate, affect the measured CCD, particularly with respect to resolution and
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co-crystallization. A calibration curve of comonomer content versus calibrated elution
temperature is constructed similarly as in TREF and CRYSTAF by using a series of narrow CCD reference materials with known comonomer content. For unknown samples,
it is highly recommended to use a calibration obtained at the same experimental conditions as the reference materials with similar MW and comonomer types.
When discussing CCD analysis, co-crystallization refers to the phenomenon when
chains with similar but different microstructures form crystals together, and/or coelute together. Co-crystallization can occur when molecules with a higher degree of
short-chain branching are “trapped” by molecules with less branching that crystallize
first. Co-crystallization can also occur when low molecular weight molecules with
less branching are prevented by entropy considerations from crystallizing with higher
molecular weight molecules of the same degree of branching. Co-crystallization is one
of the key factors limiting the resolution and accuracy of CCD analysis using any of the
crystallization-based techniques. Co-crystallization has further complicated the characterization of multiple-component systems and has made comonomer distribution modeling and de-convolution efforts less fruitful for complex systems. Co-crystallization
originates from the fact that fractionation is performed under non-thermodynamic
equilibrium conditions. A simple way to determine if there is co-crystallization in CCD
methods is as follows: 1) choose two samples with similar but different comonomer
contents, preferably with different molecular weights that can be measured by using an
online light scattering detector, 2) run the individual samples and the blend, 3) calculate the expected CCD for the blend by mathematically co-adding the individual components, and 4) compare the CCD of the blend measured with the expected one. If both
CCDs do not overlay well, it indicates the existence of co-crystallization. The common
observations of co-crystallization are as follows: 1) the expected CCD of the blend does
not overlay well with the experimentally measured CCD, 2) the individual components
in the experimentally measured CCD profile move toward each other, and 3) the peak
to valley between the two components is higher than expected. Repeating the same
cooling and heating steps, such as multiple cycle CEF or TREF [346], seems unable to
significantly reduce the extent of co-crystallization. An increase in the dynamic flow in
CEF by itself does not substantially minimize co-crystallization either. Many commercial PE samples possess broad CCDs, and behave similar to multiple component blends
with overlapping CCDs from individual components. A complete understanding of
co-crystallization and subsequent minimization of the degree of co-crystallization is
highly desired.
Another challenge associated with crystallization-based CCD analyses, and their
related two-dimensional techniques, is the narrow working range of comonomer content. All the conventional techniques are based on the ability of polyolefin chains to
crystallize from a dilute solution upon temperature change. Polyethylene with a comonomer content of up to about 8 mol% is capable of crystallizing from a dilute solution
in a reasonable time scale. This comonomer range can be extended to a certain degree
by changing solvent properties. For example, by adding a poor solvent, such as biphenyl
ether, to a good solvent, such as TCB or ODCB; the solvation ability of the liquid phase
for polyolefins may be reduced. Note that most poor solvents also cause high background absorbance, rendering IR detection unusable. However, it is possible to extend
the comonomer range to approximately 13 mol% in ODCB by using sub-ambient
Molecular Structural Characterization of Polyethylene 157
temperatures. Some elastomeric materials with higher comonomer contents are unable
to crystallize even with modified experimental conditions in a reasonable analysis time.
5.3.4 High-Temperature Liquid Chromatography (HT-LC)
High-temperature liquid chromatography (HT-LC) of polyolefins has been developed
by using a solvent gradient and fully porous graphite particles [327–330]. This technique is sometimes referred to as solvent gradient interaction chromatography (SGIC).
The eluent strength is gradually increased by varying the composition of solvents so
that the polymer chains having strong interaction with the stationary phase elute later
[355]. The separation mechanism is presumably due to the reversible interaction of ethylene units with the graphite [356]. The short side chains behave like structure defects,
preventing ethylene units from forming a close interaction on the graphite surface, thus
reducing van der Waals interactions. Because HT-LC is an adsorption, not crystallization, based technique, HT-LC overcomes the key challenges of the crystallization-based
techniques, e.g., narrow effective range for comonomer content and co-crystallization.
HT-LC has been successfully used to separate both semicrystalline and noncrystalline
PE with octene contents ranging from 0 to 100 mol% (technically 100 mol% octene is
polyoctene) [336, 357], ethylene propylene copolymers from 0 to 100 mol% propylene
[358], PP impact copolymers, and olefin block copolymers [359, 360]. The application
of HT-LC, using solvent gradients to polyolefins, has been recently reviewed by Macko
et al. [361].
The HT-LC experiment includes the following steps: 1) sample dissolution, 2) introduction of the sample solution into the graphite column at a high temperature, and
3) elution of the polymer by changing solvent strength, for example, from 100% weak
solvent to 100% strong solvent; and measuring polymer concentration by a detector or a
series of detectors. Depending on the choice of solvents, all or part of the sample in solution can be anchored onto the stationary phase before elution starts. The order of desorption with the increase in the percentage of strong solvent in the eluent, is correlated to
comonomer content, as shown in Figure 5.3 [327]. HT-LC analysis typically requires
less time than crystallization-based techniques. The analysis time can be rather short,
for example, in 10 minutes, excluding sample dissolution time [362], which also makes
HT-LC an excellent candidate for the first dimension of a two-dimensional separation.
The HT-LC separation is based on a delicate balance of the interaction between
the molecularly flat plane of graphite and the ethylene units in PE chains [357], the
interactions between mobile phase and the ethylene units in PE chains, and the interaction between mobile phase and stationary phase. The utmost resolution can be tailored by choosing a solvent pair for the fraction of interest; for example, ethylene glycol
monobutyl ether (EGMBE)/TCB [357], decanol/TCB [363]. This freedom is especially
important for industrial laboratories where identifying chemical composition of blends
or materials with complex CCD are routine tasks. In addition, this delicate balance of
interactions of HT-LC is reflected by the effect of comonomer types [362]. In other
words, there is no universal calibration for HT-LC, even for comonomer CnH2n where
n is equal or larger than six [362].
HT-LC for polyolefins has its own challenges. First is the limitation in the choice
of detectors. Generally, polyolefins only dissolve in a few solvents and only at elevated
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2000
21.7% Octene
24.3% Octene
19% Octene
1750
1500
27.7% Octene
14.8% Octene
mVolts
1250
1000
15.9% Octene
100% Octane
8.5% Octene
Gradient
750
0% Octene
500
2.6% Octene
250
0
9
10 11 12 13 14 15 16 17 18 19
Figure 5.3 Dependence of elution volume on the average octene mol% for ethylene-octene copolymers
by HT-LC [327].
temperatures. Changes in solvent composition during the solvent gradient elution limit
the usefulness of typical concentration detectors, such as differential refractive index
and infrared, and limit the utility of other commercial detectors, such as a viscometer
and light scattering detectors, for online intrinsic viscosity and molecular weight determination. One generally applicable detector is an evaporative light scattering detector
(ELSD). An ELSD detects any non-volatile component in a sample after nebulization
and removal of the solvent or mixture of solvents [364]. However, ELS detection is
known to have a nonlinear response with respect to concentration, and nebulization
efficiency is influenced by polymer molecular weight, polymer composition, and solvent composition [365, 366]. This means that ELSD must be calibrated very carefully in
order to yield reliable results [355], which can be very time-consuming and applicable
only to a narrow band of running conditions. To enable use of an infrared detector,
both the strong and weak solvents have to be relatively transparent to the polyolefin IR
absorbance wavelengths. The pairing of diphenyl ether and TCB is useable with the IR
detector to some extent; however, detector signal to noise is much less than with TCB
alone.
To expand solvent gradient problems in polymer detection, a two-dimensional
HT-LC technique (2D HT-LC) couples HT-LC with a high efficiency SEC column, or
columns, in the second dimension to separate the polymer from the weak solvent peak
to allow a better quantification [332, 336]. By using a mass – and composition-sensitive
IR-5 detector followed by a light scattering detector, abundant information about the
microstructure of PE can be quantified [358]. Taking ethylene propylene copolymer
as an example, 2D HT-LC yields a quantitative CCD profile, along with molecular
weight and comonomer composition at each comonomer content, as shown in Figure
5.4. In 2D HT-LC, more than 100 fractions from HT-LC dimension are sent for the
second dimension analysis (impractical for CEF or ATREF as the first dimension).
However, the polymer concentration in each fraction is orders of magnitude lower than
the concentration of the weak solvent. Using a higher sample concentration or larger
sample loops cannot always address this low signal to noise issue appropriately because
Molecular Structural Characterization of Polyethylene 159
IR
wt% P
Mw, g/mol
60000
1.0×106
90
8.0×105
40000
80
30000
20000
70
Wt% P
IR response
50000
100
6.0×105
4.0×105
Mw, g/mol
Ethylene-propylene copolymer
10000
60
0
2.0×105
50
–10000
0
1
3
3
4
SGIC elution volume, mL
5
Figure 5.4 2D HT-LC of ethylene propylene copolymer by using composition-sensitive IR-5 detector
and online light scattering detector [371].
injections may become irreproducible or poorly shaped peaks may result due to concentration effects. A potential solution is performing “stacked injections,” in which
multiple injections of the polymer are made before starting the desorption step, which
can significantly increase detector response and increase signal to noise [367].
The separation of polymer in SGIC is, however, not only affected by chemical composition [327, 330, 336, 357, 362] but also by molecular weight [355] and architecture
[368]. The molecular weight effect is partially because the porous graphite column
itself is an effective SEC column [369], and also because the strength of the interaction
between the polymer and stationary phase is molecular weight dependent. The ability
to measure composition and MW simultaneously after separation would be beneficial
for understanding the impact of different factors and deriving true CCD. Therefore,
2D HT-LC, using both the concentration and composition channel of the IR5 detector in a TCB and decane system, provides an “absolute” measurement of CCD and its
breadth [358]. This is a very powerful technique, especially for PP, propylene impact
copolymers, and olefinic block copolymers.
5.3.5 Thermal Gradient Interaction Chromatography (TGIC)
Thermal gradient interaction chromatography was first reported as a technique for fractionating non-olefinic polymers in 1996 using a C18 bonded silica column [370, 371].
Recently, high-temperature thermal gradient interaction chromatography (HT-TGIC)
was developed using the same porous graphitic carbon as in SGIC, in an isocratic solvent with good solvation properties and compatibility with IR detection. Rather than a
solvent gradient, HT-TGIC relies upon a temperature change to fractionate polyolefins
[331–333]. Figure 5.5 shows HT-TGIC chromatograms of a series of ethylene octene
copolymers with octene contents of 0 to 50 mol%.
The HT-TGIC technique overcomes the key issues encountered in crystallizationbased techniques and in SGIC. HT-TGIC is capable of fractionating a much wider
range of comonomer (0 to 50 mol% octene) than crystallization-based techniques
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0.0
0
20
40
60
80
100
120
Elution temperature of HT-TGIC (°C)
4.0 mol%
8.5 mol%
0.1
13.9 mol%
32.5 mol%
0.2
21.7 mol%
0.3
50.7 mol%
Concentration
Separating ethylene-octene copolymers
based on octene content by HT-TGIC
140
Figure 5.5 HT-TGIC chromatograms of ethylene-octene copolymers using Hypercarb
0.0 mol%
0.4
160
in ODCB [331].
(generally up to ~12 mol% octene). Because of a simpler solvent composition than
SGIC, HT-TGIC enables the use of many commercially available detectors, such as
infrared detectors (IR), light scattering detectors (LS), and viscometer detectors.
Adding a second dimension of separation, such as HT-TGIC followed by GPC, allows
a much simpler and more precise comprehensive microstructure characterization than
2D HTLC techniques for PE [332, 333].
The HT-TGIC analysis consists of the following steps [331]:
1. Introduction of the polymer solution into a column at high temperature.
2. Retention of the polymer chains onto the column by reducing column
temperature. PE chains are retained onto the substrate once a few ethylene units attach on the surface of the graphite. The temperature at which
polymer chains are retained on the stationary phase is significantly above
its normal crystallization temperature in a dilute solution [331, 372].
Homopolymer PE (Mw 115, 000 Daltons) becomes anchored at 140 °C
in HT-TGIC, which is 55 °C higher than its crystallization temperature
from the same solution in the absence of graphite. A similar observation
was demonstrated by using a highly amorphous ethylene-octene copolymer with 37 wt% octene (EO37). Once the polymer chains have been
anchored onto the substrate, they may subsequently crystallize upon further cooling to room temperature or lower.
3. Elution of the polymer by raising the column temperature with moderate solvent flow. The elution temperature in HT-LC with graphite is
about 50 °C higher than TREF. If in crystal form, the anchored polymer
chains may melt first with some of the ethylene units remaining attached
onto the substrate. Elution of the chains happens when the last anchored
ethylene segments detach from the substrate.
The cooling process can be made with dynamic flow as in CEF or without a dynamic
flow as in TREF. Besides the benefits discussed in CEF, dynamic cooling may have other
advantages in HT-TGIC: 1) reducing the possibility of column plugging due to localized
sample concentration buildup by the spreading of the polymer solution along the whole
Molecular Structural Characterization of Polyethylene 161
column, and 2) minimizing the possibility of forming multilayer adsorption effects by
retaining polymer chains at the same location of the column. Similarly, dynamic cooling also causes band broadening in HT-TGIC as it does in CEF. A fully porous graphite
column such as Hypercarb is an effective SEC column [369]. Increasing the length
of the Hypercarb to improve the HT-TGIC separation has an adverse result [373]. The
SEC effect can be minimized by using other nonporous or less porous graphites other
than Hypercarb [332, 233]. When dynamic cooling is used, the actual elution temperature depends on the flow rate used. The higher the dynamic flow rate, the higher the
difference in elution temperature between different fractions will be; i.e., similar to the
effect shown in Figure 5.2. A robust HT-TGIC temperature calibration method was
developed by using a blend of eicosane and HDPE to minimize environmental fluctuation and hardware performance [374]. By calibrating the thermal heating rate during
elution, eicosane and HDPE elute at 30 °C and 150 °C respectively. This calibration is
highly recommended when comparing HT-TGIC results obtained in different laboratories and long-term data integrity is of interest.
HT-TGIC has become a powerful technique to characterize CCD of PE, especially elastomeric materials and polyolefin block copolymers. For LLDPE (having a
comonomer content less than 8 to 9 mol%), currently HT-TGIC has about half the
resolution as that of the crystallization-based techniques [375]; however, optimization of experimental conditions improves the resolution [333, 376]. Other substrates,
such MoS2, WS2, boron nitride, and silica carbide, were used in HT-TGIC for the
first time in 2011 [369] and these substrates showed similar resolution as graphite
[377]. The use of different binary-solvent mixtures as the mobile phase was used to
increase the resolution in HT-TGIC of ethylene octene copolymers [378]. HT-TGIC
does not have the co-crystallization problems experienced by CEF and TREF; however, HT-TGIC has its own set of challenges. When the heating rate and cooling rate
are high, HT-TGIC has a “co-elution” issue. Co-elution is defined as the phenomena
where different chains with similar microstructures co-elute in HT-TGIC or HT-LC.
Co-elution leads to an error in CCD analysis for blended materials, where the peak
temperature of each individual component shifts toward each other, leading to a significant difference between mathematically added data and the experimentally added
data [376]. In general, the co-elution phenomenon is less significant than what is
encountered by co-crystallization. A challenge for researchers is to be able to increase
the resolution by minimizing axial dispersion effects of HT-TGIC such that it is preferred over CEF in non-elastomeric applications due to its inherent advantage of less
co-crystallization.
5.3.6 Statistical Parameters
All of the TREF, CRYSTAF, CEF, HT-LC, and HT-TGIC techniques report CCD profiles. Various statistical averages have been developed to quantify CCD by a single
number, such as weight or number average temperature (Tw and Tn) [349] composition
distribution breadth index (CDBI) [379], solubility distribution breadth index (SDBI)
[380], comonomer distribution constant (CDC) [354], and average block index (ABI)
[381], in the polyolefin industry to facilitate direct quantitative comparisons between
materials.
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5.4 PE Characterization with NMR
Nuclear magnetic resonance (NMR) is a phenomenon which occurs when some
nuclei are immersed in a static magnetic field and exposed to electromagnetic radiation. This phenomenon can be detected and analyzed to provide structural information of molecules. In PE, both 1H and 13C nuclei can have this phenomenon. NMR has
been extensively used to characterize PE through, for example, short-chain branching
(SCB)/comonomer content, sequence distribution/blockiness, regio-errors, chain end/
unsaturation, long-chain branching (LCB), and tacticity. The following sections will
briefly discuss these characterizations. At the end, some of the NMR techniques and
their recent developments will also be discussed.
Short-chain branches can be introduced by free-radical polymerization of ethylene, such as in LDPE, or by isomerization during the homopolymerization of ethylene [382, 383]. SCBs form structural defects during crystallization; thus, they strongly
affect density, crystallization rate, crystallinity, melting temperature, glass transition
temperature, and other bulk mechanical properties [384]. Generally, SCB is measured
with 13C NMR and is reported as SCB/1000 carbons, particularly when comparing to
FTIR or GPC-IR measurements. NMR has the advantage of being an absolute method;
i.e., it does not require standards. It is important to use correct 13C NMR methods
to obtain reliable results, such as using 1) a short enough pulse length for homogeneous excitation, 2) recycle time (acquisition time plus pulse delay) typically longer
than five times the longest 13C spin-lattice relaxation time (T1) in the sample, 3) the
correct ratio of acquisition to delay time to avoid transient nuclear Overhauser effect
(NOE), 4) inverse gated decoupling pulse sequence to avoid steady-state NOE build up,
5) efficient decoupling with good bandwidth, and, finally, 6) no decoupling sidebands
[385]. To decrease data acquisition time, relaxation agents, such as chromium (III) triacetylacetonate (Cr(acac)3), are often used [386]. Because the gyromagnetic ratio of
an unpaired electron is roughly 1000 times larger than that of proton, the carbon T1
relaxation time will be much shorter, being dominated by the unpaired electron. Very
recently 1H NMR was used to measure SCBs in ethylene-co-butene-co-octene, ethyleneco-hexene-co-propene, ethylene-co-hexene-co-octene, ethylene-co-hexene-co-buteneco-propene, and LDPE [387]. The authors used 1D 1H NMR with homo-decoupling
and 2Ds such as heteronuclear single quantum coherence spectroscopy (HSQC) and
heteronuclear multiple quantum coherence spectroscopy (HMQC) with a 700 MHz
NMR for SCBs quantification. It was reported that NMR acquisition time was 2 min
(1D) to 30 min (2D). It is worth noting that SCBs accuracy measured from overlapped
1D proton peaks and from the 2D NMRs should be poorer than that from 13C NMR, as
in 13C NMR, several peaks corresponding to a branch can be used for the calculation.
Another closely related concept to SCB is the comonomer content; it is often used
when SCBs are introduced by copolymerization of ethylene with short alkenes. 13C
NMR has also been the most successful analytical tool to characterize this quantity
[388]. Figure 5.6 shows a 13C NMR of a PE-propylene sample acquired with a 600 MHz
NMR and there are many peaks.
Correct and unambiguous peak assignment is the starting point of measuring comonomer content. There are a number of ways to assign carbon peaks of PEs such as
Grant-Paul chemical shift rules [389], spectra of model compounds, attached proton
Molecular Structural Characterization of Polyethylene 163
47 46 45 44 43 42 41 40 39 38 37 36 35 34 33 32 31 30 29 28 27 26 25 24 23 22 21 20
f1 (ppm)
Figure 5.6 13C NMR of a PE-propylene sample acquired with a 600 MHz NMR.
measurements, and multidimensional NMR for single-bond and multibond carbon
proton connection [390]. The last approach coupled with high-field NMR has demonstrated tremendous capability in chemical shift dispersion and has recently become
the method of choice. Peak assignments for PE-propylene have been done and updated
through the years [388, 391]. Detailed peak assignments for PE-butene were reported
by Sahoo et al. with multidimensional NMR such as HMQC and heteronuclear multiple-bond correlation (HMBC) at 750 MHz [392]. PE-hexene peak assignments can
be found in Seger and Maciel’s paper [393]. For PE-octene, peaks were assigned [390]
and the assignments were updated later [394]. Qiu et al. proposed a scheme to assign
only the carbons in the center unit of an odd n-ad to achieve the goal of counting
each carbon once and only once [394]. The comonomer content can be estimated with
simple algebraic methods if peaks corresponding to comonomer are well resolved,
for example, using branch peaks [395]. Another method is to calculate comonomer
content from triad contents obtained through collective peak assignments and algebra
calculations [388]. The advantages of collective peak assignments are that configurational splittings or long-range sequence effects (which often give rise to line broadening or poorly resolved peak splittings) can be neglected. The currently accepted “best”
approach to obtain comonomer content is through the “matrix method,” which is a
linear least squares implementation of Randall’s method [388, 396]. The matrix method
uses a vector equation s = fM where M is an assignment matrix, s is a spectrum row vector, and f is a mole fraction composition vector, wherein the triad distribution is most
often used in this composition vector f. The matrix can be solved directly to get triad
distribution [396]. Finally, the error of comonomer content analysis can be analyzed
without calibration standards with a method developed by Qiu et al. [397].
Metallocene catalysts generally produce PE- -olefin copolymers that are more
uniform in composition compared to the ones made with Ziegler-Natta catalysts.
A detailed assessment of the effectiveness of catalysts in producing a homogeneous
structure distribution hinges on the availability of a good analytical technique for measuring monomer sequence distribution. 13C NMR together with the matrix method can
determine the mole fraction of each triad sequence accurately, such as EEE, EEP, PEE,
PEP, PPP, PPE, EPP, and EPE in PE-propylene [396]. Using these triad mole fractions,
the Koenig “B” value, number average sequence length of A (La), and the number average sequence length of B (Lb) for a copolymer of A and B can be calculated with the
following equations:
164
Handbook of Industrial Polyethylene and Technology
B
([AAB] [BAB] [BBA] [ABA])
2 ([A] [1 A])
(5.5)
La
([BAB] [AAB] [BAA] [AAA])
{0.5 ([AAB] [BAA]) [BAB]}
(5.6)
Lb
([ABA] [BBA] [ABB] [BBB])
{0.5 ([BBA] [ABB]) [ABA]}
(5.7)
A B-value of 1.0 indicates a random copolymer, a B-value of 0 indicates complete
blocks of monomers A and B, and a B-value of 2 indicates an alternating copolymer
[398]. Other methods for describing sequence distribution, such as persistence ratio
and cluster index, have also been published [393].
Regio-errors can affect crystallization and melting temperature. For improvement
of kinetic models, these signals are of tremendous relevance since regio-errors affect
termination and propagation probabilities. Furthermore, quantitative information on
related signals of regio-errors will benefit the reliability of calculated catalyst characteristics, like the reactivity ratios. The probability of having regio-errors in high -olefin
concentration copolymers and block copolymers with -olefin as one of the blocks is
higher. For example, Dow’s INTUNE Olefin Block Copolymers possesses a PE block
and isotactic PP as another block. The vast majority of coordination catalysts for the
stereoregular polymerization of propylene insert sequential monomer units in a 1,2
insertion fashion. The most common regio-error observed in PP which is also one of
the components of impact copolymer (ICP) is an isolated 2, 1-misinsertion, which can
exist in one of the four diastereoisomeric forms depending on the stereochemical relationship between the methyl groups, as shown in Figure 5.7. Regio-errors 13C NMR
D
C
M
G
H
J
K
F
C
L
N
RE1
M
I
D
F G H J
K
L N
I
c
m
f
g
d
h
j
k
l
i
n
RE2
g
m
f
i
G
D
H
h j
J
K
F
f
g
L
N
RE3
H
M
F G
J
K
d
i
n
L N
I
c
h
l
D
C
I
m
k
c
C
M
d
j
k
l
d
n
RE4
m
g
h
c
i
f
j
k
l
n
Figure 5.7 3D structures and Fisher projections of 2,1-insertion regio-errors in iPP [399].
Molecular Structural Characterization of Polyethylene 165
peak assignments have become easier due to a dramatic sensitivity increase with NMR
cryoprobe technology; for example, using 2D INADEQUATE, as shown in Figure 5.8
[399]. 13C NMR studies of regio-errors in PE-octene [400] and PE-hexene [393] were
reported as well.
Polyethylenes can be terminated with a saturated or an unsaturated end group. A saturated end group can be formed if hydrogen is used in the polymer production process
and the details are listed below using a hafnium catalyst and PE-octene as an example.
Hydrogen addition to form a saturated chain end:
Ethylene as the last monomer
Hf-CH2CH2P + H2
Hf-H + CH3CH2P
(5.8)
Octene as the last monomer with 1,2 insertion
Hf-CH2CH(C6H13)P + H2
Hf-H + CH3CH(C6H13)P
(5.9)
Octene as the last monomer with 2,1 insertion
Hf-CH(C6H13)CH2P + H2
Hf-H + (C6H13)CH2CH2P
(5.10)
β-hydride elimination occurs to form an unsaturated chain end. There are
four types of unsaturated chain ends: 1) vinylene, 2) vinyl, 3) vinylidene, and
4) trisubstituted. Vinylene, vinyl and vinylidene are formed through -hydride
elimination. A terminal trisubstituted chain end is formed through isomerization of
vinylidene [401]. The mechanism of forming internal trisubstitution is not very well
understood.
Octene as the last monomer inserted in 1,2 insertion to form a vinylidene end group:
H
K G
L J
f
h
k g
j l
F
H
KG
J
L
F1(ppm)
F
45
50
55
60
65
70
42
(a)
40
38
36
34
F2(ppm)
(b)
42
40
38
36
34
F2(ppm)
42
40
38
36
34
F2(ppm)
(c)
Figure 5.8 Fragments of 2D INADEQUATE: (a) 2,1-insertion regio-errors (RE1) in the iPP made with
catalyst A, (b) 2,1-insertion regio-errors of the iPP made with catalyst A. Only resonances for the regioerror with lower concentration (RE2) are labeled, and (c) 2,1-insertion regio-error (RE3) in the iPP made
with catalyst B. NMR time is 2 to 3 days [399].
166
Handbook of Industrial Polyethylene and Technology
Thermal termination
Hf-CH2CH(C6H13)P
Hf-H + CH2=C(C6H13)P
(5.11)
Chain transfer to ethylene
Hf-CH2CH(C6H13)P + CH2=CH2
Hf-CH2CH3 + CH2=C(C6H13)P
(5.12)
Chain transfer to octene with 1,2 insertion
Hf-CH2CH(C6H13)P + CH2=CH(C6H13)
Hf-CH2CH2C6H13 + CH2=C(C6H13)P
(5.13)
Chain transfer to octene with 2,1 insertion
Hf-CH2CH(C6H13)P + CH(C6H13)=CH2
Hf-CHC6H13CH3 + CH2=C(C6H13)P
(5.14)
Octene as the last monomer inserted in 2,1 insertion to form a vinylene end group:
Thermal termination
Hf-CH(C6H13)CH2P
Hf-H + (C6H13)CH=CHP
(5.15)
Chain transfer to ethylene
Hf-CH(C6H13)CH2P + CH2=CH2
Hf-CH2CH3 + (C6H13)CH=CHP
(5.16)
Chain transfer to octene with 1,2 insertion
Hf-CH(C6H13)CH2P + CH2=CH(C6H13)
Hf-CH2CH2C6H13 + (C6H13)CH=CHP
(5.17)
Chain transfer to octene with 2,1 insertion
Hf-CH(C6H13)CH2P + CH(C6H13)=CH2
Hf-CHC6H13CH3 + (C6H13)CH=CHP
(5.18)
Ethylene as the last monomer to form a vinyl end group:
Thermal termination
Hf-CH2CH2P
Hf-H + CH2=CHP
(5.19)
Chain transfer to ethylene
Hf-CH2CH2P + CH2=CH2
Hf-CH2CH3 + CH2=CHP
(5.20)
Molecular Structural Characterization of Polyethylene 167
Chain transfer to octene with 1,2 insertion
Hf-CH2CH2P + CH2=CH(C6H13)
Hf-CH2CH2C6H13 + CH2=CHP
(5.21)
Hf-CHC6H13CH3 + CH2=CHP
(5.22)
Chain transfer to octene with 2,1 insertion
Hf-CH2CH2P + CH(C6H13)=CH2
Vinyl-terminated polymers are especially interesting for many PE materials. These
double bonds can be incorporated back into the polymer chain during polymer propagation and thus can form long-chain branches (LCBs). Furthermore, catalysts create
their unique unsaturation fingerprints which may facilitate catalyst screening and process control [402]. Unsaturation microstructure also provides important information
about the chain termination mechanism. Thermal degradation and oxidation of PEs
are two key concerns of both resin manufacturers and end users. One of the driving
forces for thermal degradation and oxidation is the unsaturation in polymer chains.
1
H NMR is a better technique to characterize the unsaturation in PE and detailed studies for PE-octene were done by Busico et al. [401] and He et al. [403]. For industry,
researches may choose to assess only the total amount of vinyl, vinylene, vinylidene,
and trisubstituted groups per 1000 carbons. These can be obtained based on the collective 1H NMR peak assignments. Figure 5.9 shows an example of PE-octene [404]. For
common saturated chain ends, 13C NMR has often been used to characterize them. For
example, – CH2CH2CH2CH3 has the chemical shifts of 29.6, 32.0, 22.9 and 14.1 ppm,
respectively, ending with CH3. Assignments for some other saturated structures can be
found in Busico’s work [405].
The advent of Ziegler-Natta catalysts has certainly made a significant impact in the
polyolefin industry, allowing tailoring of both chain length and comonomer insertion.
Extraordinary developments of homogeneous catalysts based on group IV metallocenes with the use of methylaluminoxane as a cocatalyst occurred in the 1980s [406]. By
using these new catalysts, polymers with a controlled molecular structure, such as narrow distribution of molecular weights and of comonomer composition, were formed
[407, 408]. The physical and mechanical properties, melting temperature, and crystallinity of PE are highly dependent on the type and number of branches [409]. LCBs are
formed by the incorporation of vinyl-terminated macromers during polymerization.
As these branches are typically longer than the entanglement molecular weight, their
presence strongly affects the processability of the bulk polymer, especially melt strength.
LCB can be detected and estimated using triple detector GPC (light scattering, intrinsic
viscosity, and refractive index) [410]. Rheology provides an indirect method for characterizing the presence of LCB in PE. It was found that LCBs affect the zero-shear viscosity even at very low LCB concentrations [411, 412]. The third detection method is 13C
NMR, which is good for characterizing Y-PE type LCB, as shown in Figure 5.10, and the
results are often reported as LCB/1000 carbons. Some studies used branch carbon for
detection, while others focused on carbon branch to take advantage of improved sensitivity [388]. Solvent effects on 13C NMR chemical shift of the LCB branching carbon
were also studied and it was found that naphthalene is a very good solvent to separate
the branch signals of LCB and SCB in PE-hexene and PE-octene [413].
COSY
F1 (ppm)
Handbook of Industrial Polyethylene and Technology
F1 (ppm)
168
HSQC
4.8
110
5.0
112
5.2
114
5.4
5.6
5.6
5.4
5.2
4.8 F2 (ppm)
5.0
5.1
5.0
F2 (ppm)
4.9
Vinylidene
Vinylene
5.6
5.5
Tri
5.4
5.3
Vinyl
5.2
5.1
5.0
4.9
4.8
PPM
Figure 5.9 1H NMR of an ethylene-octene copolymer obtained with the double presaturation method.
Unsaturation regions are assigned to vinylene, trisubstituted, vinyl, and vinylidene. Fragments of COSY
and HSQC are also shown [404].
H2
C
X
H2
C
C
H2
H2
C
C
H2
Br
H2
C
HC
C
H2
CH2
H2C
CH2
H2C
CH2
X
Figure 5.10 Y-PE type LCB.
H2
C
H2
C
C
H2
X
Molecular Structural Characterization of Polyethylene 169
n=4
n = 20
n = 16
n = 12
n = 10
n=8
n=6
CH(Bn)
Maching CH(B4)
38.30
38.28
38.26
38.24 [ppm]
Figure 5.11 13C NMR spectra of the methine peak region of model copolymers [416].
It is worth mentioning that 13C NMR also has its limitations; it currently only can
distinguish LCB from C20 and shorter branches, as shown by Figure 5.11 [414]. With
higher magnetic field strength, better solvents and other approaches, researchers might
be able to distinguish LCB from higher than C20 branches, but there will almost certainly still be limitations to reach C89, which is the minimum entanglement chain length
for PE [415].
Tacticity of a polyolefin can have a tremendous impact on its physical and mechanical properties [416]. For example, one of the world’s most produced plastics, isotactic
PP (iPP), is semicrystalline, but atactic PP (aPP) is waxy and amorphous, with limited commercial applications. iPP is often a part of PE products, such as in impact
copolymers (ICP) and INTUNE resins, a product of Dow Chemical. The synthesis of many different types of polymers with well-defined stereochemistry has been
actively pursued in both industrial and academic laboratories [417]. 1D 1H NMR can
be used to distinguish iPP and syndiotactic PP (sPP) [417]. 1D 13C NMR is one of the
most commonly utilized techniques for analyzing polyolefin tacticity because 13C has a
larger chemical shift dispersion [405]. For PP, tacticity can be measured with 13C NMR
signals of methyl groups without the nuclear Overhauser effect (NOE), but researchers can also take advantage of NOE to increase signal-to-noise ratio (S/N), since all of
the methyls used to calculate tacticity have the same degree of proton substitution and
have similar mobility. Recently, tacticity was also measured with band-selective 2D
HSQC, as shown in Figure 5.12, showing potential utility when the sample amount is
limited [418].
One of the most important recent NMR technique developments in the area of PE
characterization is the high temperature 10 mm cryoprobe. This probe has brought a
revolution to the chemical industry for PE characterization [419–422]. The cryoprobe
concept was proposed in the middle of the 1970s, and commercial products were developed in the 1990s. But most of the focus was on 1.7–5 mm probes for biological NMR
studies, and the probe’s temperature could only go to 80 °C. For PE characterization,
especially in industry, researchers require over 100 °C for proper polyolefin dissolution
in organic solvents and they were not limited by sample quantity, and therefore it was
2
3
4
5 6
7
8
9
10
0.9
1.0
1.1
21.6
21.2
20 8
F1
20.4
20.0
13
C (ppm)
19.6
19.2
13
1
H (ppm)
Handbook of Industrial Polyethylene and Technology
F2:
170
Figure 5.12 Pentad stereosequences and bs-HSQC of an aPP [420].
greatly desirable to have a high temperature 10 mm cryoprobe to boost NMR sensitivity
for polyolefins. It would take approximately 7 hours to obtain reliable 13C NMR spectra
with typical samples and about 17 hours for blend analysis with conventional 10 mm
probes using a 400 MHz field. Compared with a conventional 10 mm probe, a high
temperature 10 mm cryoprobe developed by Bruker increases sensitivity over 5.5 times
and now it typically takes about 20 to 30 min to acquire a high quality 13C NMR of PE
with a high temperature 10 mm cryoprobe on a 400 MHz NMR spectrometer [419].
Subsequently, the 13C NMR acquisition time of a PE sample can be improved to only
a few minutes with a 600 MHz high temperature 10 mm cryoprobe. Furthermore, the
sensitivity increase allows fractions in the amount of tens of milligrams from chromatography to easily be studied. Figure 5.13 shows 13C NMR of high temperature LC
fractions of a PE-octene block copolymer [422]. The measurement of unsaturation in
PE can also reach a few units per 1,000,000 carbons in less than 30 min with 1H NMR
and a high temperature 10 mm cryoprobe on a 400 MHz spectrometer [404]. High
quality 2D INADEQUATE, which is one of the least sensitive NMR experiments, is now
accessible within reasonable time periods [399]. Currently, the highest field cryoprobe
is a 950 MHz TCI 5 mm probe with an upper temperature limit of 80 °C. It provides
excellent resolution for some elastomer samples although its sensitivity has not been
shown to be better over a commercially available 10 mm cryoprobe in conjunction with
a 700 MHz field.
Dynamic nuclear polarization (DNP) was predicted and verified in 1953 [423], but
was not widely used until recently because of the lack of high-frequency and highpower microwave sources. During a DNP process, stable paramagnetic polarizing
agents are added to an NMR sample and the sample is irradiated with microwaves to
transfer the high polarization in the electron spin to the nuclei of interest to increase
NMR sensitivity. Theoretical sensitivity increases can be over 600 times for 1H and over
Molecular Structural Characterization of Polyethylene 171
Long ethylene
sequence
Whole polymer
HT-LC fraction 1
fraction 3
fraction 5
fraction 6
30.3
30.2
30.1
30.0
29.9
29.8 PPM
Figure 5.13 13C NMR of a PE-octene block copolymer and its HT-LC fractions [424].
2,600 times for 13C [424]. The mechanisms of the sensitivity increase for liquids and
solids are unique. The sensitivity increase can potentially be much higher if polarization takes place at low temperatures with the NMR performed at a higher temperature.
DNP was attempted directly on PE films and showed a high sensitivity increase [425].
The first solution-state DNP NMR (dissolution DNP) related to polyolefins was published recently by Hilty et al. The authors loaded catalyst in an NMR tube and put it
into the main magnet (400 MHz), polarized 1-hexene in a separated magnet at 1.4K for
4 hours, used hot toluene to dissolve the polarized 1-hexene, used He and Ar gases to
transfer the polarized 1-hexene in toluene to the NMR tube in the main magnet; the
polymerization was started and 13C NMR was acquired at 25 °C. Signal gain of 3 orders
of magnitude was achieved [426]. Although there are still many challenges in DNP, it
is certain that there will be some important applications for PE characterization in the
future.
Triple detection GPC with refractive index, light scattering, and viscometer detectors can provide polymer branching or shape in addition to molar mass distribution,
but provides limited information on the chemical composition or the microstructure
as a function of molar mass. With the addition of NMR detection, GPC-NMR, molecules (on the GPC elution volume axis) are further resolved with respect to chemical
differences by NMR (chemical shift axis). Therefore a two-dimensional separation is
achieved; although only one chromatographic method is applied. The power of NMR
detection allows for direct calculation of molar mass of each fraction without calibration
if end groups can be well detected, raising analysis possibilities in applications where
GPC separations fail or in polydisperse branching fractions. With low temperature
GPC-NMR, Ute et al. characterized ethylene-propylene-(2-ethylidene-5-norbornene)
terpolymers; showing that ethylene and propylene content could be sufficiently quantified from the on-flow experiments. The small amount of 2-ethylidene-5-norbornene
did require a stop-flow procedure in the analysis [427]. Analysis of PE by GPC-NMR
must be conducted at temperatures above 100 °C for most samples due to the fact that
they are only soluble at such high temperatures, and the GPC requirement that the molecules be fully unentangled. Therefore, not only the chromatographic fractionation but
also the NMR detection must be carried out at these temperatures. The only high temperature GPC-1H NMR work so far was reported by Hiller et al. Their NMR flow probe
172
Handbook of Industrial Polyethylene and Technology
Probe capillary
Heater circuit
Probe cell
heater
Connection block
heater
Gas flow
Flow
Capillary heater
circuit
Figure 5.14 High temperature flow probe. The scheme indicates the two separated heating circuits and
the connection to the heated transfer lines [430].
could operate at temperatures up to 130 °C. The probe was a dual inverse probe with
pulse field gradients and had a flow cell with a volume of 120 l. An interface between
SEC and NMR which contained a temperature-controlled stop-flow valve was developed. This valve was a two-position device and guided the flow either from the GPC to
the NMR or directly to the waste, as shown in Figure 5.14 [428]. This was excellent pioneering work, but the system reliability was at issue due to probe leaking. Clearly this
issue needs to be resolved before practical use can be achieved for this methodology.
With the recent Bruker development of a high temperature 10 mm cryoprobe [419],
hyphenated-13C NMR might also be possible in the future. With recent developments of
PE separation techniques, other hyphenated-NMR such as TGIC [420] – NMR might
also be possible in the future.
Zhou et al. pioneered thermal gradient NMR (TG-NMR) in which graphite was used
directly in the NMR tube as a stationary phase to study the separation mechanism of
TGIC, as shown in Figure 5.15 [404, 429]. This concept has the potential to be extended
as a simple hyphenated-NMR technique in the future.
Solid-state NMR is also one of the powerful tools for elucidating structures and
dynamics of polymers [430]. Some applications in PE can be found in the literature
[431, 432]. A recent paper reported that the amount of rigid crystalline chains in alltrans conformations, amorphous chains with increased equilibrium gauche conformer
content undergoing essentially isotropic reorientation, mobile all-trans chains (termed
mobile trans), and less mobile noncrystalline chains (termed constrained amorphous)
can be quantified by simple 13C NMR experiments on solid polymer samples. The results
are from using a 300 MHz solid-state NMR with 6 kHz magic angle spinning (MAS)
Molecular Structural Characterization of Polyethylene 173
Normalized intensities [a.u.]
1.2
PE + PDMS (coaxial insert) in ODBC-d4
1.0
0.8
Crystallization
Adsorption
0.6
Without graphite
With graphite
0.4
20
60 80 100 120 140
Temperature [°C]
40
160
Figure 5.15 Normalized intensities of (−CH2− + −CH3−) PE/(−CH3) PDMS in TG-NMR experiments
using ODCB-d4 and a temperature array from 160 30 °C in 2 °C steps [431].
L45
L400
Quantitative
Mobile only
Mobile all-trans
Crystalline all-trans
Mobile amorphous
Total interface (%)
25
20
15
10
5
0
Constrained
amorphous
600
200
400
MW (kg/mole)
1400
Rigid only
45
40
35
30
25
20 45
ppm
40
35
30
25
20
ppm
Figure 5.16 Results from modified Elimination of Artifacts in NMR SpectroscopY (EASY) experiments
[435].
speed (Figure 5.16) [433]. Recently, Bruker has developed a commercial 1 GHz solidstate NMR spectrometer with a MAS probe spinning up to 111 kHz; and the potential
for richer information and higher quality results might be available with these advanced
technologies soon, which may have significant impacts in industry.
In the future, we anticipate that the combination of high magnetic field, high MAS
spinning speed, high temperature cryoprobe, and DNP will make NMR characterization of PE easier, faster, and more accurate.
5.5 Polymer Analysis Using Vibrational Spectroscopy
Infrared (IR) and Raman are both vibrational spectroscopic techniques that are now
routinely used in most industrial laboratories for the analysis and characterization of
polymers and polymer-related materials. Both techniques provide functional group
174
Handbook of Industrial Polyethylene and Technology
information about the chemical species present within the sample and enable the elucidation of polymer structure.
5.5.1 Basic Theory of Infrared and Raman Spectroscopy
Infrared spectroscopy uses the absorption of infrared radiation to probe vibrational
energy levels. An infrared spectrum arises due to the absorption of specific wavelengths in the mid-infrared region (4000–400 cm 1) of the electromagnetic spectrum
and requires the presence of a permanent dipole moment. Consequently, IR spectroscopy is more sensitive to side groups and polar functional groups with large permanent
dipole moments and these bands are more intense in the infrared spectrum. In contrast,
Raman measurement is typically based on the inelastic scattering of monochromatic,
visible, or near-infrared radiation and depends on a change in polarizability associated
with the vibration. Raman as a technique is thus more sensitive to the bonds in the
backbone of the polymer chains and double and triple bonds in the polymer repeat
unit. In addition, aromatics, certain inorganic materials, and crystalline and highly
structured materials (both organic and inorganic) are also strong Raman scatterers and
tend to produce strong distinctive bands in the Raman spectrum.
It is important to remember that infrared and Raman spectroscopy are complimentary techniques and when possible should be pursued in conjunction as they respond
differently to different molecular vibrations. In general, symmetric vibrational modes
yield strong Raman bands whereas asymmetric vibrations result in stronger infrared
absorptions. It is not our intent to discuss the theory of infrared and Raman in any level
of detail. Several excellent books that describe both the theory and instrumentation of
infrared and Raman spectroscopy in varying levels of detail are available for interested
readers seeking additional information [434–446].
5.5.2 General Applicability of Infrared and Raman
Spectroscopy to Polymers and Related Materials
The application of infrared and Raman spectroscopy to polymers is not intrinsically
different from the application of these techniques to any other chemical system. The
basic information sought from any vibrational analysis of polymeric materials can be
classified into four general categories: qualitative (chemical) identification, quantitative analysis, structure and morphology elucidation, and online process monitoring (or
in-situ analysis). The choice of applying either technique depends on the problem and
the relative merit of each technique with regard to the problem being addressed.
Ease of sampling is often considered a major advantage of Raman spectroscopy
for analysis of solid samples like polymer pellets, powders, fibers, films, and molded
articles. These samples can be probed by simply illuminating them with a laser beam
and analyzing the 90 or 180 degree (more common) backscattered radiation. In most
cases, no sample preparation, contact, or manipulation is required. Due to modern
accessories, such as attenuated total reflectance (ATR), photoacoustics, and diffuse
reflectance, IR spectroscopy now comes very close to achieving the ease of sampling
offered by Raman [447, 448]. However, IR measurements still require some sample
handling and contact. Depending on the information being sought, this can often be
Molecular Structural Characterization of Polyethylene 175
a disadvantage in polymer analysis; i.e., sample handling can disturb sample orientation. Microspectroscopy is also an area where Raman is often the preferred sampling
method. Lateral resolution on the order of 1 m is routinely achievable with visible
(typically 532 nm) excitation using a confocal Raman microscope. In contrast, resolution of an infrared microscope is on average approx. 10 m [449–451]. The higher
spatial resolution is especially advantageous for analysis of multilayer packing films
or increasingly complex laminates composed of many thin coextruded layers. Raman
spectroscopy, however, does have one major disadvantage that arises when visible
laser radiation is used to excite the Raman scatter. That is, absorption of visible light
can result in fluorescence which in the most extreme cases completely overwhelms
the Raman signal. Highly colored samples, degraded polymers, or samples containing impurities are especially susceptible. The fluorescence problem is typically circumvented by moving to a longer excitation wavelength; but this also causes a decrease
in the spatial resolution and the Raman scattering intensity (sensitivity). Fluorescence
can also be avoided using ultraviolet (UV) excitation, although the risk of “burning”
samples is much higher [452].
In the following sections, some applications in which vibrational spectroscopy is
used to analyze PE and PE copolymers are reviewed. This review is not intended to
be comprehensive. A few key application examples from each general category are
highlighted and adequate references are provided for interested readers.
5.5.3 Qualitative Identification Using Infrared and Raman Spectroscopy
Unknown identification is a key application area of vibrational spectroscopy in industrial problem solving. It is often used to reverse engineer competitive products, troubleshoot process upsets, address product quality issues, identify reaction intermediates
and by-products, and for patent enforcement. A vibrational spectrum can be used as
a molecular fingerprint for chemical identification. Infrared spectroscopy is perhaps
more extensively use for this purpose than Raman. This, in part, is due to the availability of extensive infrared libraries, both commercial and those developed in-house.
In the simplest case, exact identification can be accomplished by acquiring a spectrum
and comparing it with a reference spectrum. This approach generally works best for
a pure component sample. Functional group analysis using group frequency correlations can be used to assign individual peaks observed in an infrared or Raman spectrum of a multicomponent blend or polymer formulation [453–455]. Some knowledge
of the application and/or process chemistry as well as the vibrational frequencies and
intensities of the functional groups is required for proper interpretation. Functional
group analysis can be employed in conjunction with spectral subtraction to aid in the
interpretation of complex multicomponent spectra. Pure component spectra can be
successively subtracted from the mixture spectrum until only one identifiable component remains (a process referred to as “spectral stripping”). Spectral subtraction can
highlight subtle spectral features that are not readily discernible. However, it is also
important to remember that this procedure will lead to some distortion since there are
small differences in peak positions and it is difficult to remove all the peaks belonging
to a particular component from the complex mixed spectrum. This is especially true if
the peaks being subtracted have high absorbance. Software that supports an automatic
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Handbook of Industrial Polyethylene and Technology
multicomponent search is also available; although search results may not always yield
the correct answer and must be interpreted with caution [456, 457].
5.5.3.1
PE Unsaturation and Short-Chain Branching
The infrared and Raman spectra of PE have been extensively studied and peak assignments are generally well established in the literature [458–461]. There are three broad
classes of PE that are segmented by the type and content of branching: HDPE, LLDPE
and LDPE. HDPE has little or no short-chain branches (comonomer). In LLDPE, shortchain branches are incorporated by feeding a comonomer like 1-butene, 1-hexene or
1-octene into the reactor along with ethylene. LDPE is polymerized using a free radical initiator in a high pressure process and has no comonomer but both short – and
long-chain branches. Ziegler-Natta (Z/N) catalyzed PEs can be classified based on the
subtle differences in the level of unsaturation and methyl bending vibration observed
in the infrared spectrum [462, 463]. Since HDPE contains very little to no comonomer,
the methyl bending vibration at approx. 1377 cm 1 is absent or barely visible as a weak
shoulder in the infrared spectrum. This peak can be used to differentiate HDPE from
LDPE and LLDPE. In LDPE the methyl peak at 1377 cm 1 is stronger than the CH2
wagging mode at 1368 cm 1, and the 888 cm 1 peak due to vinylidene unsaturation
is generally stronger than the 910 cm 1 peak due to terminal vinyl unsaturation. In
Z/N-catalyzed LLDPE the peak located at 910 cm 1 is generally stronger. In addition
to differentiating PE into broad classes, differences observed in the infrared spectrum
of HDPE, polymerized using different catalysts, have also been reported [462, 463].
Strong terminal vinyl unsaturation in HDPE (910 cm 1 and 990 cm 1) suggests polymerization using a SiO2-Al2O3-CrO3 type catalyst, whereas a single peak due to transunsaturation (965 cm 1) would indicate a g-Al2O3-MoO2 type catalyst [462]. However,
the authors rightly caution that the relative intensities of the double bonds can change
depending on the polymerization temperature, pressure, and solvent, and the spectrum must be interpreted with caution. Characterization of short-chain branching in
PE using infrared spectroscopy has also been reported for the identification of comonomer. Methyl branches were characterized by an absorbance peak at 935 cm 1, which
is tentatively assigned to the methyl rocking mode. The CH2 rocking mode of the ethyl
branch appears as a small peak at about 770 cm 1 whereas propyl and longer branches
have been reported to give rise to an absorption peak at about 890 cm 1. A distinct peak
at about 770 cm 1 has been observed for polybutene prepared in a pilot scale reactor.
However, information about the comonomer is best obtained using 13C NMR.
5.5.3.2 Ethylene Copolymers
In propylene-ethylene (P-E) copolymers, the CH2 rocking vibrations in the region
between 760–700 cm 1 in the mid-infrared region are very sensitive to the microstructure and monomer sequence distribution. The position of the CH2 rocking mode is
dependent on the length of the methylene sequence units. Absorption bands at about
751 cm 1, 734 cm 1, 726 cm 1 and 720 cm 1 have been assigned to 2, 3, 4 and ≥ 5 contiguous methylene groups, respectively [464, 465]. For example, in P-E copolymers the 734
cm 1 infrared peak arises due to a single insertion of ethylene between a head-to-tail
Molecular Structural Characterization of Polyethylene 177
arrangement of propylene units. Similarly, insertion of two or more ethylene units
results in five contiguous methylene groups and an absorption at 720 cm 1. Generally
speaking, the monomer sequence distribution is catalyst dependent, but the 751 cm 1
and 726 cm 1 peaks, which involve single insertion of ethylene, between head-to-head
or tail-to-tail arrangement of propylene units, are rarely detectable for Z/N-catalyzed
polymers. This region is further complicated by crystallization of the long methylene
runs. In such cases, the single rocking mode centered at about 720 cm 1 is replaced by
the characteristic doublet at about 719 cm 1 and 730 cm 1. This would be expected if the
copolymer is blocky. Infrared peaks in this region can be used to qualitatively distinguish between random and impact PP copolymers, and also form the basis for quantitative estimation of ethylene in these copolymers. A P-E random copolymer generally
shows a single CH2 peak centered at about 734 cm 1, and depending on the catalyst and
level of copolymerized ethylene, a weak shoulder may appear at 720 cm 1. The presence of a characteristic doublet at about 719 cm 1 and 730 cm 1, and a weak shoulder at
734 cm 1 would indicate an impact copolymer.
In addition to the type of comonomer, molecular weight (MW) and molecular weight
distribution (MWD) of the polymer; short-chain branching distribution (SCBD) or
chemical composition distribution (CCD) is also an important parameter for ethylene
a-olefin and polar ethylene copolymers, since it directly affects the properties of the
polymer. High temperature GPC systems with infrared detection are now routinely
used to measure the level of short-chain branching across the entire MWD of the polymer [466–468]. Methyl groups per 1000 carbon atoms (with correction for end groups)
or a calibration of the CH3/CH2 ratio against the wt% comonomer (measured using
NMR) is used to determine the comonomer content as a function of MW. The GPC
effluent is transferred to a high temperature liquid flow cell, placed in a FTIR bench,
using a heated transfer line. A compact, filter-based infrared detector with a built-in
flow cell is also available and replaces the benchtop FTIR spectrometer. However, only
a narrow band of wavelengths is measured, depending on the choice of the filters [469].
A GPC/FTIR interface, based on solvent evaporation and cryogenic deposition of the
solid polymer or additive onto a slowly rotating ZnSe disk, is also available [470]. The
advantage of this approach is that the entire infrared spectrum is acquired, free of solvent interference. For improved signal-to-noise (S/N) the ZnSe plate can be removed
and examined in an infrared spectrometer after the GPC run is over. However, experimental parameters need to be properly selected to ensure uniform deposition of the
solid polymer onto the ZnSe disk.
5.5.3.3 Analysis of Multilayer Films
Transmission infrared microscopy is widely used for compositional analysis of multilayer films. The standard analysis procedure involves cross-sectioning the sample using
a cryo-microtome, placing the section on a salt window, and analyzing the individual
layers. The aperture of the microscope can be adjusted appropriately to isolate the individual layers and eliminate spectral interference from the adjacent layers. However,
layers thinner than 10 m are challenging using traditional infrared microscopy and
rarely yield a pure component spectrum. In such cases, multiple layers can be measured and spectral subtraction used to identify functional groups characteristic of the
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individual layers. This is generally adequate to determine the composition of the individual layers, if a distinct spectral marker is available for each layer. Other techniques
commonly used in conjunction with infrared microscopy, to confirm layer composition in multilayer films, are hot-stage microscopy and DSC. The sample cross-section
is heated using a programmable hot stage and the melting behavior of the individual
layers is observed using a polarized light microscope equipped with an optional CCD
camera. The melting point of each layer is determined visually as indicated by a loss of
layer birefringence. This procedure only works for transparent layers devoid of fillers.
A bulk DSC measurement on the film can further aid in distinguishing between closely
melting polymers.
Slide-on micro tip attenuated total reflectance (ATR) objectives and micro ATR
imaging accessories are now commercially available and have been used for singlepoint analysis, line and area mapping, and 2D imaging of multilayer films. These accessories enhance the spatial resolution by a factor equivalent to the refractive index of the
crystal material (e.g., diamond = 2.4 and germanium, Ge = 4) [471]. The spatial resolution of an infrared (and Raman) microscope is diffraction limited. Using the Rayleigh
criterion, spatial resolution is defined as R ≥ 0.61 × /NA, where is the wavelength
and NA is the numerical aperture. The numerical aperture, NA = n sin , where n is the
refractive index (RI) of the medium in which the optics are immersed and is the half
angle of the cone of light captured by the microscope objective. Since the typical NA of
infrared microscopes is 0.6 to 0.7 the diffraction limited spatial resolution varies across
the infrared spectrum, and is on the order of the wavelength used for analysis. From
the equation shown above, the resolution of a traditional infrared microscope at 1000
cm 1 (10 m) assuming a half angle q of say 30° and n = 1 (RI of air) is 12.2 m. Since
the ATR tip behaves like a solid immersion medium, the diffraction limited resolution
for a Ge ATR tip would be approx. 1.5 m. However, the equation above only calculates
the theoretical limit of the lateral spatial resolution in far-field optical microscopy. In
practice, the optical design and aberrations in the entire optical train (mirrors, objectives, apertures, and detector geometry, etc.) degrade the achievable spatial resolution.
As mentioned earlier, confocal Raman microscopy has a distinct resolution advantage for analysis of multilayer films. A lateral spatial resolution of 1 m is easily achievable. Confocal Raman microscopes utilize 180° backscattering geometry and a confocal
pinhole placed in the back focal plane between the objective and image plane of the
microscope. The pinhole rejects the out-of-focus scatter and spatially filters the analysis
volume of the sample, thereby improving the spatial resolution. Consequently, mapping
through the thickness of a multilayer film by focusing normally onto the surface of the
film, and raising or lowering the sample stage to move the laser focus vertically through
the individual layers, has been proposed [472]. This sampling approach (z-mapping
or optical sectioning) is very attractive in an industrial setting as it would obviate the
need to cross-section the sample and expedite routine sample analysis. However, it is
our experience, as well as that noted by other practitioners, that this sampling approach
is inadequate for analysis of most real-world multilayer films and laminates; especially
when using a dry metallurgical objective [473, 474]. Most multilayer films and laminates used in food packing, for example, have multiple layers of varying thicknesses
and refractive indices. As the laser focus moves a few microns into the sample, due to
refraction both the position and depth of focus increase dramatically. The point of focus
Molecular Structural Characterization of Polyethylene 179
is actually deeper than anticipated, as discussed in recent publications [473–477]. It is
therefore difficult to identify with certainty from which layer the Raman scatter originates in samples with an unknown multilayer structure. It is likely that some layers, e.g.,
thin tie layers, may be missed entirely during the analysis. Applications of infrared and
Raman imaging of laminated or layered products, polymer blends, and other chemical systems are also described in the open literature, but will not be discussed here
[478–486].
5.5.3.4
Analysis of Polymer Gels
The term “gel” is commonly used to define any small defect or imperfection in a
fabricated polymer film or product. PE gels fall into the following main categories
[448, 487, 488]: gels can be highly oxidized (brittle black specks), 2) oxidatively crosslinked polymer, 3) thermally cross-linked, 4) highly entangled (high molecular weight)
polymer that is undispersed but not cross-linked, 5) unmelted resin from the extrusion
process, 6) filler agglomerates from masterbatches, and 7) extraneous contaminations.
Extraneous contaminants may take the form of a different type of resin, metal, wood,
cloth fibers, or dirt. Gels can be generated from many different sources; these include
the converting process, the blending of resin pellets with significantly different shear
viscosities, direct contamination, and from the resin manufacturer. Poorly designed
extrusion equipment and processes are leading causes of oxidatively degraded resins
and cross-linked gels [487].
Analytical techniques commonly used for analyzing gels include polarized light and
epi-fluorescence microscopy, thermo-microscopy, FTIR microscopy, and scanning
electron microscopy (SEM) with an energy dispersive X-ray detector (EDX) for examining inorganic contaminants [489, 490]. Raman is also a viable option for analysis of
gels due to inorganic or foreign contaminants, but is generally less useful for oxidized
and degraded polymeric gels. Protocols established for gel analysis in polymer films are
well documented in the literature [448, 487, 488, 491].
5.5.4 Quantitative Analysis Using Infrared and Raman Spectroscopy
Infrared spectroscopy provides a fast and simple alternative for quantitative analysis,
and is therefore widely used in plant and manufacturing laboratories for quality control and product release. Several ASTM methods have been published for quantitative
analysis of unsaturation [492, 493], branching [494], structure [495], and comonomers
[496–498] in PE. Additionally, infrared methods for measuring additives, density, and
anhydride grafting levels are also reported in the open literature [448, 499–501]. Infrared
transmission is perhaps the simplest and most preferred sampling mode for most quantitative analyses. Depending on the level and nature of the analyte, polymer films of the
appropriate thickness are compression molded using a hydraulic press. Care must be
taken to eliminate interference fringes, especially when thin films are used for analysis.
Fringing occurs due to multiple internal reflections and manifests itself as sinusoidal
variation in the baseline. This reduces the precision of the analysis. Fringing is easily
eliminated by pressing a film between two matted sheets of Teflon or dull sides of a sheet
of aluminum foil. Since infrared (and Raman) are not primary techniques, quantitative
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Handbook of Industrial Polyethylene and Technology
analysis using these techniques typically involves measuring the peak height or area
of the appropriate infrared peak(s), normalizing for film thickness, and establishing
a Beer’s law correlation (calibration curve) based on a primary technique like nuclear
magnetic resonance (NMR), liquid chromatography (LC), or titration. However, semiquantitative analysis used to monitor relative levels of degradation in PE pipes, and to
study additive migration in LLDPE films, has also been reported [502, 503]. For good
quantitation, the absorbance of the analyte and thickness normalization bands must
fall between 0.2–1.0 absorbance units (AU) to ensure linearity of the detector response.
Consequently, thick plaques (500–1000 m) are generally required for the estimation
of low levels of additive while very thin films are needed for quantitative analysis of
polar ethylene comonomers with strong carbonyl absorptions (e.g., EVA or EAA). In
some instances, depending on the level of comonomer, overtone bands must be used
for quantitative estimation [504]. The content of vinyl acetate (VA) in EVA is generally
determined using the carbonyl stretching band at 1740 cm 1. However, the VA content
needs to be below 10 wt% [448]. For content up to 55 wt% the carbonyl overtone at
3460 cm 1 has been recommended [504]. Multivariate calibration techniques, like partial least squares (PLS) [505], have also been utilized in quantitative analysis using mid
– and near-infrared and Raman data [499, 506–510]. PLS is especially useful for calibrating complex spectra, or a component in a mixture with highly overlapping peaks.
5.5.5 PE Morphology
5.5.5.1 Crystallinity
The 1000–1500 cm 1 spectral region shows several distinct peaks in the Raman spectrum of PE. The peak at 1060 cm 1 is assigned to the – C-C – asymmetric stretch associated with the crystalline and amorphous regions of the polymer. The band at 1080 cm 1
is assigned to the gauche conformers present in the amorphous region and can be used
to quantitate the amorphous content in the sample. Peaks from 1400 to 1460 cm 1 correspond to methylene bending vibrations. In this region, the peak at about 1418 cm 1
has been unanimously assigned to the crystalline phase in PE [460, 461, 511, 512]. This
peak, which has been used to approximate the crystallinity [513–516], is a result of
crystal field splitting of the two components of the methylene bending vibration, when
the unit cell is occupied by two structural units, as in the orthorhombic lattices. It was
first reported by Strobl and Hagedron that a simple two-phase model (amorphous +
crystalline) was inadequate for describing all the features observed in the Raman spectrum of PE [513]. The presence of an additional disordered phase of anisotropic nature,
where the chains are in trans conformation but have lost their lateral order, was postulated and has been confirmed by NMR measurements [517, 518]. This disordered
phase is located at the interface between the crystalline and amorphous regions, and
has been associated with chain loops and entangled chain segments. In addition, these
authors also noted that the integrated area of the Raman twisting bands, in the 1295–
1305 cm 1 spectral region, remained the same for the melt as well as a fully crystallized
sample. By using the integrated area of the 1295–1305 cm 1 spectral region as an internal standard to normalize all band intensities, and by curve fitting the Raman spectrum,
these authors were also able to quantitate the three phases. Although conflicting views
Molecular Structural Characterization of Polyethylene 181
have been reported on the reliability of quantitative estimates of the disordered phase
obtained from the decomposition of the Raman spectrum [519–521], the estimation of
crystallinity is believed to be accurate as confirmed in multiple studies [514–516, 519–
521]. The PE crystallinity can be estimated using the relationship; % c = (A1418/AN ×
0.44) × 100, where A1418 is the total integrated area of the 1418 cm 1 Raman band and
AN is the total integrated area of the twisting region [514]. At this point it is important
to note that sample orientation has a significant effect on the Raman peak intensities
and will therefore introduce error in any Raman crystallinity measurement. The error
can be significant for highly oriented samples such as fibers. If the sample is highly oriented or an isotropic sample is otherwise unavailable, the incident laser radiation and
Raman scatter can be depolarized using polarization scramblers. The other alternative
is to slowly spin the sample under the laser beam using a small motor, while acquiring
the Raman spectrum.
In the infrared spectrum of PE, the 720 and 730 cm 1 doublet arises due to correlation splitting, and these peaks can be used to determine the crystallinity [436, 522].
However, the 1894 cm 1 combination band, due to the Raman active CH2 fundamental vibration at 1170 cm 1 and the infrared active CH2 rocking mode at 731 cm 1, is
generally preferred for the determination of crystallinity. The 1894 cm 1 peak is weak
and hence is amenable for estimation of crystallinity in thick samples. The 910 cm 1
vinyl unsaturation peak can be used for path length normalization [436]. The ratio of
these two peaks can be correlated with crystallinity, measured via DSC, X-ray diffraction (XRD), or density. Alternatively, the amorphous content in PE can be determined
using the CH2 wagging mode at 1303 cm 1. This peak is relatively insensitive to orientation and calibration can be accomplished using a liquid paraffinic hydrocarbon like
hexadecane [436, 523].
5.5.5.2 Orientation
Orientation strongly influences the mechanical and optical properties of the polymer,
and is generally defined as the degree of alignment of the polymer chains with respect to
some unique direction. Several techniques are available to characterize the orientation
of the molecular chains in polymers. Molecular orientation can be derived from birefringence, sonic modulus, X-ray (WAXD and SAXS) and polarized infrared, Raman,
and fluorescence [434, 435, 524–526]. Birefringence measures the average (amorphous + crystalline) orientation whereas X-ray diffraction probes orientation in the
crystalline phases. Infrared and Raman spectroscopies can simultaneously probe the
orientation of both phases in the same sample. However, what is actually being measured is the orientation of the transition dipole and the Raman tensor for the individual
vibration, respectively. To determine orientation of the molecules, the angle between
the chain axis and transition dipole or Raman tensor must also be known.
Infrared dichroism has been widely used to quantify orientation in PE films in order
to understand film properties [527–531, 533]. For samples oriented uniaxially in the
machine direction (MD), the procedure involves measuring two infrared spectra; one
with the infrared radiation polarized parallel (AII) to the MD and the other with the
radiation polarized perpendicular to the MD (A ), and calculating the dichroic ratio
(D) as shown in Figure 5.17.
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Handbook of Industrial Polyethylene and Technology
Z
Ch
ain
ax
is
MD
EII
E
ole
n
itio
dip
ns
Tra
Y
ND
TD
D=
X
P2 = P2 cos
=
P2 [chain] = P2 cos
=
1
3 cos2 – 1
2
2
D–1
3 cos2 – 1
D+2
Figure 5.17 Co-ordinate system representing uniaxial orientation.
In order to determine the orientation of the polymer chain, P2 [chain], the angle ( )
between the transition dipole and chain axis must be known. For the two limiting cases
where the transition dipole is either parallel or perpendicular to the chain axis P2 is
given by:
D 1
for
D 2
P2
P2
2
D 1
for
D 2
=0
= 90
(5.23)
(5.24)
The equation for P2 shown in Figure 5.1, is analogous to the Fraser and Hermans orientation function f. P2 varies from −0.5 for perfect alignment perpendicular to MD ( =
90°) to zero for random orientation, to 1 for perfect alignment along the MD. The derivations of these equations and a discussion on orientation analysis in biaxially oriented
samples are available in the literature and will not be repeated here [524, 525, 532, 533].
For vibrations where is 55°, D = 0, and such peaks are insensitive to orientation.
Infrared peaks commonly used to measure crystalline and amorphous orientation in
PE are compiled in Table 5.1 [436].
For quantitative orientation measurement, polarized Raman spectroscopy has an
advantage over infrared in that it can determine both the second P2 (<P2 cosq>) and
fourth P4 (<P4 cosq>) moment of the orientation distribution function (ODF). Being
able to measure P2 and P4 simultaneously helps to characterize the orientation distribution a little better. Although Raman spectroscopy is potentially more revealing for
orientation analysis, it is much more complex, both theoretically and experimentally.
Molecular Structural Characterization of Polyethylene 183
Table 5.1 Transition dipole angles and peak assignments for select PE peaks.
Peak,
cm 1
Transition dipole angle
( ), deg.
Assignment
720
parallel to b-axis
CH2 rock
730
parallel to a-axis
CH2 rock
1078
90°
Amorphous, gauche C-C stretch + CH2 wag
1303
0°
Amorphous, asymmetry CH2 wag (GTG’)
1368
0°
Amorphous, symmetry CH2 wag (GTG’)
1894
perpendicular to c-axis
Crystalline, combination CH2 rock
2016
0°
Crystalline + amorphous, combination CH2
(twist + rock)
This is true for uniaxially oriented samples as well. Typically, multiple polarized Raman
measurements must be made in different scattering geometries and sample orientations
[460, 461]. For biaxially oriented systems, the experiment and data analysis is even more
complex [532, 534]. Consequently, relatively few quantitative orientation studies have
been reported with Raman spectroscopy and most of these have focused on well-characterized polymers like PE [460, 461] and polyethylene-terephthalate (PET) [535–537].
A discussion of the theory and the potential experimental errors in quantitative Raman
orientation measurements, is beyond the scope of this chapter but are discussed elsewhere [460, 461, 532, 534]. Intensity ratios from polarized Raman spectra can suffice in
most applications. They can be used to rank orientation in a given set of samples and/or
to study the relative change in orientation as a function of processing conditions in order
to optimize some desired property. For the localized skeletal deformation vibrations
being considered, a dominant scattering polarizability element may exist. Relative to
the polarization of the incident electric field, the scattered Raman intensity may be very
strong when a bond is parallel to it. Conversely, the scattering intensity (polarized component) may be much smaller and weaker when a bond is perpendicular to it. However,
the entire scattering polarizability tensor should be considered in the analysis of Raman
intensity. If the terms along the bonds dominate, the relative scattering intensity may be
used quite similar to the dichroic ratio generally used in infrared spectroscopy, although
the dichroism observed is not the same as from infrared measurement. In the extreme
case when a single tensor component dominates the scattering, Everall et al. [538] have
shown that the Fraser orientation function f is determined using the equation:
f
R 1
R 4
(5.25)
where R = IZZ/IYY. The subscripts ZZ and YY denote the polarization direction of the
laser beam and the polarization analyzer.
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Handbook of Industrial Polyethylene and Technology
5.5.5.3 In-Situ Monitoring
Generally speaking, mid-IR spectroscopy is not very amenable for online sampling
and suffers from very limited options and poor performance of the chalcogenide fiber
optics in this region. Remote monitoring using these fibers is thus limited to a few
meters. Near-infrared (NIR) sampling is also much better than mid-IR, since fiber
optics similar to Raman can be used. However, the spectral information in the NIR
is generally much less than that available in the mid-IR or Raman, due to broad peaks
and spectral overlap in this region. NIR is therefore typically used in combination with
some multivariate (chemometric) technique to extract the desired information. Raman
essentially offers the best of both techniques. It provides the high degree of spectral
information that is found in the mid-IR, and even better sampling ease than that found
in the NIR, thereby making it the technique of choice for online monitoring. Raman
typically involves the use of visible or near-infrared laser radiation and is consequently
very amenable to fiber-optic coupling using standard, low-cost silica fibers. Using the
appropriate working distance objective, the laser can be focused directly into the process through a transparent window for noninvasive analysis. It is therefore possible
to analyze samples that are normally inaccessible, or process streams situated very
remotely from the spectrometer. Due to the advent of compact diode lasers, notch filters, and sensitive charge coupled detectors (CCD), portable Raman spectrometers are
commercially available from several vendors along with non-contact, filtered fiber optic
Raman probes. Popular Raman process instruments and probe head configurations are
described in recent publications [539–541]. For polymer processing and related applications, this development opens up the possibility of in-situ process monitoring and
studying morphology (crystallinity and orientation) development during fiber and film
fabrication [542–544]. Raman spectroscopy has been used for in-process monitoring of
copolymer composition in a single-screw extruder using a noninvasive Raman probe to
acquire spectra of the melt. A series of ethylene vinyl acetate (EVA) random copolymers
with varying vinyl acetate (VA) content ranging from 2.0 to 43.1 wt% were monitored
and characterized using a robust multivariate calibration suitable for real-time prediction of VA content during processing [545]. Reactive extrusion of glycidyl methacrylate
(GMA) with low density PE (LDPE) was monitored in situ, along the length of the
extruder. The monitoring was done through glass windows on a counter-rotating, twinscrew extruder, using a fiber optic Raman probe. The concentration of GMA monomer
was monitored in different zones, and disappearance of the GMA peaks was assumed to
indicate grafting, homopolymerization, or evaporation of the methacrylate monomer.
Subsequent Raman analysis of the sample at room temperature confirmed GMA grafting onto the LDPE [546]. A Raman system for closed-loop control during the manufacture of titanium dioxide (TiO2) filler has also been reported [547]. The spectroscopic
instrumentation was integrated with a powder sampling system and installed on four,
full-scale manufacturing plants. Fiber coupling (up to 100 m) was employed, and each
spectrometer system was multiplexed to three sampling heads. Online Raman spectra
of LLDPE pellets acquired using a wide area illumination probe were used in conjunction with a multivariate partial least squares calibration model to establish the feasibility of measuring PE pellet density in packed and free-flowing polymer pellets. This
study illustrates the possibility of directly measuring physical properties of polymer
Molecular Structural Characterization of Polyethylene 185
pellets, in situ, during solid conveying used in the chemical industry [548]. Finally,
a fiber optic Raman system was used to monitor LDPE crystallization in real time. A
partial least squares (PLS) model was developed to predict crystallinity, using DSC and
light depolarization measurements as primary techniques. This method allowed crystallinity changes to be measured with a time resolution of 10 s [549].
In an attempt to understand how fiber crystallinity develops in a spin line, online
X-ray patterns, as a function of distance from the spinneret, have been obtained during actual spinning of HDPE fibers [550]. In addition, the crystal structure [551] and
morphology of HDPE and its blends (LDPE/HDPE) have also been studied in real time
[552, 553]. However, most of this work has been accomplished using high-intensity
synchrotron X-ray radiation. Raman spectroscopy does not require an elaborate experimental setup like online X-ray diffraction (XRD) and can provide information on both
the amorphous and crystalline phases of the polymer, provided that a spectral marker
(peak) characteristic of each phase can be unambiguously identified. Since X-ray and
Raman techniques can provide complementary information on the molecular structure
evolution in both crystalline and amorphous phases, simultaneous X-ray diffraction
and Raman measurement systems (both benchtop and synchrotron based) have also
been developed and applied to study crystal phase transitions in PE, n-alkanes and PP
fibers [554, 555]. It is worth mentioning that although there is no absolute method for
the determination of the degree of crystallinity, vibrational spectroscopic techniques
(Raman and infrared) are not considered primary techniques for the measurement
of crystallinity. These techniques are only sensitive to conformational order/disorder
(regularity) along the chain length and are less sensitive to the three-dimensional order
of the crystalline phase. However, if an appropriate spectral marker for the crystalline
phase can be found, it generally correlates well with an independent measure of crystallinity such as density, DSC, or wide-angle XRD [542].
Online Raman spectra, acquired at different points along a fiber spin line, have been
reported and used to monitor the development of crystallinity in HDPE fibers [514].
In PE, the 1418 cm 1 Raman peak is assigned to the orthorhombic crystalline phase
[451, 460, 511–513, 556]. A spectral decomposition procedure proposed by Strobl and
Hagedron [513] was applied to the online Raman spectra, acquired during fiber spinning, to estimate percent crystallinity, as a function of distance from the spinneret [514].
A calibration curve for propylene crystallinity, developed using propylene-ethylene
copolymers, with varying amounts of ethylene, and homo PP resins has also been
reported. This curve was used to monitor spin-line crystallinity development and validate a stress-induced crystallization model [557, 558]. The effect of throughput, spinning speed and resin melt index on the rate and onset of crystallization were studied
[557, 558]. Online Raman has also been used to study crystallinity development in low
LDPE [559–561], LLDPE [561–563], HDPE [560] and LDPE/PP bi-component blown
films [564–566]. Results from some of these studies have been compared with predictions from a flow-enhanced crystallization model. Results were also used to estimate
non-isothermal crystallization half times and to study the effect of crystallization kinetics on the resulting microstructure and mechanical properties of the film [560, 561, 563].
As explained earlier, orientation analysis using Raman is much more complex, both
theoretically and experimentally, even for a uniaxially oriented sample [460, 461].
Consequently, only a few orientation studies have been reported using online Raman
186
Handbook of Industrial Polyethylene and Technology
data [536, 537, 567, 568]. Needless to say, due to the experimental complexity and significant potential measurement errors involved, online quantitative Raman orientation
measurements must be attempted with caution. In most cases, measuring relative orientation using Raman dichroic ratio may be adequate to rank orientation in a given set
of samples and/or study the relative change in orientation as a function of processing
conditions in order to optimize some desired property [541].
As previously noted, mid-IR spectroscopy is less amenable for online analysis using
fiber optic sampling, although significant improvements have been made in recent years,
with the development of better mid-IR optical fibers and ATR tip probes used to monitor reactions in situ. In the application space for polymer processing, several examples
of online mid-IR have been reported for analysis of additives, anti-blocks, lubricants,
slip, copolymer compositions, and fillers in polymer melts [506, 569–572]. A sample
cell to perform noninvasive infrared spectroscopic analysis on moving process streams
containing a polymer melt was patented by Harvey [573]. An infrared spectrum is
acquired by diverting and driving a small stream of the polymer melt through the infrared transmission cell via a transfer line. Fidler et al. have reported online Fourier transform infrared (FTIR) analysis of silica and erucamide in LDPE melt using a 750 m
path length high temperature and pressure flow cell [571, 574]. In addition to speed and
the potential to reduce off-specification production, the online method was reported to
have better precision and was less susceptible to interference from other fillers like talc,
unlike the incumbent off-line gravimetric ashing procedure. Analyses of antioxidants
in HDPE and LLDPE, fillers and vinyl acetate content in ethylene vinyl acetate (EVA)
copolymer have also been reported [569]. Chalmers has reported compositional monitoring of propylene/ethylene copolymer melt streams through a specially designed
1 mm path length transmission cell, heated to 240 to 260 °C and fitted with 6 mm thick
zinc selenide (ZnSe) windows. A set of forty samples was used to develop an off-line
PLS model using NMR as the primary method. This model was used to predict ethylene
content in the copolymer melt stream [506]. Applications of online near-infrared (NIR)
spectroscopy for quantitative analysis of additives and fillers and quality monitoring of
recycled plastic waste during extrusion have also been reported [570, 575].
5.6 Emerging Techniques
In atomic force microscope infrared (AFM-IR) spectroscopy a sample is mounted on
an infrared transparent prism, in a configuration similar to a conventional ATR, and
illuminated with a pulsed tunable mid-infrared laser. The laser is continuously tunable
over a wide mid-infrared range from 3600–1200 cm 1. An AFM cantilever tip placed
above the sample is used to detect the absorbed infrared radiation, thereby providing
nanoscale spatial resolution [576, 577]. The principle of AFM-IR is based on photothermal induced resonance (PTIR) [578, 579]. Absorption of the infrared radiation,
by the sample, causes the sample to heat up, resulting in rapid thermal expansion
which causes resonant oscillation of the cantilever. The position of the cantilever tip is
detected using a laser beam and a position-sensitive detector, like in a traditional AFM.
The induced cantilever oscillations decay in a characteristic ring down, which is Fourier
transformed to extract the amplitudes and frequencies of the oscillations. A plot of
Molecular Structural Characterization of Polyethylene 187
the amplitudes of the cantilever as a function of the source wavelength creates a local
absorption spectrum. AFM-IR absorption spectra are direct measurements of the sample absorption and correlate well with conventional, bulk IR spectra. Application of
AFM-IR to study multilayer films, polymer composites and nanodomains in commercial impact copolymers have been reported recently [580, 581, 582].
Acknowledgments
We would like to thank Dr. Xiaohua Qiu of The Dow Chemical Company for helpful
comments and discussions.
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518. Kitamura, R., Horri, F., and Murayama, K., Phase Structure of Lamellar Crystalline
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537. Chase, B., Kennedy, A., Abstracts of 24th FACSS Conference, Paper 408, 169, 1997.
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541. Paradkar, R.P., Sakhalkar, S.S., He, X., and Ellison, M.S., On-Line Estimation of Molecular
Orientation in Polypropylene Fibers Using Polarized Raman Spectroscopy, Appl. Spectrosc.,
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542. Paradkar, R.P., Leugers, A., Patel, R., Chen, X., Meyers, G., and Lipp, E., Raman Spectroscopy
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Sons Ltd.: Hoboken, NJ, 2009.
543. Bart J.C., Plastics Additives: Advanced Industrial Analysis, p. 663, ISO Press, 2006.
544. Gururajan, G., and Oagle A.A., Raman Applications in Polymer Films for Real-time
Characterization, in: Raman Spectroscopy for Soft Matter Applications, Amer, M.S. (Ed.), p.
33, John Wiley & Sons, Inc.: Hoboken, NJ, 2009.
545. Barnes, S.E., Brown, E.C., Sibley, M.G., Edwards, H.G.M., and Coates, P.D., Vibrational
Spectroscopic and Ultrasound Analysis for the In-Process Monitoring of Poly(ethylene
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546. Jarukumjorn, K., and Min K., SPE ANTEC Tech. Paper, 1006, 2064, 2000.
547. Clegg, I.M., Everall, N.J., King, B., Melvin, H., and Norton, C., On-Line Analysis Using
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548. Kim, J., Kim, Y., and Chung, H., Direct On-Line Raman Measurement of Flying Solid
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551. Liu, L., Paradkar, R., and Bensason, S., Crystal Structure Transformations and Orientation in
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552. Clements J., Zachmann H.G., and Ward, I.M., Study of the Melting Behaviour of UltraHigh Modulus Linear Polyethylene Using Synchrotron Radiation, Polymer, 29, 1929,
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553. Song, H.H., Stein, R.S., Wu, D.Q., Ree, M., Phillips, J.C., LeGrand, A., and Chu, B.,
Time-Resolved SAXS on Crystallization of a Low-Density Polyethylene/High Density
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554. Tashiro, K., Kariyo, S., Nishimori, A., Fujii, T., Saragai, S., Nakamoto, S., Kawaguchi, T.,
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of X-Ray Diffraction and Raman Spectra: Application to Structural Study of CrystallinePhase Transitions of Chain Molecules, J. Polym. Sci., Part B: Polym. Phys., 40, 495, 2002.
555. Ran, S., Fang, D., Sics, I., Toki, S., Hsiao. B. S., and Chu, B., Combined Techniques of
Raman Spectroscopy and Synchrotron Two-Dimensional X-ray Diffraction for In Situ
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52, 3425, 1970.
557. Paradkar, R.P., Patel, R.M., Knickerbocker, E., and Doufas, A., Raman Spectroscopy for
Spinline Crystallinity Measurements. I. Experimental Studies, J. Appl. Polym. Sci., 109,
3413, 2008.
558. Patel, R.M., Doufas, A.K., and Paradkar, R.P., Raman Spectroscopy for Spinline Crystallinity
Measurements. II. Validation of Fundamental Fiber-Spinning Models, J. Appl. Polym. Sci.,
109, 3398, 2008.
559. Cakmak, M., Serhatkulu, F., Graves, M., and Galay, J., SPE ANTEC Tech. Papers, 43, 1794,
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561. Henrichsen, L.K., McHugh, A.J., Cherukupalli, S.S., and Ogale, A.A., Microstructure
and Kinematic Aspects of Blown Film Extrusion Process: II. Numerical Modelling and
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Spectroscopy During Blown Film Extrusion of a Linear Low-Density Polyethylene, Polym.
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563. Cherukupalli, S.S., Gottlieb, S.E., and Ogale A.A., Real-Time Raman Spectroscopic
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6
Thermal Analysis of Polyethylene
Kevin Menard1* and Noah Menard2
1
Mettler Toledo, Columbus, Ohio, USA
Veritas Testing and Consulting LLC, Denton, Texas, USA
2
Contents
6.1 Introduction ............................................................................................................................... 218
6.2 Differential Scanning Calorimetry (DSC) ............................................................................. 218
6.2.1 Glass Transition and Melting Temperature ............................................................... 221
6.2.2 Heat Capacity Measurements ...................................................................................... 224
6.2.3 Crystallization Studies .................................................................................................. 225
6.2.4 Oxidative Induction Time (OIT) ................................................................................ 226
6.3 Thermogravimetric Analysis (TGA) ...................................................................................... 228
6.4 Thermomechanical Analysis (TMA) ...................................................................................... 228
6.4.1 Coefficient of Thermal Expansion............................................................................... 229
6.4.2 Softening Point, Heat Distortion and Other Tests .................................................... 230
6.5 Dynamic Mechanical Analysis (DMA) .................................................................................. 230
6.5.1 Temperature Scans – Modulus and Transition Temperatures ................................. 231
6.5.2 Frequency and Other Scans ......................................................................................... 232
6.6 Coupled Thermal Techniques.................................................................................................. 233
6.6.1 Spectral DSC .................................................................................................................. 233
6.6.2 Evolved Gas Analysis (EGA) ....................................................................................... 235
6.7 Conclusions................................................................................................................................ 235
References ............................................................................................................................................ 236
Abstract
Thermal analysis is a key technique for the characterization of polyethylene (PE) and its copolymers. A brief summary of each technique and the major applications of DSC, TGA, TMA
and DMA are discussed. Examples of key data and interpretation are shown. Hyphenated or
coupled techniques useful for the analysis are also introduced. References to standard methods
are included.
Keywords: Thermal analysis, DSC, TGA, TMA, DMA, TG-IR, TG-MS, TG-GCMS
*Corresponding author: kmenard@gmail.com
Mark A. Spalding and Ananda M. Chatterjee (eds.) Handbook of Industrial Polyethylene and Technology, (217–238)
© 2018 Scrivener Publishing LLC
217
218
Handbook of Industrial Polyethylene and Technology
6.1 Introduction
Polyethylene and its copolymers have long been studied using thermal analysis techniques. However, ever since PE became commercially available in 1939 [1], many of
the testing methods developed for it occurred prior to the development of modern
thermal analysis techniques. Techniques like melt index (MI) and solid density [2]
remain important in the industry despite the advantages modern rheological and thermal methods may bring. As the industry developed, thermal techniques became more
common and today a range of thermal instruments are used by the polyolefin industry. Because PE is available in a wide range of densities, which all have different melting behavior, infrared spectroscopy is not able to identify the materials as the various
grades of PE. So another method is needed, such as differential scanning calorimetry
(DSC), which allows the identification of linear low density PE (LLDPE), low density
PE (LDPE), medium density PE (MDPE), high density PE (HDPE), and ultra high
density PE (UHDPE) by its melting behavior.
In general, the majority of the methods used in the thermal analysis of PE and its
copolymers are either the American Society for Testing Materials (ASTM) methods,
their International Organization for Standardization (ISO) derivatives, or an industrial
variation of them. Table 6.1 lists by technique the ASTM methods for thermal analysis
in general as well as those specific to PE. These important test methods are discussed
in the next sections.
6.2 Differential Scanning Calorimetry (DSC)
In general, DSC or its analog, differential thermal analysis (DTA), remain the primary
thermal measurement tool used in the polymer industry. For PE and its copolymers,
they are used for a variety of experiments to look at the thermal transitions, crystallization behavior, and resistance to oxidation [3]. The instrumental principles behind
DSC vary between power-compensated and heat flux models: the results of the measurements are the same and the data are treated identically. Heat flux DSC measures
the temperature difference between a sample and a reference, which is normally an
empty pan or crucible using a single furnace with two sensors. Power-compensated
DSC measures the energy needed to keep a sample-containing furnace at the same
temperature as a reference furnace by constantly adjusting the sample furnace to
match. This gives improved responsiveness but at the cost of baseline stability and
sample size. As both instruments give the same information, we will only differentiate
when necessary.
Calibration methods are normally used to either the manufacturer’s specifications
or to an appropriate standard. Temperature and enthalpy calibration are normally done
with pure metals. The common metals include indium with a melting temperature
of 156.6 °C, tin with a melting temperature of 231.9 °C, and zinc with a melting temperature of 419.5 °C. Gallium is often used for standardization near ambient standard
since it has a melting temperature of 29.7 °C [4]. For low temperature work, mercury
(−38.83 °C), which has several safety issues, has been replaced with compounds like
adamantane, water, or n-octane. Indium is available as a NIST traceable standard for
Thermal Analysis of Polyethylene 219
Table 6.1a ASTM thermal methods for PE: differential scanning calorimetry (DSC).
D3418
Transition Temperatures and Enthalpies of Fusion and Crystallization of Polymers
by DSC
D3895
Oxidative-Induction Time of Polyolefins by DSC
D7426
Procedure for Determining Tg of a Polymer or an Elastomeric Compound by DSC
E967
Temperature Calibration of Differential Scanning Calorimeters and Differential
Thermal Analyzers
E968
Heat Flow Calibration of Differential Scanning Calorimeters
E1269
Determining Specific Heat Capacity by Differential Scanning Calorimetry
E1356
Assignment of the Glass Transition Temperatures by Differential Scanning
Calorimetry
E2070
Kinetic Parameters by Differential Scanning Calorimetry Using Isothermal
Methods
E2253
Enthalpy Measurement Validation of Differential Scanning Calorimeters
F2625
Measurement of Enthalpy of Fusion, % Crystallinity, and Melting Point of
UHMWPE by DSC
Table 6.1b ASTM thermal methods for PE: thermogravimetric analysis (TGA).
D3850
Rapid Decomposition of Electrical Insulating Material by TGA
E1131
Method for Compositional Analysis by TGA
E1582
Calibration of Temperature Scale for TGA
E2105
Standard Practice for TG-IR
E2401
Method for Mass Loss and Residue Measurement Validation of TGA
E2550
Test Method for Thermal Stability by TGA
those needing ISO certification.1 Heat flow and heat capacity are also calibrated with
pure metals or with sapphire.
Calibration requirements depend on several factors specific to the user and to the
application. In general, DSC calibrations are done on a less frequent basis as required by
manufacturer’s recommendations or by the appropriate operating standard. The application can greatly affect this: simple melting runs on clean samples normally allow both
low levels of maintenance and fairly infrequent calibrations. In contrast, oxidative studies or running contaminated samples will mean frequent cleaning and recalibrations.
Calibration checks, however, must be frequent and checking standards either daily or
1
A traceable standard implies a material that has a valid traceable history to the National Institute for
Standards and Testing (NIST). A material from a vendor or from a round robin test that lacks this history
is not considered traceable. This becomes important to meet laboratory quality requirements for ISO and
other certification.
220
Handbook of Industrial Polyethylene and Technology
Table 6.1c ASTM thermal methods for PE: thermomechanical analysis (TMA).
D3386
CTE of Electrical Insulating Materials by TMA
E 228
CTE by TMA with Silica Dilatometer
E 831
CTE of Solids by TMA
E1363–03
Standard Test Method for Temperature Calibration of TMA
E1545–05
Standard Test Method for Assignment of the Glass Transition Temperature by
TMA
E 1824
Tg by TMA in Tension
E1953–07
Standard Practice for Description of Thermal Analysis and Rheology Apparatus
E2092–04
Standard Test Method for Distortion Temperature in Three-Point Bending by
TMA
E2113–04
Standard Test Method for Length Change Calibration of TMA
E2206–06
Standard Test Method for Force Calibration Of TMA
E2347–05
Standard Test Method for Indentation Softening Temperature by TMA
E831–06
Standard Test Method for Linear Thermal Expansion of Solid Materials by
TMA
Table 6.1d ASTM thermal methods for PE: dynamic mechanical analysis (DMA).
D4065
Determining DMA Properties Terminology
D4092
Terminology for DMA Tests
D4440
Measurement of Polymer Melts
D4473
Cure of Thermosetting Resins
D5023
DMA in Three Point Bending Tests
D5024
DMA in Compression
D5026
DMA in Tension
D5279
DMA of Plastics in Torsion
D5418
DMA in Dual Cantilever
D7028
Glass Transition Temperature of Polymer Matrix Composites
D7750
Cure Behavior of Thermosetting Resins by DMA using an Encapsulated Specimen
Rheometer
E 473
Terminology for Thermal Analysis
E-756
Measuring Vibration-Damping Properties of Materials
E 1640
Tg by DMA
E 1867
Temperature Calibration for DMA
Thermal Analysis of Polyethylene 221
every 10 to 20 runs is not uncommon. As long as the melting onset and heat of fusion
remain within acceptable limits, recalibration is unnecessary.
The most common concern with calibration is how precise the calibration has to be.
This is dependent to a great degree on the application for precision and accuracy. If the
instrument cannot meet the manufacturer’s specification, one should look for a cause.
However, the values the manufacturer gives may be too narrow or too wide for an
application and time between calibrations may need to be adjusted for that. Similarly,
routine maintenance such as cleaning the furnace and the sensor, the underside of the
furnace lid, and possibly the exit gas line, should be done as needed. Applications that
produce soot or smoke will require more frequently cleaning of the chambers as compared to cleaner analysis such as melting.
Data storage requirements have changed over the years as disk drives have become
larger and cheaper. While instruments allow data sampling rates of fractions of a second, the availability of large drives means this is seldom a limiting factor. Even so, frequent backups are still a good idea particularly if using older instruments. Normally
collecting one point per second is adequate except for the extreme rates used in fast
scanning DSC. Data security varies depending on whether the data is stored in a file
system or in a database. In any case, original data should not be alterable: any changes
or data smoothing should be detectable. For those working in the food or drug areas,
most systems offer 21 CFR 11 compliant software.2
Generally, operation of a DSC is straightforward. A sample is prepared and placed
in an appropriate pan, not filling it more than half way. Pan type, like sample preparation, is dependent on the form of the material. Powders are added as is and loaded into
a standard crucible with a lid. Pellets can be handled the same way after cutting into
small pieces. Alternatively they can be cold pressed into flat shapes. Films and fibers
are cut to fit the crucible and normally loaded in a crucible that allows the upper lid
to compress them slightly to prevent them from pulling up and away from the pan on
heating. Normally a heat-cool-heat cycle is used: the initial heat tells us about the state
of the material as it was received. The second heat, after a controlled cooling from the
melt, tells us what the material looks like with a standardized heat history. This allows
comparisons between batches of material on the same basis.
6.2.1 Glass Transition and Melting Temperature
Two fundamental measurements of polymer in DSC are the glass transition and melting temperatures. Both are strong indicators of molecular structure. The glass transition
temperatures in PE are subambient and have been reported to range from −120 °C for
HDPE to −60 °C for LLDPE [5]. Figure 6.1 shows a typical Tg analyzed by DSC. Note
two concerns with the glass transition. First is that the glass transition lists multiple
ways it can be calculated, even in DSC. This is more complicated when other methods
are used [39]. While many give similar values, this is not always true and it is important
to define which method is used. As linear PE tends to be highly crystalline, transitions
2
Electronic recordkeeping requirements of the FDA are defined in this US government document. Full
text of the Code of Federal Regulations Title 21, Chapter 1, Part 11 can be found at https://www.accessdata.
fda.gov/scripts/cdrh/cfdocs/cfcfr/CFRSearch.cfm?CFRPart=11
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Handbook of Industrial Polyethylene and Technology
Heat flow, mW
Integral
–1147.40 mJ
Normalized –65.13 Jg–1
Onset
102.75
Peak
114.29
Glass transition
Onset
–63.29 °C
–60.83 °C
Midpoint ISO
–60.82 °C
Midpoint DIN
Midpoint ASTM, IEC –62.94 °C
Midpoint Richardson –62.96 °C
0.213 Jg–1K–1
Delta cp DIN
0.178 Jg–1K–1
Delta cp ASTM, IEC
Delta cp Richardson 0.178 Jg–1K–1
–140–120–100 –80 –60 –40 –20
0
20
40
60
80 100 120
Temperature, °C
Figure 6.1 The glass transition and melting peak of MDPE is shown on heating (solid line). Endotherms
are down unless otherwise noted. (Reprinted with permission from Mettler-Toledo, LLC, Columbus, Ohio)
are weak and depending on the instrument and particular polymer, fast scanning techniques or DMA may be needed to observe it. The second issue concerns the value for
the Tg itself. This is a fairly complex topic and multiple values are reported. Schneider
[6] summaries the current understanding as the PE Tg is reported as 2 values, a Tg(U)
and a Tg(L); U is for upper and L is for lower. The Tg(U) appears to be centered around
−70 °C while the Tg(L) appears around −120 °C. Boyer [7] attributes the Tg(L) to a sub
Tg transition. Various sources report differing values experimentally: This complexity
and the requirement of cooling to below a temperature of −140 °C to meet ASTM testing protocol make the use of the melting peak more common.
The general procedure for measuring a Tg as recommended by the ASTM and ISO
methods is to cool the material to 30 to 40 °C below the expected transition to allow for
stabilization and the start-up transition, and then scan to at least 20 °C above at a rate of
10 to 20 °C/min. Cooling to −140 °C requires liquid nitrogen cooling for the DSC and
many labs prefer not to use liquid nitrogen due to expense and safety concerns. Also,
historically, PE has been classified by density and melting point and these values remain
popular today due in part to the large database of these values. Values for the Tg(U) and
the melting point are shown in Table 6.2. One should note, however, that sometimes
only the Tg(L) is seen, as we shall see in Figure 6.10.
The melting point of various grades of PE is probably the best indicator of both
the grade and its ability to be processed under certain conditions [13]. Particularly
where blending of grades is used to obtain certain properties or in recycling where
films may contain multiple layers, DSC represents a more precise way of determining
the grade and/or percent of that grade in a batch. Figure 6.2 shows an overlay of several grades of PE while examples of the melting temperature are tabulated in Table 6.1.
A rough estimate of percentages can also be calculated on blends using the fraction of
the peak areas, as shown in Figure 6.3. This blend of 65% PE and 35% propylene (PP)
was a three-layer film with the PP layer in the middle. It was checked by both DSC, as
shown in Figure 6.3, and also by FTIR using an ATR. The DSC scan indicated that the
Thermal Analysis of Polyethylene 223
Table 6.2 Examples of PE Tg and Tm values (values taken from refs [8–12]).
PE grade
Tg in C
Tm in C
LDPE
133 to 103a,b/ 85e
105 to 118a
LLDPE
110a/ 80 to 50e
120 to126a,b
m-LLDPE
45 to 30d
110 to 115c, d
MDPE
110a/ – 80 to 50e
115 to125a
HDPE
120 to 110a,b/ 70 to 50e
125 to 135a
UHMWPE
110a
130 to 135a
Heat flow
Superscripts refer to specific sources in references [8–12].
LDPE
Peak 108.81 °C
LLDPE
Peak 123.48 °C
HDPE
Peak 132.89 °C
60
70
80
90
100
110
120
130
140
150
Temperature, °C
Figure 6.2 Overlay of grades of PE showing the ability of the DSC to differentiate them. (Courtesy of
Veritas Testing and Consulting, LLC)
4
Heat flow, mW
6
161.8°c
8
Total Area 181.7 - Delta H102.1 J/g
Percent to 134.8 C - Area 121.4 - % 66.80
10
100
127.8°c
110
120
130
140
150
160
170
180
190
200
Temperature, °C
Figure 6.3 Melting curve of a PE blend showing the use of peak areas to estimate the amount of each polymer
in the mixture. (Reprinted with permission from Veritas Testing and Consulting, LLC, Denton, Texas)
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Handbook of Industrial Polyethylene and Technology
composition was 67% PE. The FTIR scan showed at first only PE. After pressing the
material into a film thin enough to get transmission data, the FTIR showed both PE
and PP. The DSC scan was considerably quicker and gave reasonable agreement to the
original blend.
The enthalpy of melting (ΔHf , also called heat of fusion) which is reported in
Figure 6.3 is related to the degree of crystallinity of the polymer. While more important
for semicrystalline polymers like PET that exhibit cold crystallization, it can be used to
estimate the degree of crystallinity of PE or of a copolymer. Since the theoretical value
of purely crystalline PE has been calculated, one can ratio the measured enthalpy to
the theoretical enthalpy for 100% crystalline PE and get a mass fraction crystallinity
for the sample. This detection of crystallinity by the DSC makes it useful for studies of
crystalline behavior.
The calculation is straightforward. The mass fraction crystallinity (Xc) of PE is determined from the DSC run by the equation using the enthalpy value measured:
Xc
H f ,m
H f ,100
(6.1)
where Hf,m is the measured heat of fusion of the PE sample and Hf,100 is the heat of fusion
of 100% crystalline PE. The value of Hf,100 for PE is known to be 293 J/g [14]. So for a
measured sample with a heat of fusion of 278 J/g, we would calculate 94.8% crystallinity. Software programs do exist that will do this for you and also calculate the percent
crystallinity as a function of time.
6.2.2 Heat Capacity Measurements
One of the most important uses of DSC in addition to determining the melting point of
PE (and hence the grade) is the measurement of the heat capacity (Cp) of the material.
Heat capacity and coefficient of thermal expansion for a blend are needed to develop
models for optimizing mold design and extruder performance [15]. Heat capacity is defined as the number of J of heat needed to increase a kg of material by 1 °C.
Traditionally, it is measured in the DSC by either a scanning or an isothermal method.
In the scanning method, an accurately measure weight of material is heated across the
region of interest at 1 to 2 °C/min in pans of a known Ro; where Ro is the thermal resistance of the pans. Pans should be closely matched by weight. A baseline using the empty
pans and a sapphire standard are also run at the same conditions. Ideally, the sapphire
standard is run in the same pans used for the baseline run. From this, the heat capacity of the material is normally calculated using the ASTM defined method E-1269. The
isothermal method is slower but gives better data as the sample is heated to a series of
set temperatures and then held there until it has equilibrated. Values are calculated as
before using a baseline and sapphire run.
With the development of modulated temperature techniques, many analysts are
using these techniques to simplify the measurement of the heat capacity. When modulate temperature DSC methods are used, a non-linear heating rate is applied to the
sample. This can be a heating step followed by an isothermal hold, a single frequency
Thermal Analysis of Polyethylene 225
oscillation, or a multiple frequency oscillation depending on the method used [16].
The latter technique has the advantage of allowing one to determine the frequency
dependence of the glass transition and has been applied to UHMWPE composites
[17, 18]. One can calculate a Cp value equivalent to a traditionally measured value from
the run. These values have been reported to be in good agreement with the 3-curve
method above and are simpler to run [19, 20]. In PE, variations in the polymer’s degree
of crystallinity, branching, cross-linking, the blend of grades used, and the additive
packet can all affect the measured Cp [21].
6.2.3 Crystallization Studies
Many commercial processes require an understanding of the rate of crystallization in
the PE being used. As the commercial polymer includes an additive package that has
antioxidants, nucleation agents, flow enhancers, and others, these can show strong differences in behavior. As this behavior is driven by the cooling rate, excellent control in
the DSC on cooling is required. Two types of studies are commonly done.
In the first type of cooling study, the sample is heated above the melting temperature,
held a short time (about 2 to 4 min) to melt any residual crystals, and then cooled at a
controlled rate. The peak temperature during cooling is taken as the crystallization temperature (Tc). On reheating, one can calculate the resulting crystallinity [4]. Figure 6.4
shows an example. Cooling tends to be a much more sensitive probe of the material
than heating and it is not uncommon to see materials that show a single peak on heating
show multiple peaks on cooling. This method is limited to the maximum cooling rate
of the instrument, the best of which was 750 °C/min in the PerkinElmer DSC8500 .
However, with the introduction of chip calorimeters like the Mettler-Toledo Flash
DSC , extremely high rates of up to 2,400,000 °C/min can be obtained. Work with
these instruments has produced some very interesting effects [22, 23] for PE, although
completely amorphous PE has not yet been obtained [24]. Highly amorphous specimens have been produced with this instrument. The results agree closely with work by
Geil et al. where drops of molten PE were quenched in liquid nitrogen [25, 26].
916.65 mJ
Integral
Normalized 52.04 Jg–1
101.90 °C
Onset
96.31 °C
Peak
Onset 103.61 °C
Peak
99.80 °C
Heat flow, mW
Heat flow, mW
Peak 43.87.16 °C
Cooling at 5 °C/min
Heating at 5 °C/min
Peak 69.16 °C
Heating at 100 °C/s
Integral
–824.97 mJ
Normalized –46.83 Jg–1
103.20 °C
Onset
114.29 °C
Peak
0
(a)
10 20 30 40 50 60 70 80 90 100 110 120
Temperature, °C
Cooling at 100 °C/s
Onset 107.73 °C
Peak 118.03 °C
40
(b)
50
60
70
80 90 100 110 120 130
Temperature, °C
Figure 6.4 Controlled cooling from the melt and reheating to study the crystallization of PE: (a) scan is
a run in standard DSC mode at 5 °C/min cooling and heating, and (b) a Flash DSC run at 6000 °C/min
cooling and heating. Note the presence of a lower temperature peak in the cooling curve not seen in the
heating curve.
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Handbook of Industrial Polyethylene and Technology
The second study type is an isothermal recrystallization [4]. The data are presented
as isothermal crystallization half times (ICHT). The basic experiment heats the polymer to above the melting temperature in the DSC and holds it there for a short time
(about 2 to 4 minutes) to allow it to fully melt. The material is then cooled as quickly
as possible, ideally at a controlled rate so the process can be modeled, to a temperature below the melting temperature. The crystallization of the material at that temperature is then tracked over time as shown in Figure 6.5. The time corresponding
to the crystallization peak (that is, when half of the crystallization is completed) is
the measured ICHT. Figure 6.5a shows a scan with an isothermal hold portion of the
experiment for a high density PE sample run on a conventional power-compensated
DSC at a cooling rate of 500 °C/min. Figure 6.5b shows a scan of the same material
run using a power-compensated chip calorimeter at a cooling rate of 60,000 °C/min.
The faster cooling rate allows lower isothermal temperatures to be used: maximum
cooling rates available today are 750 °C/min in conventional units and 240,000 °C/
min in chip calorimeters such as the Mettler-Toledo Flash DSC. While convention
power compensation measures ICHT in the range of 1.9 to 12 minutes, the chip calorimeter measures from 0.01 to 0.05 minutes. Use of faster cooling rates is possible
with subambient cooling.
Both Tc and ICHT measured by DSC are useful parameters for characterizing crystallization behavior; e.g., nucleating efficiency. As Tc increases and correspondingly the
ICHT decreases, this indicates faster crystallization rates. Studies on nucleation effects
in polyolefins have shown these values can be correlated with nucleation density as well
as each other [27].
6.2.4 Oxidative Induction Time (OIT)
A serious concern with any polymer is its oxidation during both processing and use
of the product. Because of the concerns with loss of desirable properties on oxidation, oxidative induction time (OIT) testing is a common use of DSC for PE products,
15
109 °C
121 °C
Heat flow, mW
16
Heat flow, mW
17
122 °C
18
19
110 °C
111 °C
123 °C
112 °C
20
21
(a)
0
5
10
15
Time in minutes
20
0.00
(b)
0.02
0.04
0.06
0.08
Time in minutes
0.10
0.12
Figure 6.5 Isothermal crystallization half time (ICHT) measurement of a high density PE: (a) a
conventional DSC at cooling rate of 500 °C/min, and (b) a scan from a chip calorimeter at a cooling rate
of 60,000 °C/min. Temperatures reported are the isothermal holds.
Thermal Analysis of Polyethylene 227
including food packaging, cable jackets, electronics, consumer packaging, and many
others [28]. This test is commonly used to track the effect of in-service aging such
as PE pipes that are exposed to hot water, or food containers that are exposed to
chlorine and detergents. The parts are typically tested after limited periods of exposure. As the stressing agent affects the material, these changes can be seen in decreased
OIT compared to the unexposed material. There is an inverse relationship with OIT
and branching; i.e., lower density grades will degrade faster under thermal oxidative
stress [29].
OIT tests are relatively simple to run although data analysis can be tricky, and erroneous and unexpected results are not uncommon. OIT is commonly done by two techniques: DSC and high pressure DSC (HP DSC). Conventional DSC experiments are
run as follows: under an inert gas, heat the material to the test temperature and hold.
After a designated time sufficient to obtain a stable baseline, normally 1 to 2 minutes,
the purge gas is switched to either air or oxygen (preferred), and then measure the
time until oxidation is observed, as shown in Figure 6.6. The HP DSC version uses
an elevated pressure of oxygen to shorten testing times: this is normally done by rapidly bringing the sample to the temperature of interest at 1 atm and then increasing
the pressure to that of the test. Both methods are valid and comparisons of the two
methods have been reported [30].
A less common but more precise method of measuring OIT in PE is the use of chemiluminescene to detect the start of oxidation. Like most polymers, PE is known to
emit weak light during oxidation [31, 32]. As increasing oxygen pressure increases the
light emitted, it is often run in a pressure DSC. Light emission is tracked using a specialized camera. The appearance and intensity of the light emitted is then used to calculate the OIT value. Both scanning and isothermal methods are used. Studies on the
thermal and photostability of PE by chemiluminescene have been used to examine the
efficiency of antioxidants for both ultra violet (UV) [33] and thermal degradation [34].
The main disadvantage of the technique is the expense of the system and the need for
manual operation.
10
Heat flow mW
14
18
Oxygen ON
OIT = 9.0
22
0
2
OIT = 3.4 min
4
6
8
10
12
14
Time in minutes at 100 °C
Figure 6.6 Oxidative induction time (OIT) on a PE sample at 210 °C. HDPE was exposed to high
temperatures for 240 hours (solid line) and 24 hours exposure (dashed).
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Handbook of Industrial Polyethylene and Technology
6.3 Thermogravimetric Analysis (TGA)
Thermogravimetric analysis (TGA) uses a high sensitivity balance in a controlled temperature and atmosphere oven to measure the weight loss or gain of a sample under
controlled conditions [14]. In many cases, these instruments also measure heat flow at
the same time and are referred to as a simultaneous thermal analyzer (STA). Calibration
can be done by using the Curie points of materials such as iron (780 °C) and nickel
(354 °C) [35] or by melting standards as in DSC. As traceable Curie standards are not
currently available and the typical accuracies of Curie point values are ±3.0 °C as compared to ±0.5 °C for melting points, many laboratories prefer to use melting standards.
Although it is occasionally used for OIT measurements, the vast majority of uses
of TGA for PE is for compositional analysis. In general, for PE and its copolymers,
TGA is used to detect filler content although there is some use for the method for
energy applications [36]. An example of compositional analysis is shown in Figure 6.7
where the amount of glass fiber is determined in a HDPE composite. Methods tend to
be straightforward: heating ramps to 1000 °C with a switch to air or oxygen at about
600 °C. Weight loss (or gains in the case of OIT) is calculated for each step seen and
for any remaining residue. In this example, the unfilled material (solid line) has a small
amount of loss before the burn, and then burns to almost no residual. The filled material has a residual of about 31%, corresponding to the glass fiber composition.
6.4 Thermomechanical Analysis (TMA)
Thermomechanical analysis (TMA) was once used extensively to mimic older tests like
heat distortion, softening point, or Vicat softening point. However, with the use of DSC
for determination of Tg and Tm, today TMA is used mostly for the measurement of the
coefficient of thermal expansion. Tg measurements are still used for weak transition,
for example, in blends or highly filled samples, as the technique is more sensitive than
120
Delta Y=1.38%
Onset = 421.2
Percent weight loss (%)
100
80
60
40
Residual = 31.21%
Glass filled HDPE
20
Unfilled HDPE
0
Residual = 0.14%
0
100
200
300
400
Temperature in °C
500
600
700
Figure 6.7 TGA on a sample of pure HDPE and an overlay of a glass-filled (dashed) sample. (Reprinted
with permission from Veritas Testing and Consulting, LLC, Denton, Texas).
Thermal Analysis of Polyethylene 229
DSC. While both linear and volumetric expansions can be measured in the TMA, most
PE samples are measured for linear expansion. This can be done in the compression or
expansion as well as the extension or tensile geometry (see Figure 6.8 below).
Using TMA the size change in a sample is measured while under controlled conditions of temperature and atmosphere. While this normally shows an expansion with
temperature until it reaches the melt,3 under conditions where the material changes
by curing, burning out of a component, and others, shrinkage can occur. Again, temperature calibration is by melting point standards, and both the force and the probe
position are calibrated using reference materials. While traceable weights and standards like gauge blocks can be obtained, normally both are supplied by the instrument
manufacturer.
6.4.1 Coefficient of Thermal Expansion
The biggest application of TMA currently is the determination of the linear coefficient
of thermal expansion (CTE). Along with Cp, CTE is used, in engineering and designing
parts so that their expansion and contraction as temperatures change can be matched
to other components to prevent binding of parts, wrinkling of films, or leaking joints,
as well as for tuning processes for maximum efficiency. Measurement of the CTE of a
PE is shown in Figure 6.9. The thin line indicates the range used. Literature values range
in the 100–260 ppm/°C range [37]. A slow heating ramp of 2 to 5 °C/minute is applied
with as minimal a force on the probe as possible. In practice, this is 1–5 mN in compression and 5–10 mN in extension. Fixtures are normally made of quartz glass, which has
a CTE of 0.4 ppm/°C. (Note that in extension, stainless steel clamps are used to hold
the specimen and a baseline run must be made to be able to subtract their effect from
the data. This is particularly important in highly filled PE composites where the fillers
strongly affect expansion.)
Expansion
Penatration
Extension
Flexure
Dilatometry
Figure 6.8 Geometries and configurations used for TMA and DMA.
3
Or until the viscosity drops to the point below which it can’t support the probe. This is seen in polymers
like polystyrene that don’t have a melt but on heating in the TMA the glass transition will reach a point
where the probe drops. While this looks like a melt in the TMA scan, it is really a decrease in viscosity with
temperature to the point where the sample mass can no longer support the probe’s weight.
230
Handbook of Industrial Polyethylene and Technology
1.8
Proben position (mm)
1.75
1.7
1.65
1.6
1.55
CTE = 244.2 ppm/°C
1.5
1.45
1.4
–50
–30
–10
10
30
Temperature in °C
50
70
Figure 6.9 Coefficient of thermal expansion on a PE part. In this case, the CTE is calculated from
−40 to 70 °C.
6.4.2 Softening Point, Heat Distortion and Other Tests
There are several older tests that measure the softening point of a polymer by where it
bends when heated, where a probe tip will sink in and indent it, and others. Many of
these are still used because they predate the use of thermal analysis equipment, and
operating specifications have continued to refer to them. These tests have generally
been successfully done using TMA [38]. Most involve heating the sample through a
particular range and those interested are referred to reference [39].
6.5 Dynamic Mechanical Analysis (DMA)
Dynamic mechanical analysis is probably the second most important thermal technique
used for PE after DSC. By applying an oscillatory signal to a sample, DMA allows one to
obtain storage modulus, loss modulus and damping for a material across a wide range
of temperatures in a reasonable time scale. Changes in modulus and damping (also
called tan δ or loss factor) indicate the presence of the glass transition and the melting
temperature as well as other transitions that are undetected in DSC or TMA. A glass
transition in the DMA shows up as a decrease in the storage modulus as well as a peak
in both the loss modulus and tan d curves, as shown in Figure 6.10. (For information on
the melt rheology of PE, please refer to the chapter on rheology.) DMA is estimated to
be 100 times more sensitive to the presence of the glass transition than DSC and gives
modulus information as well. This makes it a very complimentary technique to DSC.
Because DMA instrumentation can be a complex topic, various designs and
approaches are used. Most common in the thermal analysis laboratory is the axial design
where force is applied along the z axis of the instruments [40]. Calibration remains an
issue as there is no primary standard for modulus or damping: manufacturers have
their own procedures to validate performance and supply reference materials. These
are normally well-characterized polymers but are not primary standards nor traceable
to any organization like the National Institute of Standards and Testing (NIST, USA) or
equivalent bodies. Temperature standards are the same melting standards as previously
Thermal Analysis of Polyethylene 231
0.3
Onset= –119.1°C
Storage modulus
0.25
1.00E+09
0.15
tan
Modulus in Pa
0.2
1.00E+08
0.1
Peak = –109.7°C
0.05
tan
1.00E+07
–150
–100
0
–50
Temperature in °C
50
100
0
Figure 6.10 DMA run on a sample of HDPE showing the sensitivity to transitions possible in the DMA.
The Tg is shown in the storage modulus and onset temperature at −119 to −110 °C. (Reprinted with
permission from Veritas Testing and Consulting, LLC, Denton, Texas)
discussed. Because deformation (or load) and frequency can affect the results greatly,
it is important that these values be recorded along with heating rates. For example, a
change in the testing frequency of one decade can shift a glass transition by 5 to 7 °C. In
addition, the method of measuring a transition is more variable than in DSC or TMA
and this method also needs to be defined [28]. Using the peak of the tan δ instead of the
onset of the modulus drop can raise the measured value for the Tg by as much as 25 °C
[8–12]. A Tg measured by DMA should have the geometry mode, the frequency, and
other conditions available, as these influence the value. For example, the Tg was −65 °C
by DMA (peak of tan delta, single cantilever, 5 °C/min heating rate, 1 Hertz). This gives
them the minimum needed to allow comparison to other methods.
6.5.1 Temperature Scans – Modulus and Transition Temperatures
The most common type of DMA run is that of a temperature scan where the material
is heated at 2 to 5 °C/min under inert gas (e.g., nitrogen) while being deformed by a
set amount or under a set load. This can be at either single or multiple frequencies. The
latter is often preferred and a set of 0.1, 1, and 10 hertz is often used to allow one to
characterize any frequency dependence in the material. The testing geometry or mode
is often single cantilever as it is very easy to load and samples can be cut from standard
ASTM “dogbones.” Films are run in tension and foams in compressions. Shear can be
used when an extremely wide range of temperature is needed. For PE, the runs are
often started at subambient temperatures, −150 °C, to start below the Tg and increase
to just below the melting temperature, unless the run is being done in shear. The shear
geometry allows testing into the melt. Often testing in the melt is done in torsional
rheometers [41, 42]. The glass transition is easily detected even in highly crystalline
linear PE, as shown by Figure 6.10. The larger modulus decrease and peak in tan d on
the right of the graph corresponds to the melt. The Tg is calculated by the onset of the
232
Handbook of Industrial Polyethylene and Technology
drop in storage modulus (solid) or the peak of tan d (dashed). Other methods are also
used in various industries [39].
Changes in the Tg temperature, the modulus value, and damping are monitored.
These changes are particularly indicative of changes in the crystallinity of the material,
of exposure to stress agents, or to thermal aging. Many modern instruments, like the
DMA, allow these tests to be done under exposure to UV light, humidity, and while
immersed in a solvent bath. Radiation sterilization, exposure to cleaning agents in the
manufacture, elevated temperatures, and other environmental stresses can cause serious degradation of properties in PE [43–47], and being able to test while exposed can
be advantageous.
Changes in modulus and damping make for a weaker, more brittle product that is
more susceptible to failure in use. Changes in the additive package used in the polymer
can also induce changes. Figure 6.11 shows the effect of changing the amount of nucleation agent on the DMA data [48]. Similarly one would see changes in the Hf values
for the samples in the DSC.
6.5.2 Frequency and Other Scans
Storage modulus, Pa
Most DMA are capable of other tests like frequency scans at isothermal temperatures,
creep-recovery, stress-relaxation, and limited stress-strain scans. Of the other testing
modes available, the most useful is the frequency test. Testing PE products like pipes
and geotextiles [49] is often done by frequency scans, where the material is held isothermal while the frequency of oscillation is varied and then heated to another temperature for similar testing. The process then repeats [8–12]. The data collected from this
can be used in time-temperature superposition to generate master-curves, predicting
processes over a much wider frequency rate or longer time than actually run [50, 51].
Once the frequency scans are collected and the master curve generated, the shift
factors and Cole-Cole plot are used to check on the validity of the results [39]. Once
one is reasonably sure that the data is valid, the master curve can be used to predict
testing at frequencies unattainable in the DMA, such as very fast frequencies like 1010
hertz or those so slow as to be unmeasurable in reasonable times. As frequency can
be expressed as reciprocal seconds, the curve can be inverted to estimate time. Other
Sample 3
1.000E+09
Sample 2
1.000E+08
Increasing levels
of nucleation agents
Sample 1
1.000E+07
–20
0
20
40
60
Temperature, °C
80
100
Figure 6.11 DMA runs on a series of PE samples with different nucleation agents. Modulus increases as
the amount of crystallinity in the sample increases. Data from unpublished work [40].
Thermal Analysis of Polyethylene 233
uses of the master curve include estimating chain branching, molecular distribution,
and molecular weight. All of these are beyond the scope of this chapter and require a
fair amount of sophistication [50, 52]. The reader is referred to Wrana’s text [53] for a
complete discussion.
While most DMA instruments are also capable of tests like stress-strain, creep recovery, and stress-relaxation measurement, these are beyond the scope of this discussion,
and the reader is referred to the appropriate reference [39].
6.6 Coupled Thermal Techniques
Coupled or hyphenated thermal techniques [54–56] are terms used to described the
connecting of two instruments together to analyze the same sample or the by-products
of the first analysis. The approach has been used for the detection of evolved gases since
the 1960s when the first mass spectrometer (MS) was coupled to a TGA [57]. Since
then, thermal instruments have been coupled to a wide range of other instruments to
allow measurement of everything from the metal content in a polymer to its crystallization rate, to the off-gases. Of this range of options, two techniques have applications to
PE: spectral DSC and evolved gas analysis.
6.6.1 Spectral DSC
The coupling of DSC to either a Raman spectrometer or to a near-infrared (NIR) spectrometer allows one a greater understanding of the behavior of a material on crystallization. Raman spectroscopy particularly has shown the ability to detect changes in a PE
sample [58, 59], as is shown in Figure 6.11. Here a good (flexible) and bad (brittle) sample of PE film were isothermally crystallized in the DSC while monitored by Raman.
The figure on the left shows the ICHT scans on the two PE samples. Note the faster
crystallization and higher DHf of the “bad” material. The figure on the right shows the
–1.221
–2
0.203
Good material
Peak = 1.200 min
Area = –44.962 mJ
Delta H = –8.9923 J/g
–6
–8
Bad material
–20
Peak = 0.733 min
Area = –71.002 mJ
Delta H = –14.2003J/g
–12
–14
–16
–18
–20
–21.92
0.0
0.5
1.0
1.5
Time, min
0.18
0.16
0.14
0.12
0.10
0.08
0.06
0.04
0.02
0.00
–0.02
–0.04
–0.06
–0.08
–0.093
Bad
material
Bad Material
Good
material
Good Material
PC1PC(75%)
1 (75%)
Heat flow endo up, mW
–4
2.0
2.5 2.69
10
50
100
Index
Index
150
184.7
Figure 6.12 DSC (left) and Raman data (right) from a spectral DSC experiment tracking the isothermal
crystallization of a PE sample. Endotherms are up in this figure. (Reprinted with permission from
PerkinElmer, Inc., Shelton, CT)
Real time
Destructive on vapor
Qualitative
Extremely sensitive
Same limits as GCMS
Extremely sensitive
Very Sensitive
and FTIR methods
Columns can be 02 or halo- However limited by chosen
Complex due to overlapping
type of column (not
fragmentation patterns
gen sensitive
Good for exploratory work
good for unknowns)
AMU range and
High mass wash-over concerns
Non-destructive on vapor
Qualitative
IR sensitivity limited
Mixtures can be difficult
CO2 and H20 problematic
Complex to operate
Analysis tricky
Easy to operate
Analysis straightforward
Easy to operate
Analysis straightforward
Libraries available
Pyrolysis libraries useful
Libraries available
Libraries available
Specific EGA libraries now
available
Quantitative
Destructive on vapor
Requires long times following TGA experiment
Time depends on number
of samples collected
Complex to operate
Analysis varies
Detector dependent
Traditional GC techniques
often required.
Quantitative
Destructive on vapor
Near real time
(uGC runs ~ 2–3 mins)
Very complex to operate
Complex analysis
Need both GCMS and
FTIR libraries
Quantitative
Destructive on vapor
Real Time for IR
GCMS analysis takes
longer
Resolves overlapping
evens while restoring
time info
Real time
Requires knowledge of
what will be evolved.
Wide range of column
types and detectors
available
Detection of specific materials quickly
Depends on column type
Small molecules and
fragments
but can range from small
Larger molecules and mixtures
to large.
get complex
Good for separating
mixtures
Resolve overlapping events.
Large molecules.
Symmetric molecules
invisible
TG-IR-GCMS
Mass ios and fragment
analysis.
Isotope detection
TG-uGC
Functional group analysis
TG-GCMS
TG-MS
TG-IR
Table 6.3 Comparison of various EGA techniques for polymers.
234
Handbook of Industrial Polyethylene and Technology
Thermal Analysis of Polyethylene 235
Mix of products
HCL bands
TGA of PE
elastomer
SO2
SO2
Abs
0.25
PE combustion products
0.00
10
Wa
ve
nu
mb
er
500
Mi
ns
0
Figure 6.13 TG-IR data for a chlorinated PE elastomer that had also been sulfonated. (Reprinted with
permission from Veritas Testing and Consulting, LLC, Denton, TX)
principal component analysis of the Raman spectra corresponding to the formation of
crystals in PE. One sees, in the Raman data, how the crystal phases forms slower and to
a less extent in the “good” or more flexible sample.
6.6.2 Evolved Gas Analysis (EGA)
Evolved gas analysis (EGA) couples the off-gas from a TGA or STA to a Fourier transform infrared spectrometer (FTIR), a mass spectrometer (MS), a gas chromatographmass spectrometer (GCMS), or some combination of these. These techniques are often
abbreviated as TG-IR, TG-MS, and TG-GCMS. Each technique has advantages and disadvantages, as shown in Table 6.3. The most common technique used is TG-IR, because
of the traditional association of thermal and infrared techniques in polymer labs.
Coupled techniques [60] are most commonly used in two applications for PE and
its copolymers or derivatives. The first is to quantify what is emitted when the polymer burns. With simple PE of any grade, it is a pretty clean burn [61]. However, when
chlorinated PE elastomers or sulfonated PE are burned, a much more complex burn
profile results. Figure 6.13 shows the results of a TG-IR study on a sulfonated chlorinated PE elastomer; note the strong signals from both the HCl and SO2 formed from
the combustion of the sample.
The second application is the detection of leachates or contaminates from the material enclosed in a PE container into the polymer. The leachates can degrade polymer
performance by increasing brittleness, degrading antioxidants, or decoloring the material. By running a TG-GCMS experiment, one can concentrate low-level contaminates
on the column and then obtain a GCMS of a specific weight loss.
6.7 Conclusions
Thermal analysis techniques continue to be an important tool in the industrial use of
PE and its copolymers. As the amount of PE recycled increases, DSC has become more
236
Handbook of Industrial Polyethylene and Technology
important in that area due to FTIR’s inability to separate grades of PE. The complex
blending of materials for specific properties makes traditional methods less dependable. Coupled techniques expand the ability of thermal methods to solve material and
processing problems.
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Technique: Application to the Glass Transition of Polymers. J. Therm. Anal. Calorim., 87,
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tpl.fpv.ukf.sk/engl_vers/thermophys/proceedings/simon.pdf.
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21. Gaur, U., and Wunderlich, B., Heat Capacity and Other Thermodynamic Properties of
Linear Macromolecules. II. Polyethylene, J. Phys. Chem., 10, 119, 1981.
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23. Androsch, R., Di Lorenzo, M.L., Schick, C., and Wunderlich, B., Mesophases in Polyethylene,
Polypropylene, and Poly(1-butene), Polymer, 51, 4639, 2010.
24. Schick, C., personal communication, 2015.
25. Jones, J., Barenberg, S., and Geil, P., Amorphous Linear Polyethylene: Electron Diffraction,
Morphology, and Thermal Analysis, J. Macromol. Sci., Phys., B15, 329, 1977.
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Polym. Bull., 1, 127, 1978.
27. Icenogle, R.D., and Chatterjee, A.M., Effect of Nucleating Agents on the TemperatureDependent Melt Crystallization Kinetics of Poly(1-butene), J. Appl. Polym. Sci., 31, 1859,
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28. Ehrenstein, G.W., Riedel, G., and Trawiel, P., Thermal Analysis of Plastics: Theory and
Practice, Hanser: Munich, 2004.
29. Wright, D., Failure of Plastic and Rubber Products, RAPRA Technology Ltd: UK, 2001.
30. Hsuan, G., and Guan, Z., Evaluation of the Oxidation Behavior of Polyethylene
Geomembranes Using Oxidative Induction Time Tests, in: Oxidative Behavior of Materials
by Thermal Analytical Methods, ASTM STP 1326, Riga, A., and Patterson, G. (Eds.), 76,
ASTM: West Conshohocken, PA, 1997.
31. Ashby, G.E., Oxyluminescence from Polypropylene, J. Polym. Sci., 50, 99, 1961.
32. Schard, M.P., and Russel, C.A., Oxyluminescence of Polymers. I. General Behavior of
Polymers, J. Appl. Polym. Sci., 8, 985, 1964.
33. Jipa, S., Zaharescu, T., Setnescu, R., Setnescu, T., Brites, M.J.S., Silva, A.M.G., MarceloCurto, M.J., and Gigante, B., Chemiluminescence Study of Thermal and Photostability of
Polyethylene, Polym. Int., 48, 414, 1999.
34. Catalina, F., Peinado, C., Allen, N.S., and Corrales, T., Chemiluminescence of Polyethylene:
The Comparative Antioxidant Effectiveness of Phenolic Stabilizers in Low-Density
Polyethylene, J. Polym. Sci. Part A Polym. Chem., 40, 3312, 2002.
35. Wunderlich, B., Thermal Analysis of Polymeric Materials, Springer: New York City, NY,
2005.
36. Kayacan, I., and Dogan, O.M., Pyrolysis of Low and High Density Polyethylene. Part I:
Non-isothermal Pyrolysis Kinetics, Energy Sources, Part A, 30, 385, 2008.
37. Lampman, S. (Ed.), Characterization and Failure Analysis of Plastics, ASM International:
Materials Park, OH, 2003.
38. Riga, A., and Neag, C.M. (Eds.), Materials Characterization by Thermomechanical Analysis,
ASTM: Philadelphia, PA, 1991.
39. Menard, K., Dynamic Mechanical Analysis: A Practical Introduction, 2nd ed., CRC Press:
NYC, 2008.
40. Menard, K., and Menard, N., Dynamic Mechanical Analysis in the Analysis of Polymers
and Rubbers, in: Encyclopedia of Polymer Science and Technology, 1–33, Wiley: New York
City, NY, 2015.
41. Rohn, C.L., Analytical Polymer Rheology, Hanser Gardiner Publications: Cincinnati, OH,
1995.
42. Sousa, R., Mater. Sci. Appl., 5, 923, 2014.
43. Sardashti, P., Scott, A.J., Tzoganakis, C., Polak, M.A., and Penlidis, A., Effect of Temperature
on Environmental Stress Cracking Resistance and Crystal Structure of Polyethylene,
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Handbook of Industrial Polyethylene and Technology
44. Sardashti, P., Tzoganakis, C., Zatloukal, M., Polak, M.A., and Penlidis, A., Rheological
Indicators for Environmental Stress Cracking Resistance of Polyethylene, Int. Polym. Proc.,
30, 70, 2015.
45. Sardashti, P., Tzoganakis, C., Polak, M.A., and Penlidis, A., Radiation Induced Long Chain
Branching in High Density Polyethylene through a Reactive Extrusion Process, Macromol.
React. Eng., 8, 100, 2014.
46. Kissin, Y.V., Polyethylene: End-Use Properties and their Physical Meaning, Hanser Publishers:
Munich, Germany, 2012.
47. Wright, D., Failure of Plastics and Rubber Products, Rapra Technology: Shopshire, 2009.
48. Brostow, W., and Menard, B., personal communication, 2016.
49. Hsuan, Y., Koerner, R., and Lord, A., Notched Constant Tensile Load (NCTL) Test for HighDensity Polyethylene Geomembranes, Geotech. Test. J., 16, 6, 1993.
50. Dealy, J., and Plazek, D., Time-Temperature Superposition – A Users Guide, Rheology
Bulletin, 78, 16, 2009.
51. Razarvi-Nouri, M., and Hay, J., Iran. Polym. J., 13, 360, 2004.
52. Plazek, D., J. Rheology, 40, 28, 1996.
53. Wrana, C., Introduction to Polymer Physics (English Edition), Lanxess: Leverkuessen, 2015.
54. Proveder, T., Urban, M.W., and Barth, H.G. (Ed.), Hyphenated Techniques in Polymer
Characterization, ACS Symposium Series 581, ACS: Washington DC, 1994.
55. McClennen, W.H., Buchanan, R.M., Arnold, N.S., Dworzanski, J.P., and Meuzelaar, H.L.C.,
Thermogravimetry/Gas Chromatography/Mass Spectrometry and Thermogravimetry/Gas
Chromatography/Fourier Transform Infrared Spectroscopy: Novel Hyphenated Methods
In Thermal Analysis. Anal. Chem., 65, 2819, 1993.
56. Menard, K., and Waters, J.L., in: Symposium, Pittsburgh Conference Proceedings, Waters
Lecture 3, 2015.
57. Schwenker Jr., P., and Garn, P., Thermal Analysis, Academic Press: New York City, NY, 1969.
58. Menard, K., in: Proceedings of the PE TopCon, SPE: Brookfield, 29, 2009.
59. Menard, K., and Spragg, R., Combining Raman Spectroscopy and Differential Scanning
Calorimetry, Spectroscopy, 17(6), 62, 2009. http://spectroscopyonline.findanalytichem.
com/spectroscopy/Raman+Spectroscopy/Combining-Raman-Spectroscopy-andDifferential-Scan/ArticleStandard/Article/detail/600239.
60. Litvinov, V., Spectroscopy of Rubbers and Rubbery Materials, p. 14, RAPRA: Shopshire, 2002.
61. Serio, M.A., Bassilakis, R., and Solomon, P.R., Use of TGFTIR Analysis for the
Characterization of Fuels and Resources, in: ACS Preprints, New Orleans, 1994. https://web.
anl.gov/PCS/acsfuel/preprint%20archive/Files/41_1_NEW%20ORLEANS_03–96_0043.
pdf.
7
Rheology of Polyethylene
Gregory W. Kamykowski
TA Instruments – Waters LLC, Wood Dale, Illinois, USA
Contents
7.1 Rheology Fundamentals ........................................................................................................... 240
7.1.1 Flow Testing ................................................................................................................... 242
7.1.2 Deformation Testing ..................................................................................................... 245
7.1.3 Dynamic Testing: Fundamentals, Dynamic Strain Sweeps, Frequency Sweeps,
Dynamic Temperature Ramps ..................................................................................... 246
7.1.3.1 Fundamentals .................................................................................................. 246
7.1.3.2 Dynamic Strain Sweeps .................................................................................. 249
7.1.3.3 Dynamic Frequency Sweeps .......................................................................... 250
7.1.3.4 Dynamic Temperature Ramps ...................................................................... 250
7.2 Melt Rheology............................................................................................................................ 251
7.2.1 Extrusion Plastometer .................................................................................................. 252
7.2.2 Rotational Rheometry .................................................................................................. 254
7.2.2.1 Polymer Flow ................................................................................................... 254
7.2.2.2 Dynamic Frequency Sweeps .......................................................................... 257
7.2.2.3 Dynamic Frequency Sweeps, Time Temperature Superposition.............. 260
7.2.2.4 Creep Testing ................................................................................................... 266
7.2.2.5 Creep, Time Temperature Superposition .................................................... 267
7.2.2.6 Extensional Testing ......................................................................................... 269
7.2.3 Capillary Rheometry..................................................................................................... 270
7.2.4 Time Temperature Superposition with Capillary Data ............................................ 273
7.3 Dynamic Mechanical Testing on Solids and Solid-like Materials ...................................... 273
7.3.1 Dynamic Testing with Rotational Deformation ........................................................ 273
7.3.2 Dynamic Testing in Linear Deformation ................................................................... 275
7.3.3 Dynamic Temperature Ramps ..................................................................................... 276
7.3.4 Other Tests on a DMA .................................................................................................. 279
7.4 Conclusions................................................................................................................................ 280
References ............................................................................................................................................ 280
Corresponding author: gkamykowski@tainstruments.com
Mark A. Spalding and Ananda M. Chatterjee (eds.) Handbook of Industrial Polyethylene and Technology, (239–282)
© 2018 Scrivener Publishing LLC
239
240
Handbook of Industrial Polyethylene and Technology
Abstract
Rheological testing is a practical and sensitive way of characterizing polymeric materials,
including polyethylene (PE) in its molten and solid forms. Rheological characterizations reveal
information about the molecular weight, molecular weight distribution, and branching information of PE resins. From these characterizations, the behavior of a given material, in terms of
its processability, and end-use performance can be determined.
In this chapter, rheology fundamentals and their applicability to PE are explained. We also
discuss the rheological methods that are commonly used by PE producers and processors and
show how these relate the molecular architecture to ultimate performance. These methods
include flow, dynamic oscillatory testing, and transient testing like creep. Commonly used geometric configurations, such as capillary geometries and rotational geometries, are shown.
Keywords: Rheology, creep, dynamic testing, capillary testing, viscosity, viscoelasticity,
processability
7.1 Rheology Fundamentals
Rheology is the study of the flow and deformation of matter. The word was coined by
Professor Eugene C. Bingham of Lafayette College in 1920. When we examine the flow
properties of matter, we are describing the behavior of liquids and gases. We illustrate
this behavior below in Figure 7.1 with a dashpot on the left and examples that apply to
commercial rheometers on the right. The key feature of fluid behavior illustrated with
the dashpot is that the Force (F) is dependent on the velocity (v) and is independent of
the position (x) in the dashpot. The corollary for this would be that, as an example, for a
parallel plate geometry on a rotational rheometer, the torque is dependent on the angular
velocity and independent of the angular displacement. For the capillary setup (far right),
the force depends on the rate and is independent of the vertical position in the barrel.
For deformation properties, we are describing the behavior of solid-like materials.
The object that illustrates this behavior appropriately is a spring, where the force is
dependent on the position and is independent of the velocity, just the opposite of the
fluid described above. The deformation system is shown by Figure 7.2.
Most industrial polymers are not just simple fluids or simple solids, although viscous
properties may dominate when the materials are in the molten state and elastic properties may dominate when the plastic part has been made and used for its intended
purpose. We use the term viscoelastic to describe the behavior shown by most plastic
materials. The term, viscoelastic, means that the materials exhibit both fluid and solidlike behavior.
F
Flow: Fluid behavior;
Viscous nature
F = F(v); F F(x)
Figure 7.1 Illustration of fluid behavior. The dashpot is shown on the left. Examples of a parallel plate
and capillary rheometer geometry are shown on the right.
Rheology of Polyethylene 241
F
Deformation: Solid behavior
Elastic nature
F = F(x); F F(v)
0 1 2 3
X
Figure 7.2 Illustration of solid behavior. The spring is shown on the left. Examples of torsional and
tensile testing of a solid material are shown on the right.
F
Maxwell model
F
Kelvin (Voigt) model
Figure 7.3 Illustration of viscoelasticity by the Maxwell model (left) and the Kelvin model (right).
A common way to illustrate the viscoelasticity is by combining the spring and the
dashpot models used above. If the spring and dashpot are combined in series, we refer
to this as the Maxwell model. If we combine the spring and dashpot in parallel, we refer
to this as the Kelvin or Voigt model. These models are shown in Figure 7.3.
The model to select is often based on the type of test that is being performed. For
example, if a stress relaxation test is performed, an instantaneous displacement of the
free end of a viscoelastic element is applied, and the more appropriate model is the
Maxwell model. The behavior would be predicted by Equation 7.1.
G(t ) G0e
t
t0
(7.1)
where G(t) is the modulus (i.e., stress/strain) at time t; G0 is the spring constant, and t0
is the ratio of the dashpot fluid viscosity to the spring constant. In this case the material
is not cross-linked.
For creep tests, an instantaneous force is imposed and the displacement with time
is monitored, just the opposite of a stress relaxation test, and the Kelvin model would
be the more appropriate model. In this case the behavior would be described by
Equation 7.2.
J (t )
J0
Je 1 e
t
t0
t
(7.2)
0
where J(t) is the compliance (strain/stress) at time t, J0 is an instantaneous compliance,
Je is the equilibrium recoverable compliance, t0 is the ratio of the dashpot fluid viscosity
to the spring constant, and 0 is the steady-state viscosity.
When analysts use these models to fit the data from rheological testing, it is common to need several spring and dashpot elements to describe adequately the behavior
observed.
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Handbook of Industrial Polyethylene and Technology
7.1.1 Flow Testing
In the introductory section above, we have already used a few rheological terms that
are part of the rheologist’s lexicon, such as viscosity, modulus, and compliance. At this
time, it is useful to mention and describe the terms that are used most frequently in
rheology and will be used throughout the rest of this section of the book.
We begin with examining the simple flow pattern for steady laminar flow, as shown
in Figure 7.4. In this system, we have a fluid between two parallel plates. We put the
upper plate in motion with a velocity V0. The bottom plate remains stationary. If a vertical stripe with a dye was placed on the outer edge of the system prior to flow initiation
and we observed the flow pattern, it would look like the schematic shown in Figure 7.4.
The flow behavior can be described by the relationship shown in Figure 7.4:
vx
V0
y
H
(7.3)
Note that the velocity of a fluid particle in contact with the upper moving plate is V0,
and the velocity of a particle in contact with the lower stationary plate is 0. Essentially,
we are saying that these particles are not slipping from their respective plates.
We can now define the fundamental flow parameter in rheology and viscometry,
which is the shear rate, . The shear rate is the change in the x-velocity with respect to
y; i.e.,
dv x /dy. By using the shear rate as the flow parameter, we can, in theory, correlate flow behavior in a multitude of flow configurations, such as parallel plate or coneand-plate rotational flow, concentric cylinder flow, and capillary flow. Parallel plate and
cone-and-plate geometries are used in rotational rheometers; concentric cylinder flow
is approximated in extruders; and capillary flow is employed in capillary rheometers
and extrusion plastometers. The unit for shear rate is 1/seconds.
Just as we normalized flow by utilizing a fundamental flow parameter like the shear
rate, we also want to normalize the force to take into account the type of geometry and
the dimensions of the system. In the laminar flow system shown by Figure 7.4, this is
done simply by dividing the force by the area of the plate in contact with the fluid. There
are appropriate formulas for different geometries. These formulas will not be derived
Steady simple shear flow
Top plate velocity = V0; Area = A; Force = F
H
Bottom plate
Velocity = 0
y
x
vx = (y/H)*V0
.
= dvx/dy = V0/H
= F/A
= ./ .
Figure 7.4 Laminar flow of a fluid.
Shear rate, sec–1
Shear stress, Pascals
Viscosity, Pa-sec
Rheology of Polyethylene 243
here. The reader can go to references for thorough treatments on the derivations of
these formulas for common geometric configurations [1, 2].
For shear flow, the most commonly used symbol for shear stress is . However, many
publications and books have used for shear stress as well. It will be clear from the
context what the symbol is being used for. The units of stress are Pa in the SI system,
and dynes/cm2 in the cgs system. SI units are the preferred units. 1 Pa = 10 dynes/cm2.
Finally, the ratio of the shear stress to the shear rate is the shear viscosity. The symbol
used in plastics rheology is . The SI units are Pa∙s, and the cgs units are Poise. 1 Pa ∙ s =
10 Poise. It is this parameter that is used most often to show the overall flow behavior
of a molten plastic.
The most common way of displaying a molten thermoplastic’s flow behavior is by
plotting the viscosity as a function of shear rate, as shown by Figure 7.5. Both viscosity
and shear rate are plotted on logarithmic axes.
Figure 7.5 identifies certain flow regimes. In the low shear rate region, a neat polymer melt has a viscosity that is independent of shear rate. This is the Newtonian region,
and the viscosity in this area is called the Newtonian viscosity or the zero-shear viscosity 0. The zero-shear rate viscosity is strongly dependent on the molecular weight
of the plastic. The relationship is expressed in Equation 7.4. The actual exponent may
range from 3.2 to 3.9, but the most common exponent in this relationship is 3.4 [3].
So, if the molecular weight of one material is 10% higher than another, its zero-shear
viscosity will be 40% higher. Clearly the viscosity measurements are very sensitive to
changes in molecular weight. Note that the weight average molecular weight (Mw) is
used in this relationship.
0
K Mw
3. 4
(7.4)
At higher shear rates, the viscosity decreases. This decrease in viscosity with increasing shear rate is known as shear thinning behavior. The shear rate where the resin
exhibits non-Newtonian behavior is dependent on its molecular weight and molecular
1.E+05
Pseudoplastic
behavior
Newtonian
region
Viscosity (Pa-sec)
1.E+04
0 < n < 1 (Newtonian)
1.E+03
=
Transition
region
0
Power law
region
1.E+02
1.E+01
.
= m* (n–1)
1.E+00
1.E–02
0
MW
3.4
1.E–01
1.E+00
1.E+01
1.E+02
Shear rate (sec–1)
Figure 7.5 Representative flow curve of a thermoplastic melt.
1.E+03
1.E+04
1.E+05
244
Handbook of Industrial Polyethylene and Technology
weight distribution. For many resins, a region exists where the log viscosity versus log
shear rate plot becomes a straight line. Since this is a log-log plot, the relationship in
this region can be expressed as Equation 7.5.
m (n 1)
(7.5)
where m is a fitting constant and (n-1) is the slope of the line in this region, which we
call the power law region. This is an important region for processing polymer melts.
The relationship is alternately expressed in terms of the shear stress, :
m n
(7.6)
The typical range for n is 0.15 to 0.60 [4]. This is sometimes referred to as the power
law model of Ostwald and de Waele. The fact that molten polymers exhibit shear thinning behavior is beneficial from a processing standpoint in industrial applications.
For example, a low viscosity in the high shear rate region will facilitate processing by
decreasing the discharge temperature, a desirable trait for extrusion processing resin
for blow molding operations. When the parison for blow molding is formed, having a
high viscosity at a low shear rate will minimize parison sag before the mold can clamp
onto the parison.
For materials of a similar family and similar molecular weight distribution, the effect
of molecular weight can be seen in Figure 7.6. Note that the differences are the greatest
at the lowest shear rates and that the flow curves are converging at higher shear rates.
One thing to keep in mind is that resins encounter a wide range of shear rates as they
are being processed. Just knowing the viscosity at a low shear rate, an intermediate
shear rate, or a high shear rate is not enough to predict whether a resin will perform
satisfactorily in a given application.
The relationship between instrument parameters and rheological parameters will
be presented later. For rotational rheometers, this will include relating angular velocity to shear rate, angular displacement to strain, and torque to stress. For capillary
1.E+05
Viscosity (Pa-sec)
Low
Medium
1.E+04
High
1.E+03
1.E+02
1.E+01
1.E–02 1.E–01 1.E+00 1.E+01 1.E+02 1.E+03 1.E+04
Shear rate (1/sec)
Figure 7.6 Viscosity versus shear rate for a given material with different molecular weight values.
Rheology of Polyethylene 245
rheometers, cross-head speed will be related to the shear rate, and the force will be
related to the stress.
7.1.2 Deformation Testing
This section will present transient testing on viscoelastic polymers. Instead of flow testing, the focus is primarily on the deformation properties of the material. The starting
point for this testing is similar to what was done for flow testing, except that we focus
on deformation. The simplest deformation pattern is shown in Figure 7.7.
This is identical to what we showed for flow testing, except that now we have an
instantaneous deformation, not steady flow. The appropriate fundamental deformation
parameter is the strain, which is the change in x-deformation with respect to y. The
symbol for strain is , which is unitless. The shear stress, , is the same as it was before,
and the ratio of stress to strain, / , is the modulus G. The units of modulus are Pa in
the SI system and dynes/cm2 in the cgs system.
The test described by this approach is called the stress relaxation test. Here, a strain
is applied and held constant and then the modulus is monitored. The typical behavior
of a polymer melt is shown in Figure 7.8.
Simple shear deformation
Top plate displacement = X0; Area = A; Force = F
H
y
Bottom plate
Displacement = 0
x
x = (y/H)*X0
= dxx/dy = X0/H
Shear strain, unitless
= F/A
G= /
Modulus, Pa
Shear stress, pascals
Figure 7.7 Simple shear deformation.
106
100.00
105
10.00
104
1.00
103
0.10
102
0.01
0.10
1.00
Step time ts (s)
Figure 7.8 Stress relaxation of a polymer melt.
10.00
0.01
100.00
Strain (%)
Modulus G(t) (Pa)
0.1 MI HDPE 190 C
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Handbook of Industrial Polyethylene and Technology
6×10–3
1000
0.1 MI HDPE
4×10–3
100
3×10–3
Stress
(Pa)
Compliance J(t) (1/Pa)
5×10–3
2×10–3
10
1×10–3
0×100
0
1000
2000
3000
1
5000
4000
Step time ts (s)
Figure 7.9 Creep testing of a polymer melt.
Table 7.1 Summary of the properties presented in this section. All of the testing is
unidirectional.
Rheological property
Symbol
Units
Shear Rate
1/s
Shear Stress
Pa
Viscosity
/
Strain
Pa∙s
–
Modulus
G(t) = (t)/
Pa
Compliance
J(t) = (t)/
1/Pa
The other transient test is called the creep test. It is the opposite of the stress relaxation test in that, in a creep test, a stress, , is applied instantaneously, and the strain,
(t), is monitored as a function of time. We define the compliance, J(t), as the ratio of
the strain, (t), to the stress. The units of compliance are 1/Pa. An example of a creep
test is shown in Figure 7.9.
The terms and rheological properties presented in this section are summarized in
Table 7.1
7.1.3 Dynamic Testing: Fundamentals, Dynamic Strain Sweeps,
Frequency Sweeps, Dynamic Temperature Ramps
7.1.3.1
Fundamentals
The test descriptions so far have focused on unidirectional testing. However, with commercially available rotational rheometers and dynamic mechanical analyzers, a very
important type of testing that is often performed is dynamic oscillatory testing, whereby
Rheology of Polyethylene 247
Strain
Stress
Time
Figure 7.10 Dynamic oscillation testing on an elastic material.
the test specimen is subjected to sinusoidal strains or stresses. In the majority of cases,
dynamic oscillatory testing is the preferred way of measuring viscoelastic properties.
It is usually more convenient than stress relaxation or creep testing.
To begin the discussion on dynamic testing, consider a sinusoidal application of
strain. A sinusoidal application of stress could also be used, but the key points will be
presented by just considering sinusoidal strain imposition. The strain is described by
Equation 7.7.
(t )
0
sin( t )
(7.7)
Note that the strain history is described by 2 parameters, the strain amplitude 0, and
the frequency , which is in radians/second.
The first situation to consider is that of a perfectly elastic material, like a spring.
As discussed in Section 7.1.1, the stress is directly proportional to the strain in this
situation
(t ) k (t ) k 0 sin( t )
0
sin( t )
(7.8)
A plot of stress and strain versus time is shown in Figure 7.10. Note that the stress
and strain are perfectly in phase with each other. The peaks and the valleys occur at the
same time.
The second situation is that of a viscous fluid. Again, going back to Section 7.1.1, the
stress would be directly proportional to the shear rate .
(t )
A (t )
A 0 cos( t )
(7.9)
where A is the viscosity. Note that, since the stress is dependent on the cosine, it is 90° or
/2 radians out of phase with the strain. For a viscous fluid, the plots of stress and strain
would look like Figure 7.11. Note that the peaks and valleys are out of phase.
For most of the polymeric materials that are tested by dynamic testing, the phase
shift is between the limits of 0° for a perfectly elastic solid and 90° or /2 radians for a
perfectly viscous fluid and has a value of , as shown by Figure 7.12. The lower the value
for , the greater the elasticity and the less the viscous nature of the material.
For dynamic oscillatory testing, we can summarize our discussion and define some
of the key parameters that are used in dynamic oscillatory testing, namely, the storage
modulus (G , E ), the loss modulus (G , E ), tan , the complex modulus (G*, E*), and
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Handbook of Industrial Polyethylene and Technology
Strain
Stress
Time
This time corresponds
to /2 radians.
Figure 7.11 Dynamic oscillation testing on a viscous material.
Strain
Stress
This time corresponds
to radians.
Figure 7.12 Response of a viscoelastic material to a sinusoidal strain.
Table 7.2 Summary of dynamic oscillatory testing and parameters frequently used in dynamic
testing.
Parameter
Shear
Elongation
Units
Strain
= 0sin( t)
= 0sin( t)
–
Stress
=
= 0sin( t + )
Pa
0
sin( t
)
Storage Modulus
(Elasticity)
G = ( 0/ 0)cos
E = ( 0/ 0)cos
Pa
Loss Modulus
(Viscous Nature)
G = ( 0/ 0)sin
E = ( 0/ 0)sin
Pa
Tan
G /G
E /E
–
Complex Modulus
G* = (G 2
Complex Viscosity
* = G*/
G 2)0.5
E* = (E 2
* = E*/
E
E 2)0.5
Pa
Pa-sec
the complex viscosity ( *). Note that dynamic testing can be performed in shear deformation as well as in elongational, flexural, and compression modes. Shear, tension, and
compression are the three fundamental modes of deformation. Flexural deformation is a
composite of tension and compression. The descriptions and the parameters are shown
in Table 7.2. An analogous description for the elongational parameters is also provided.
Rheology of Polyethylene 249
104
LDPE 190
105
Storage modulus ‡ (Pa)
Loss modulus ^ (Pa)
103
102
101
Oscillation stress E (Pa)
104
100
103
10–2
10–1
100
101
Oscillation strain R (%)
102
10–1
103
Figure 7.13 Strain sweep of nominal 10 dg/min MI LDPE at 190 °C and a frequency of 1 Hz.
7.1.3.2
Dynamic Strain Sweeps
The first test that is usually done when performing dynamic testing is a dynamic strain
or stress sweep. This is typically done at a frequency of 1 Hz (6.28 rad/s), but some
operators use a frequency of 10 rad/s (1.59 Hz) or 10 Hz.
The goal is to determine what is called the linear viscoelastic region (LVR) of the
material. A graph of a strain sweep for a nominal 10 dg/min MI low density PE (LDPE)
resin (190 °C, 2.16 kg) is shown in Figure 7.13. The testing on the LDPE was performed
at 190 °C and at a frequency of 1 Hz. A graph like that shown in Figure 7.13 is useful by
itself, but the information gained from this type of graph is often used for subsequent
frequency sweeps, temperature sweeps, and temperature ramps.
The area of interest is where the G and G are independent of strain. From Figure 7.13,
it appears that LDPE is in the LVR until a strain of about 37%. Beyond this point, the
moduli are not constant, and a breakdown of the material, slippage, edge fracture, or
other phenomenon has resulted in the material showing changes in the values of G and
G . The LVR region is called this because the plot of the oscillation stress versus the
oscillation strain is a straight line in this region. For Figure 7.13, the graph of oscillation
stress versus oscillation strain will have a slope of 1 in the LVR since both parameters
are plotted on a logarithmic scale.
The LVR is typically the region where subsequent testing is performed such as frequency sweeps and temperature ramps. This is where the reproducibility is the best and
where the most fundamental properties of the material are obtained. In recent years,
more attention has been given to the nonlinear viscoelastic region in dynamic testing,
but this should always be done in addition to, not in place of, testing in the LVR. Testing
in the nonlinear viscoelastic range is referred to as large amplitude oscillatory shear
(LAOS). The term that is often used with LAOS testing is the “fingerprinting” of a material. This is a developing area in the field of rheology, and there are many papers coming
out that deal with this approach, including those that apply to polymer melts [5, 6].
Of the two dynamic parameters (G , G ), G is the more sensitive. In the molten state,
it is common for neat materials to exhibit linear behavior to strains of 50% or more.
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Handbook of Industrial Polyethylene and Technology
The maximum strain in the LVR will depend on the molecular weight of the material,
the temperature, and the frequency. Properties that make the material more solid-like
(higher molecular weight, lower temperature, and higher frequency) reduce the strain
at which the material enters the nonlinear viscoelastic region. The material is considered to be out of the LVR when its storage modulus is 5% lower than its value in the
LVR, although some researchers prefer to use a decrease of 10% as their criterion.
7.1.3.3 Dynamic Frequency Sweeps
For polymer melts, the most common oscillatory test performed is the dynamic frequency sweep from 100 to 0.1 rad/s. This is illustrated in Figure 7.14. There are some
salient features that can be seen in this figure. First, at the highest frequencies, the storage modulus is greater than the loss modulus, which indicates that the elastic nature of
the material is greater than its viscous nature in the high frequency (short time) region.
At the other end of the time scale, at long times, which corresponds to low frequencies,
the viscous nature dominates. Also note that the shape of the complex viscosity curve is
similar to the flow curve shown in Figure 7.5. As we will see, often, the complex viscosity versus frequency curve is a good approximation of the steady shear viscosity versus
shear rate curve. Frequency sweeps are, by far, the most common way of characterizing
polymer melts on rotational rheometers.
7.1.3.4 Dynamic Temperature Ramps
Another common dynamic test that is performed on rotational rheometers or on
dynamic mechanical analyzers (DMAs) is the dynamic temperature ramp. Like the
dynamic strain sweep, the most common frequency employed is 1 Hz (6.28 rad/s).
The temperature ramp is a rapid way of determining the temperature dependence of
the rheological properties of a material. The temperature range can be determined over
which a given material can contribute to the mechanical strength of a product. Glass
transition temperatures can also be determined from this type of testing. Before the
G (Pa)
G (Pa)
* (Pa.s)
106
0.8 MI HDPE
105
Cross-over modulus: 34,718 Pa
Cross-over freq: 10.20 rad/s
104
103
10–1
100
(rad/s)
101
102
Figure 7.14 Isothermal frequency sweep on 0.8 dg/min MI high density PE (HDPE) (190 °C, 2.16 kg).
Rheology of Polyethylene 251
test is started, a strain sweep test described in an earlier section should be performed.
This would be performed at the lowest temperature in the temperature ramp because
the LVR will be the shortest at the lowest temperature. Strains that are in the LVR at the
lowest temperature will also be in the LVR at higher temperatures. The most common
heating rate is 3 °C/min. An example of a DMA scan of HDPE is shown in Figure 7.15.
There some key features to note from this DMA scan. First, the storage modulus
exhibits an inflection point at around 50 °C, which corresponds to a transition temperature. Secondly, the storage modulus and the loss modulus decrease dramatically
around 125 °C. This represents the maximum temperature at which HDPE can contribute mechanically to a product. Because of the quick decrease in mechanical behavior,
PE cannot be used in applications where the contents are boiled. Note the two and
possibly three peaks in the loss modulus. These can be designated as glass transition
temperatures. There is a loss modulus peak at about −130 °C, a possible peak at −40 °C,
and a broad peak at 25 °C. There would be corresponding peaks for tan too. These can
be used to identify the glass transitions of the material. In general, the progression for
glass transition temperature (Tg) is the following:
Tg (DSC) ≤ Tg (storage modulus onset) ≤ Tg (loss modulus peak) ≤ Tg (tan Peak)
If the Tg is determined using a DMA scan, the technician must specify which parameter is being used and the test frequency. The heating rate should also be given.
DMA testing like that described above can be performed on rotational rheometers in
a torsional mode and can also be performed on dynamic mechanical analyzers (DMAs)
that impose a linear type of deformation on the specimen. The linear deformation can
be tensile, flexural, or compressive.
7.2 Melt Rheology
The following section will discuss the test geometries and test methods that are used to
perform rheological characterizations of polymeric materials.
Moduli G and G (Pa)
1010
G
109
108
107
G
106
105
104
–150 –100 –50
0
50
Temperature (°C)
Figure 7.15 DMA scan of HDPE at 10 rad/s [7].
100
150
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Handbook of Industrial Polyethylene and Technology
7.2.1 Extrusion Plastometer
Although the approach in the previous sections has been to describe rheological behavior in terms of fundamental rheological terms, the reality of the situation for industrial
PE resins is that the most commonly performed melt rheology test is the melt flow rate
determination. This is described in ASTM D1238 [8]. The ISO version is ISO 1133. The
melt flow rate (MFR) test is illustrated in Figure 7.16.
The plunger, barrel, and die are all made to standard dimensions. In this procedure,
the operator loads the barrel with pellets of the test resin, then applies the plunger on
top of the pellets that are now melting and, lastly, places a weight on top of the plunger.
After a suitable equilibration period, the operator cuts the molten strand and simultaneously starts a stopwatch. After a suitable amount of flow time, the operator cuts the
strand again and weighs the resultant extrudate. The weight is then converted to the
melt flow rate (MFR), which is expressed in grams per 10 minutes of extrusion flow;
i.e., dg/min. In the case where the load is 2.16 kg and the temperature is 190 °C, the
MFR is called the melt index or MI for short. The load of 2.16 kg and temperature of
190 °C is the standard condition E for PE. In some cases where the molecular weight
is very high, it is more convenient to use a load of 21.6 kg. The flow rate thus obtained
is referred to as the high load melt index (HLMI) or I21.6. Another weight that is used
occasionally is 10 kg.
At times, running melt flow rate tests with different loads and taking melt flow rate
ratios can provide an idea of the shear sensitivity and molecular weight distribution of
the material. For example, Lue reported ratios of the melt flow rates obtained with the
load of 21.6 kg to those obtained with the standard load of 2.16 kg for selected materials. He obtained values of 15.4 for a linear PE from a metallocene catalyst, 25.7 for a
linear PE from a Ziegler catalyst, 29.6 for a PE from a metallocene catalyst in which
there was long chain branching, and 171 for conventional LDPE [9].
Sometimes operators will use the term melt index even when the load is different
from 2.16 kg or the temperature is not 190 °C. It is advisable to verify the conditions
under which the MFR was determined. An example of a common notation would be
MFR (190 °C, 2.16 kg). This indicates the temperature at which the MFR was obtained
and the load.
The method described above is referred to as Method A. It is a completely manual method. In some cases, the determination has been automated so that the time it
Load
Piston
Polymer melt
Figure 7.16 Diagram of melt flow rate testing using an extrusion plastometer.
Rheology of Polyethylene 253
takes for the plunger with the weight to descend a certain distance is determined by an
electronic device on the instrument. This is referred to as Method B. The most common length of travel that is timed is 1 inch. The relationship between the MI for PE,
as determined by Method A, and the time in seconds to travel 1 inch, tB, is given in
Equation 7.10. The load is assumed to be 2.16 kg in both methods.
MFRA
1089
m
tB
(7.10)
where m is the melt density in g/cm3. This equation also implies that the melt density
can be obtained by performing Method A and Method B simultaneously and then rearranging Equation 7.10 to Equation 7.11.
m
MFRA
tB
1089
(7.11)
The main utility of the melt index test for PE is as a quality assurance test. The device
can be used to measure the shear rate, apparent shear stress, and apparent viscosity of
the molten polymer, but this is rarely done. One complicating factor is the fact that
the length/diameter (L/D) ratio of the standard die is only 3.7. With a small L/D ratio,
a substantial amount of pressure can be consumed as entrance pressure decrease, as
with low density PE, which will cause the output rates to be smaller than would be predicted from more fundamental viscosity measurements. Thus, comparisons between
rheological properties obtained from melt flow rate testing and other methods, such as
rotational rheometry and regular capillary rheometry, with L/D > 20, may not result in
the viscosity values agreeing with each other. Nevertheless, for materials like linear low
density PE (LLDPE) and HDPE, where there is not much long chain branching, a good
correlation can be obtained between MFR testing and viscosity measurements from
rotational or other capillary geometries [9].
Another drawback with MFR testing is that it is typically done as just a single point
measurement. As mentioned above, the typical load for the test is 2.16 kg, which corresponds to an apparent shear stress of 19.5 kPa. The resultant shear rate can vary,
depending on the molecular weight of the PE. For example, a 1 MI material would
have a shear rate of about 2.4 1/s, while a 10 MI material would have a shear rate of
about 24 1/s.
Even with its limitations, melt index determination will continue to be used as a quality control test. A useful empirical relationship for blends is shown in Equation 7.12 [1].
m
log( MIblend )
wi log(MIi )
(7.12)
i 1
where wi is the weight fraction of component i and MIi is its melt index. For example, a
50% blend of a 1 dg/min MI resin and a 10 dg/min MI resin, the resultant blend would
have an MI of 3.2 dg/min.
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Handbook of Industrial Polyethylene and Technology
7.2.2 Rotational Rheometry
7.2.2.1 Polymer Flow
We will now consider other methods for measuring rheological properties of molten
PE. We will do this by considering the lower shear rate region first, then will work up
to the high shear rate region.
There are two main geometries used in the low shear rate region. These are the cone
and plate geometry and the parallel plate geometry. In this section and in succeeding
sections, we will focus on fundamental rheological parameters, not on the somewhat
arbitrary flow parameter used with melt flow rate testing.
The formulas for converting angular velocity to shear rate and torque to shear stress
for these two commonly used geometries are provided in Table 7.3.
For both geometries, is the angular velocity in rad/sec, M is the torque in N-m, R is
the geometry radius, and H is the gap for the parallel plate geometry, in meters. For the
cone-and-plate geometry, is the cone angle (rad). For the derivations of these formulas, the reader is referred to references [1, 2]. The general dimensions that are used for
PE melt rheology are 25 mm diameters for both cone-and-plate and parallel plate geometry. A common angle for the cone is 2° (0.035 radians). The angles for commercial
cones typically range from as low as 0.5° (0.009 radians) to as high as 6° (0.105 radians).
There are some features to point out regarding these geometries. For the cone and
plate geometry, the shear rate and shear stress are constant throughout the gap. For the
parallel plate geometry, the shear rate and shear stress are 0 along the vertical axis. The
shear rate then increases in a linear fashion from 0 at the center to the maximum, R ,
at the edge. It is the shear rate at the edge that is used when reporting data from the
parallel plate geometry.
If the material is behaving in a Newtonian fashion, the shear stress will increase
linearly from 0 at the center to R at the edge for the parallel plate geometry, and its value
is calculated from the formula given in Table 7.3. If the material is non-Newtonian, then
the dependence of the shear stress on the radial distance will be more complex. In general, when steady flow tests are performed with the parallel plate geometry, the apparent
shear stress is used, calculated by the formula given below, as the shear stress for the
system. It is this shear stress value that is used to calculate the apparent viscosity.
To compare cone and plate data with parallel plate data, a correction to the shear
stress in the parallel plate data needs to be performed. This is shown in Equation 7.13.
Table 7.3 Relationship between machine parameters and rheological parameters.
Cone-and-plate
Parallel plate
Shear Rate ( )
Ω/β
ΩR/H
Shear Stress (s)
3M/(2πR3)
2M/(πR3)
Diagram
Rheology of Polyethylene 255
This correction is readily available on the software programs of the major commercial
rheometer manufacturers.
R Cor
R App
d log R app
1
3
4
d log R
(7.13)
where R-Cor is the corrected shear stress at the radius R and R-App is the apparent shear
stress at radius R.
For cone and plate geometry, there are 3 numbers that the operator needs to know:
the cone diameter, the cone angle, and the truncation gap. Although we refer to the
geometry as the cone-and-plate geometry, there is actually a truncation distance
machined into the cone, and it is at this gap that the cone should be set and run during
the experiment. The truncation gap increases as the cone angle increases. For example,
for a 2-degree cone, which is fairly standard, the truncation can be about 50 microns.
For a 4-degree cone, the truncation gap can be about 120 microns.
The fact that the shear rate and shear stress are constant everywhere in the gap is a
plus for the cone and plate geometry. However, with materials that are somewhat high
in molecular weight, it becomes increasingly difficult to position the cone at the specified truncation gap due to the high axial forces that are encountered. Often, these axial
forces do not dissipate in a reasonable amount of time. The way around this is to use the
parallel plate geometry. The most common gap for parallel plate geometry testing is 1
mm, but the operator does have the freedom to use larger or smaller gaps. With parallel plate geometry, the usual minimum gap that is recommended is 0.5 mm. The usual
maximum gap is 3 mm. As mentioned above, the most common diameter is 25 mm. For
low viscosity materials, the use of a 40 mm or a 50 mm diameter plate geometry is common. For very high viscosity materials, an 8-mm diameter plate geometry can be used.
With the background in rotational melt rheology established, let us look at the flow
curve for a 2 dg/min MI LLDPE (1-butene copolymer) resin in Figure 7.17. In this case,
the test temperature was 150 °C, and the geometry was a 25-mm parallel plate with a
gap of 1 mm. The shear rate spacing is typically 5 points per logarithmic decade. Note
that the material is Newtonian in the low shear rate region. Then it shows some shear
105
LLDPE
104
105
103
Stress (Pa)
Viscosity (Pa.s)
105
102
103
10–3
10–2
10–1
Shear rate (1/s)
100
101
101
Figure 7.17 Flow curve for LLDPE at 150 °C with 25 mm parallel plates and 1 mm gap.
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Handbook of Industrial Polyethylene and Technology
thinning. The measurements go well until a shear rate of 6.3 1/s is reached. Then there
is flow instability, and no more data can be obtained.
In the situation shown here, there was an unexpected huge decrease in the viscosity
at a shear rate of 6.3 1/s, which suggested a flow instability. This was fairly obvious. In
some situations, entering the region of flow instability is not as obvious. The best way to
observe this is to plot stress versus shear rate, as was done on Figure 7.17 on the right y
axis of this graph. A decrease in shear stress with increasing shear rate such as shown in
Figure 7.17 is an indicator that flow instability is occurring. The data should be omitted
from that point onward.
In the flow curve shown in Figure 7.17, the focus was on obtaining the viscosity
versus shear rate curve. This enables us to characterize the material and perhaps make
some predictions regarding the processability of the material. Another measurement
that some rheologists have found useful is the measurement of normal forces. When
polymer melts are sheared, there will be torque generated that is related to the viscosity of the material. However, there is also the generation of forces perpendicular to the
direction of flow. These forces try to separate the upper geometry from the lower one,
as shown in Figure 7.18. Eventually enough force develops, causing material to flow out
of the gap and ending the experiment. The material exits the fixture since the upper
geometry and lower geometry are maintained at a constant gap.
In the limited region where normal forces can be measured, the forces are very sensitive to the molecular architecture. In conventional viscosity measurements, the working parameters are shear stress and viscosity. When measuring normal forces with the
cone-and-plate geometry, the working parameters are the normal stress, the normal
stress difference, N1, and the normal stress coefficient . In regular viscometry, the
. For normal force measurements the relationship
shear stress has the relation,
2
. Whereas the steady zero-shear viscosity is dependent on the molecular
is N1
weight, Mw, to about the 3.4 power, the normal stress coefficient, in the limit, depends
on Mw to the 6.8 power.
There are experimental issues that make the normal force measurement challenging.
In the low shear rate region, it is difficult to get enough of a meaningful normal force
signal to be considered reliable. The minimum force that is measurable is about 0.1 N.
On the high force end, the same issue with flow instability that was present with the
102
LLDPE
104
101
103
100
102
10–3
10–2
10–1
Shear rate (1/s)
100
Figure 7.18 Flow curve of LLDPE with Normal Force measurements included.
10–1
101
Axial force (N)
Viscosity (Pa.s)
Normal stress coefficient (Pa.s2)
105
Rheology of Polyethylene 257
shear measurements can occur. Another issue is that most people who do rheological
testing are not familiar with these normal force measurements, so they do not explore
whether they would be useful or not. There have been cases where viscosity values have
been fairly similar, but clear cut differences were seen in the normal stress measurements, and operators have made more informed decisions regarding the suitability of a
given material for their application.
As shown by Figure 7.18, the last acceptable data point occurred at a shear rate
of 4 1/s, with unacceptable data occurring at higher shear rates. Overall, this is not
unusual. For much of the steady flow data on PE resins, flow instabilities start at shear
rates somewhere between 1 and 10 1/s. This is rather limiting if we want to characterize
our material over a wider range of shear rates and want to get closer to shear rates that
are actually experienced in the processing of the material.
As presented in the next section, a more extensive characterization of resins can be
obtained by performing dynamic frequency sweeps instead of the steady shear rate
sweeps that was just described. Dynamic frequency sweep data can be used to predict
steady shear properties.
7.2.2.2 Dynamic Frequency Sweeps
In Section 7.1.3.1, the fundamentals of dynamic oscillatory testing was provided. When
preparing to perform a frequency sweep on a given set of materials, the first step is
to perform a dynamic strain sweep or a dynamic stress sweep at 1 Hz to find a suitable strain to use for a subsequent frequency sweep. The goal is to select a strain that
will result in enough torque to give a reasonable signal and will also be in the LVR.
Once a strain has been selected, the frequency sweep can be performed. For PE melts,
the selected strain is usually between 1% and 10%. If the material is in the LVR, there
should be no difference between the data obtained from parallel plate geometries and
data obtained from cone and plate geometries.
Frequency sweep testing is described in ASTM D4440 [10]. As mentioned above, the
most common geometry is the 25 mm parallel plate with a 1.0 mm gap. The most common frequency range is 100 rad/s to 0.1 rad/s with 5 points per logarithmic decade. It
should not matter whether the measurements start at the high frequencies and then go
to low frequencies or start low and go high. The total amount of time will be about the
same. Sometimes it is preferred to start at high frequencies to get the initial data more
quickly. Each of the first few data points takes just a few seconds. Typically, the data
point at 0.1 rad/s takes about 3.5 minutes. Also, if there is a question of stability, more
good data points can be measured if the scan starts at the higher frequencies, which
take less time to obtain. It is recommended to use a nitrogen atmosphere when testing
PE melts to inhibit degradation.
The most common temperature for testing molten PE is 190 °C. Data on an LLDPE
resin and an LDPE resin are shown in Figure 7.19. The data displayed are the storage
modulus (G ) and the loss modulus (G ). In most cases, a single material would be
characterized by plotting storage modulus, the loss modulus, and the complex viscosity.
There are a few features to point out in Figure 7.19. First of all, by doing a frequency
sweep, we obtain information on the elastic and viscous properties of the material, not
just the viscous properties, as would be obtained by using a capillary rheometer. This
can provide useful information regarding the molecular architecture of the material.
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Handbook of Industrial Polyethylene and Technology
Storage modulus (Pa)
Loss modulus (Pa)
106
LLDPE 190
LDPE 190
105
104
103
102
10–1
100
101
102
Angular frequency (rad/s)
103
Figure 7.19 Graph of storage modulus (G ) and loss modulus (G ) for LLDPE and LDPE at 190 °C.
A parameter that rheologists use in characterizing materials is the cross-over point.
The cross-over point is where the G and G intersect. The cross-over point for the
LDPE resin in Figure 7.19 occurs at a frequency of about 100 rad/s, while for the LLDPE
resin it occurs at a slightly higher frequency. Note that the viscous nature dominates the
response at low frequencies. As the frequency is increased, the differences between G
and G diminish and eventually the elastic response will dominate over the viscous
response, creating the cross-over point.
In general, the cross-over modulus is related in an inverse fashion to the molecular
weight distribution (MWD) of the material, and the cross-over frequency is related in
an inverse way to the weight average molecular weight (Mw). For a given material, these
relationships need to be verified. In the example shown in Figure 7.19, the cross-over
modulus of the LDPE is lower than the extrapolated cross-over modulus of the LLDPE,
so we would expect the LDPE to have a broader MWD, which is reasonable. However,
the expected cross-over frequency is higher for the LLDPE, which would suggest a
lower molecular weight. Judging by the low frequency data, however, we would predict
that the LLDPE has the higher molecular weight. Focusing on zero shear viscosities is
generally more reliable than cross-over frequencies.
Another feature to note is that, if we model our polymers as a number of Maxwell
spring and dashpot elements in parallel, we should ultimately see the slope of log(G )
versus log( ) approach 2 and the slope of log(G ) versus log( ) approach 1. At this
point, the log( *) versus log( ) behavior should be Newtonian or very nearly so. It is the
complex viscosity in this region that should correlate the best with the Mw of the material.
When it comes to reporting data from frequency sweeps, the operator can pick from
a number of options, depending on what correlates the best to behavior in the field or
known properties of the material. The storage modulus, loss modulus, and complex
viscosity can be reported at one frequency, at multiple frequencies, and one can take
ratios of the properties at two different frequencies. The values at the lowest frequencies
are the most sensitive in distinguishing two or more resins.
In some studies, materials of the same type all fall on a master curve. Another way of
plotting data from frequency sweeps is the method of Van Gorp and Palmen where the
phase angle is plotted as a function of the complex modulus [11]. One of these graphs is
shown by Figure 7.20. Note that the conventional catalysts for LLDPE produced resins
[°]
Rheology of Polyethylene 259
1.7
1.6
1.5
1.4
1.3
LLDPE/butene
1.2
LLDPE/hexene
1.1
mLLDPE/hexene
1
LLDPE/octene
0.9
0.8
1.E+04
1.E+02
1.E+03
G* [Pa]
1.E+05
1.E+06
Figure 7.20 Van Gorp-Palmen plot of LLDPE and mLLDPE [11].
Viscosity (Pa.s)
105
LLDPE 190 Frq
LDPE 190 Frq
LLDPE 190 steady
LDPE 190 steady
104
103
102
101
10–3
10–2
10–1
100
101
Shear rate (1/s)
102
103
Figure 7.21 Complex viscosity and steady shear viscosity as a function of shear rate (frequency) for
LLDPE and LDPE at 190 °C.
whose Van Gorp-Palmen graphs were very similar while the resin from the metallocene catalyst gave a unique response.
Next, we will look at the graphs of the complex viscosity versus frequency for the
LLDPE and LDPE that were shown in Figure 7.19. We will do that in the context of a
very useful relationship in polymer melt rheology known as the Cox-Merz relationship:
*( )
( ) where
(7.14)
This is an empirical relationship that shows how to correlate the complex viscosity
from frequency sweep data with steady shear viscosity from flow data. As was shown
in the previous section, the shear rate range over which valid viscosity data can be
obtained is limited, with the maximum shear rate where the response is valid being
somewhere between 1 and 10 1/s. It is fairly routine to obtain frequency sweep data at
100 rad/s and sometimes even higher.
In Figure 7.21, the complex viscosity data for the LLDPE and LDPE were plotted as
a function of the shear rate. The flow data are also plotted in Figure 7.21. Note that the
steady flow and dynamic data agree well where they overlap, from 0.1 to 4 1/s. In this
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Handbook of Industrial Polyethylene and Technology
graph, we have shown the questionable high shear rate data for LDPE to illustrate the
improvement in the data when dynamic frequency sweeps are performed.
As shown by Figure 7.21, the zero-shear viscosity of the LLDPE is higher than that
of the LDPE: 4609 Pa∙s for LLDPE and 2456 Pa∙s for LDPE. Moreover, the LDPE exhibits shear thinning behavior at lower shear rates and to a greater degree than does the
LLDPE. This is in line with what we would expect.
In this example, an instability occurred in steady flow testing at about 8 1/s while the
data point looks good at the maximum frequency of 100 rad/s for the dynamic data.
Using the Cox-Merz relationship with dynamic data allows us to predict the processability of many resins where steady rotational flow testing was unable to due to the early
onset of flow instabilities.
At the lower end of the shear rate spectrum, Lue has shown that there is a good correlation between dynamic oscillatory data and melt flow rate, allowing an estimate of
the MFR from dynamic frequency sweeps. This is useful to verify MFR data or if one
does not have ready access to an extrusion plastometer [9].
7.2.2.3 Dynamic Frequency Sweeps, Time Temperature Superposition
Rheological properties are temperature dependent. For homogeneous, molten polymers, the temperature dependence is systematic. Using the concept that temperature
and time are interrelated, the use of time temperature superposition to predict the
behavior of a material in time regimes that are outside the typical range of measurement frequencies is possible. As was previously mentioned, the range of the typical
frequency sweep is from 100 to 0.1 rad/s. At times, measurements at about 628 rad/s
(100 Hz) are possible, but there can be inertia issues at higher frequencies. At the other
end of the frequency spectrum, measurements below 0.1 rad/s take an exceedingly long
time to obtain, so it would be useful to obtain this information in a more convenient,
less time-consuming way. Time temperature superposition allows rheological data to
be shifted to provide access to lower and higher frequency ranges.
The complex modulus as a function of angular frequency for an LDPE material at
different experimental temperatures is shown by Figure 7.22. The complex modulus
Complex modulus (Pa)
105
104
LDPE 190
LDPE 170
LDPE 150
LDPE 130
103
102
10–1
100
101
102
Angular frequency (rad/s)
103
Figure 7.22 Complex modulus as a function of angular frequency at selected temperatures.
Rheology of Polyethylene 261
is used here, but time temperature superposition could have been performed with the
storage modulus or the loss modulus. If the storage modulus or loss modulus were
shifted, the process would work best with the parameter that is the more dominant.
Typically, the loss modulus will dominate at the low frequencies and the storage modulus will dominate at the high frequencies. If the shifting is performed with the complex
modulus, both moduli contribute with the loss modulus providing a greater contribution at the low frequencies and the storage modulus at the high frequencies.
In our previous section, we saw that, at 190 °C, we were able to get to the Newtonian
region with conventional flow testing, so, to extend our knowledge in the other direction, namely, the high frequency range, we will set 190 °C as our reference temperature
and, sequentially, horizontally shift the Complex Modulus data from the lower temperatures onto the data at 190 °C. This will have the effect of giving us data at 190 °C at
higher frequencies than those actually employed. If we had done additional testing at
temperatures above 190 °C, then it would have the effect of going to lower frequencies
at 190 °C.
The shifted complex modulus data is shown in Figure 7.23. Note that setting the
reference temperature 190 °C and shifting the lower temperature curves so that they
superimpose into one master curve has the effect of extending our knowledge of the
viscoelastic properties at 190 °C up to almost 1000 rad/s, which would not be possible
on the rheometer at 190 °C. In this case, shifting was just performed horizontally. At
times, it is more appropriate to shift both horizontally and vertically. For the vertical
shifting, often, this consists of corrections for temperature and density. Other times,
the vertical shifting is just done empirically. On the figure, the x-axis is the angular
frequency . However, since these are shifted data, it is more appropriate to display the
x-axis as *aT. , where aT is the horizontal shifting to get the graph at temperature, T, to
superpose with the data at the reference temperature.
The individual components of the complex modulus are shown in Figure 7.24; i.e.,
the storage modulus and the loss modulus. Individually these curves shifted well. A
cross-over point for G and G is apparent at a frequency of about 50 rad/s. As previously discussed, the cross-over modulus and cross-over frequency can often be used
to characterize a material in terms of its molecular weight distribution and molecular
weight.
Complex modulus (Pa)
105
104
*** 190.0
170.0
150.0
130.0
103
102
10–1
100
101
102
Angular frequency (rad/s)
103
Figure 7.23 Shifted complex modulus data for LDPE with 190 °C as the reference temperature.
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Handbook of Industrial Polyethylene and Technology
Storage modulus (Pa)
Loss modulus (Pa)
105
104
103
LDPE 190 (storage modulus)
LDPE 190 (loss modulus)
Cross-over modulus: 19607.8 Pa
Cross-over x value: 7.76e1 rad/s
102
101
10–1
100
101
102
Angular frequency (rad/s)
103
Figure 7.24 Master curves of G and G with 190 °C as the reference temperature.
aT (x variable)
10
1
120
LDPE 190 (aT (x variable))
TTS Arrhenius
Activation energy: 54.7796 kJ/mol
Tref: 189.996 °C
TTS WLF
c1: 6.08590
c2: 457.023 K
Tref: 189.996 °C
130
140
150
160
Temperature (°C)
170
180
190
Figure 7.25 Plot of shift factors aT relative to a reference temperature of 190 °C. The shift factors were
calculated using both the WLF model and the Arrhenius model.
In addition to obtaining the master curves, it is also useful to plot the shift factors
against temperature to show the temperature dependence of the rheological properties. Figure 7.25 shows the shift factors at each temperature. Note that the shift factor is
1 for 190 °C, which indicates that this is our reference temperature and the curve at this
temperature has not been shifted.
There are two commonly used models for fitting the aT data. One is the WilliamsLandel-Ferry (WLF) model as provided by Equation 7.15.
log aT
C1 (T Tref )
(C2 T Tref )
(7.15)
The other is the Arrhenius model, which is given in Equation 7.16.
log aT
Eact 1
2. 3R T
1
Tref
(7.16)
Rheology of Polyethylene 263
104
Complec viscosity (Pa.s)
LDPE 190 (1) (complex viscosity)
103
102
101
10–1
100
101
102
Angular frequency (rad/s)
103
Figure 7.26 Complex Viscosity versus aT for LDPE at 190 °C.
104
Viscosity (Pa.s)
LDPE 190
103
Analysis: Williamson
Zero-rate viscosity:2607.89 Pa.s
Consistency: 0.267097 s
Rate index: 0.608764
R2: 0.999416
102
101
10–3
10–2
10–1
100
101
Shear rate (1/s)
102
103
Figure 7.27 Extended flow curve for LDPE with data fit to the Williamson model.
Overall, the WLF model seems to fit many aT-temperature graphs better than the
Arrhenius model, but the Arrhenius model has the benefit of being a single-parameter
model; i.e., the flow activation energy, Eact. In the graph shown in Figure 7.25, both
models fit the data very well. For this LDPE, the flow activation energy is 54.8 kJ/mol,
which is within the range expected for LDPE.
The complex viscosity versus frequency curve, which we will use to construct the
extended flow curve by making use of the Cox-Merz relationship for the dynamic
data and adding the steady viscosity data from the steady flow testing, is shown in
Figure 7.26. Next, the extended flow curve for this LDPE resin is shown in Figure 7.27.
This flow curve was made using the Cox-Merz relationship with the dynamic data and
including the valid flow data. The flow curve was then fit to a model. A fairly simple
model is the 3-parameter Williamson model, whose curve fitting results are shown in
Figure 7.27.
The Williamson equation is shown in Equation 7.17:
( )
0
(1 (
)C )
(7.17)
where o is the zero-shear viscosity, which is related to the weight average molecular weight, Mw, is called the consistency and can be related to the molecular weight
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Handbook of Industrial Polyethylene and Technology
Table 7.4 Constants for the PE resins just discussed for the Williamson-Arrhenius models.
Parameter
LDPE
A (Pa∙s)
2.09
Eact (kJ/mol-K)
54.8
crit
C
(Pa)
8.84
LLDPE
1.01
10 3
32.5
3
10
8.48
0.61
104
0.56
distribution (MWD) of the resin, and C is the rate index, which relates to the processability of the resin. In some literature articles, this is called the Cross model, but the
Cross model referred to here is
( 0
(1 (
( )
)
C
) )
(7.18)
As is shown in the graph, the Williamson model parameters are the following: o =
2608 Pa∙s; = 0.27 s; and C = 0.61. The parameter, , can also be expressed as o/ crit,
where o is the zero-shear viscosity at a particular temperature, and crit is the critical
shear stress where = o/2. We assume that this is independent of temperature. With
these changes, we can write Equation 7.17 as the following:
0
C
1
0
(7.19)
crit
Furthermore, the zero-shear rate viscosity can be written as the following:
0
Ae
Eact
RT
(7.20)
where Eact is the flow activation energy from fitting the data to the Arrhenius model, R
is the universal gas law constant, and T is the temperature in Kelvin. For these resins,
the Williamson-Arrhenius model constants are provided in Table 7.4. Using these constants would give the viscosity curves in Figure 7.28. Other models include the following [1, 2]: Cross, Carreau, Carreau-Yasuda, and Sisko.
In summary, the most commonly performed test on polymer melts with a rotational
rheometer is the dynamic frequency sweep from 100 to 0.1 rad/s. Time-temperature
superposition (TTS) can be used to predict rheological properties at frequencies that
are normally outside the range of conventional testing. By going to temperatures below
the reference temperature, shifts can be made to higher frequencies. The lower temperature limit, however, is the temperature at which crystallization begins. There must
be no crystallinity for TTS to work properly. At the other end of the frequency range,
Rheology of Polyethylene 265
Steady or complex viscosity (Pa-sec)
1.0E+04
LDPE
LLDPE
1.0E+03
1.0E+02
1.0E+01
1.0E–02 1.0E–01 1.0E+00 1.0E+01 1.0E+02 1.0E+03
Reduced shear rate (1/sec) or frequency (rad/sec)
Figure 7.28 Data from Williamson-Arrhenius model for LDPE and LLDPE at 190 °C.
rheological properties can be predicted at lower frequencies by going to temperatures
above the reference temperature. The limiting factor here would be not to go to temperatures that would cause degradation to occur too rapidly.
A discrete or continuous relaxation spectrum in which spring and dashpot elements
are determined can be developed to model the storage and loss modulus as a function
of frequency. The element, Gi, represents the elasticity of the i th element. The element,
, represents the viscosity of the same element. However, it is more common to divide
i
the viscosity by the modulus to determine a relaxation time. Several researchers have
taken these relaxation spectra and related them to the entire molecular weight distribution of the polymer melt, not just Mw or Mw/Mn. For more information, the reader is
referred to an article by Thimm et al. [12].
While some relationships have been shown between rheological data and separation methods like gel permeation chromatography (GPC) that provide direct molecular
weight distribution chromatograms, it is difficult to get exact and complete correlation
between the two methods. Small differences in GPC data that are within experimental
error of this method can have huge implications in estimating the rheological properties of a polymer melt [13]. Moreover, a wide range of frequencies is not typically accessible with conventional rheological testing to completely predict the molecular weight
distribution that would be expected from GPC data.
Dynamic testing on melts can also be useful in determining the stability of a given
resin. The most common test for this is the dynamic time sweep. Degradation can be
accelerated by using an air atmosphere and monitoring the change in rheological properties with time. In general, G will exhibit more change than G . The changes in G can
occur in both directions. For many materials, including most PEs, an increase in G will
occur with time as oxidation, chain extension, and cross-linking occur. For some materials such as polypropylene (PP), the rheological properties will decrease with time,
indicating chain scission.
Rotational rheometry, with small specimen requirements, is also very useful in characterizing parts that have come back from the field if they have failed.
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Handbook of Industrial Polyethylene and Technology
2.0×10–2
Compliance (1/Pa)
HDPE
Pure viscous
response
Viscoelastic response
1.5×10–2
Pure elastic response
1.0×10–2
5.0×10–3
0.0×100
Stress
0
20000
40000
Step time (s)
60000
80000
Figure 7.29 Diagram of a creep test that shows constant stress, an elastic response, a viscous response,
and a viscoelastic response.
7.2.2.4 Creep Testing
Creep testing was described previously in Sections 7.1.1 and 7.1.2.2. Along with stress
relaxation, it is a unidirectional type of test that gives the viscoelastic properties of
a material. It consists of the instantaneous imposition of a shear stress and then the
monitoring of the resultant strain as a function of time. In creep testing, the key parameter is the compliance, J(t), which is the ratio of the time-dependent strain, (t), to the
imposed, constant shear stress, . The stress history and the characteristic responses for
a solid, a fluid, and a viscoelastic material are shown in Figure 7.29.
As in dynamic oscillatory testing and stress relaxation, operation in the LVR is preferred where the compliance is independent of stress.
We generalize Equation 7.2 to include several Kelvin elements in series that would
give the following relationship:
m
J (t )
J0
Ji 1 e
i 1
t
ti
t
(7.21)
0
As time progresses, the behavior will be governed by the final term, t/ o, so, in the
terminal region, we should be able to obtain a viscosity of the PE sample. Note that we
use the notation, 0, which indicates that this is the zero-shear or Newtonian viscosity.
If we are in the LVR, then the viscosity we obtain will be the Newtonian viscosity, which
is directly related to the weight-average molecular weight of the material.
If we are in the terminal region, where J(t) has become a straight line with a slope pf
1/ 0, then we can rewrite Equation 7.21 for this region as the following:
J (t )
J e0
t
(7.22)
0
In addition to obtaining the zero-shear viscosity, the recoverable compliance, Je0, is
obtained. The recoverable compliance is related to the molecular weight distribution of
the material, specifically, the quantity (Mz/Mw)2 [4].
Rheology of Polyethylene 267
3×10–3
Compliance (1/Pa)
PE at 190 C
Analysis: Discrete retardation spectrum
Je: 9.94E–4 1/Pa
n0 = 2.54E6 Pa-sec
2×10–3
Analysis: Straight line
a: y-intecept: 9.11116e–4
b: slope: 4.09717e–7
1×10–3
n0 = 2.44E6 Pa-sec
0×100
0
1000
2000
Step time (s)
3000
4000
Figure 7.30 Creep of a high molecular weight HDPE at 190 °C. Mw is approximately 6×105 g/mol.
A creep curve for a high molecular weight PE is shown in Figure 7.30. Note that the
slope of the curve did not become constant until about 2000 seconds (or 33 minutes).
For some high molecular weight resins, creep testing is a useful way of characterizing
them for their elasticity and their zero-shear viscosity.
The software packages associated with most commercial rheometers can fit the creep
data to a series of Kelvin elements and display the constants for the individual elements;
Ji and ti. The parameter ti is the ratio of the viscosity of the i th element, i, to the spring
constant, 1/Ji. The recoverable compliance, Je0, is the sum of the instantaneous compliance, J0, and the contributions of each of the Kelvin elements. Je0 is also the intercept
of straight line that is drawn using the points that are in the steady flow regime. In the
example shown here, the agreement between the two methods of obtaining Je0 and o is
very good. Using the software and making the calculations from the discrete retardation
spectrum, the parameters are Je0 = 9.94 × 10–4 1/Pa and o = 2.54 × 106 Pa∙s. By just using
the last 2 data points shown here and drawing a straight line, an intercept was obtained;
Je0 = 9.11 × 10–4 1/Pa, and by taking the reciprocal of the slope of the line, a zero-shear
viscosity was obtained; o, = 2.44 × 106 Pa∙s. Testing on several PE grades that ranged in
MI from 0.65 to 4.4 showed that the zero shear viscosity obtained by the creep method
agreed well with the zero shear viscosity obtained by dynamic frequency sweeps [14].
Finally, another way to obtain Je0 is to perform a creep recovery test where the applied
stress is removed. The limiting recoverable compliance value should also equal Je0. At
times, the Je0 values will not agree. This usually indicates that the creep test did not get
completely in to the steady flow regime or that the recovery was not given sufficient
time to reach its limiting value.
7.2.2.5 Creep, Time Temperature Superposition
Time temperature superposition (TTS) can also be performed with creep data.
Moreover, TTS can also be performed with stress relaxation testing and flow testing.
For example, compliance data for LDPE is plotted as a function of time at selected
temperatures as shown in Figure 7.31. Note that for TTS on creep data, the plot is made
with the log of the compliance against the log of time.
Like before, the data are shifted to produce a master curve. A reference temperature of 190 °C was used. When the TTS shifting is performed, the x-axis will be log
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Handbook of Industrial Polyethylene and Technology
100
Compliance (1/Pa)
10–1
10–2
LDPE 190
LDPE 170
LDPE 150
LDPE 130
10–3
10–4
10–5
10–1
100
101
Step time (s)
102
103
Figure 7.31 Log compliance versus log time for LDPE at several temperatures.
Compliance (1/Pa)
100
LDPE 190
10–1
10–2
10–3
10–4
10–5
10–2
10–1
100
101
Step time (s)
102
103
Figure 7.32 Master curve of creep data on LDPE for a reference temperature of 190 °C.
10
aT (x variable)
LDPE 190
Analysis: TTS Arrhenius
Activation energy: 56.8035 kJ/mol
Tref: 189.996 °C
R2: 0.999649
1
120
130
140
150
160
Temperature (°C)
170
180
190
Figure 7.33 Shift factors for LDPE creep data at a reference temperature of 190 °C using the Arrhenius
model.
t/aT. Recall that, when we did TTS with frequency sweeps, the x-axis was log aT.
Figure 7.32 shows the master curve for the creep data. Again, testing at temperatures
below the reference temperature will assist in providing data at shorter times, while
testing at temperatures above the reference temperature will help provide data at longer
times. As expected, the shift factors shown by Figure 7.33 are nearly the same as those
obtained with the dynamic frequency sweep data shown in Figure 7.25.
Rheology of Polyethylene 269
Figure 7.34 SER fixture for testing polymer melts in elongation [15].
Rotating drum
X-ducer shaft
Pair of gears
Motor axis
Fixed outer axis
Figure 7.35 EVF fixture for testing polymer melts in elongation [16].
7.2.2.6 Extensional Testing
In recent years, fixtures have been developed to determine elongational viscosity on
polymer melts with rotational rheometers, with much of the pioneering work being
done by Sentmanat [15] with the Sentmanat extension rheometer (SER) fixtures that he
developed. These work by elongating a molten polymer strip, secured with clamps that
fit inside a drum, with counter-rotating drums. Prior to this development, determining
the extensional viscosity of polymer melts was more challenging. Some experimenters
have performed tension testing on melt extrudate from capillary rheometers. Others
have done direct tension testing or stretching between counter rotating rollers. In these
situations, the operator often had to use a silicone fluid that was the same temperature
and approximately the same density as the polymer melt to perform the test.
The design of the SER fixture is shown in Figure 7.34, with its counter-rotating
drums. Another design is the extensional viscosity fixture (EVF), which has one stationary drum in the center and the other drum rotating around the stationary drum while
rotating about its own axis. A schematic of the EVF design is shown in Figure 7.35. This
is specifically designed for the ARES and ARES-G2 rheometers, which are made by TA
Instruments.
With both of these systems, the design allows the operator to use the oven from the
rotational rheometer to heat the specimen into its melt regime. The temperature must
be high enough to completely melt all the crystallites but not too high to cause sagging during the experiment. For these types of experiments, it is recommended to have
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Handbook of Industrial Polyethylene and Technology
10
LDPE 494A
106
LLDPE PL1880
Elongation viscosity
E
*(t) [Pa s]
Comparison LDPE 494A and LLDPE affinity
7
5
10
LDPE (branched)
Extension rate 0.1 s–1
Extension rate 1 s–1
LLDPE (linear)
Extension rate 0.1 s–1
Extension rate 1 s–1
104
103
102
0.01
0.1
1
Time t [s]
10
Figure 7.36 Graph of elongational data for LDPE 494A and LLDPE AFFINITY PL 1880 resin [15].
video capabilities to watch the specimen during the experiment to verify that sagging
has not occurred.
In both systems, the maximum strain that can be achieved before the specimen
begins to wind around itself is about 3.4. The most common test is the constant elongation rate test. The practical range for the elongation rates is from 0.01 to 10 1/s. At the
lowest rate, the test will take 3.4/0.01 = 340 seconds (5:40 minutes). At the highest rate,
the test will take only 0.34 seconds. In most labs, the common test rates are 0.1, 0.3, and
1 1/s, which take 34 seconds, 11.1 seconds, and 3.4 seconds, respectively.
The elongational viscosity is plotted against time, with both axes being logarithmic,
as shown by Figure 7.36 [16]. The materials shown here are a LDPE 494A resin and
LLDPE AFFINITYTM PL 1880 resin (both from The Dow Chemical Company). The
key feature is the strain hardening behavior of the LDPE resin, which is known to have
long chain branching. The LLDPE resin, which has only short chain branches from the
1-octene comonomer, shows a rather flat curve with little indication of strain hardening.
7.2.3 Capillary Rheometry
Capillary rheometry has some similarities to melt index determinations in that molten
resin is pushed through a capillary die via a plunger in a heated barrel. However, in its
usual mode, capillary rheometry differs from melt index determination in a few respects.
First, most capillary rheometers are rate driven, not force driven as it is in a melt flow
rate determination. For capillary rheometry, the operator sets the cross-head speed, and
after a suitable equilibration time, records the equilibrium force at that particular speed. A
schematic of a capillary rheometry is provided in Figure 7.1. Secondly, the dies in capillary
rheometry are designed with higher length/diameter ratios to minimize the effect of die
entrance pressure losses. These entrance pressure losses play a role in melt flow rate determination, where the length/diameter ratio is only 3.7. Thus, trying to correlate melt flow
rate with genuine viscosity values can be difficult if entrance pressure drops are significant. Dies in capillary rheometry typically have a length/diameter ratio of greater than 20.
Often times, operators use the dual bore design capillary rheometer, with one of the bores
having a typical length/diameter ratio of 20 or greater, and a second die with length/diameter ratio of, essentially, 0. With this setup, the entrance pressure drop can be determined
Rheology of Polyethylene 271
Table 7.5 Equations for fundamental rheology parameters for capillary rheometry.
Parameter
Shear Rate,
Equation
= 32Q/( DC3)
(1/s) – apparent
Shear Stress, (Pa) – uncorrected
P/(4(L/DC))
Viscosity,
P DC4/(128LQ)
/ – apparent
directly and then subtracted from the corresponding measurements of the regular die.
The subtraction of the entrance pressure effects is known as the Bagley correction [1].
This has been related to elongational flow [17]. A third difference is that, with capillary
rheometry, the behavior of the molten material is determined at a number of rates, not just
one set of conditions as is melt flow rate testing. Depending on the design, the flow properties of a molten plastic over a very wide range of shear rates can be characterized, often
including the high shear rates that are encountered in operations like injection molding.
Testing with a capillary rheometer can also provide knowledge regarding flow phenomena of sharkskin and melt fracture [18]. With a capillary rheometer, some operators have added visual accessories or other devices to characterize extrudate swell.
Extrudate swell or die swell is a phenomenon whereby the molten polymer extrudate
has a diameter greater than the die from which it was extruded. This is important in
blow molding and injection molding as it relates to the uniformity of the parts that are
produced as well as the amount of excess material that is discarded due to excessive
extrudate swell. Common ways of measuring extrudate swell include camera recordings, laser micrometers, and simple weighing of a specified length of extrudate and
conversion to extrudate swell measurements.
Another application is the Rheotens-type of test where extrudate from a capillary
rheometer is stretched to determine the strength of the melt as a function of the drawdown ratio. This provides information on the suitability of a resin for applications like
fiber spinning, blown film, and blow molding [19].
In their simplest form, the equations that are used in capillary rheometry are those
shown in Table 7.5 for the conventional circular die [15].
In these equations, Q is the extrusion rate (cm3/s); Dc is the die diameter (cm); P is
the pressure needed to extrude the molten plastic at a given rate; and L is the length
(cm) of the die. Furthermore, Q can be seen to equal V (DB2/2), where V is the velocity of the plunger through the barrel, and DB is the barrel diameter. The units here
would be cm/s for the plunger velocity and cm for the barrel diameter. Test methods are
described in ASTM D 3835 [18] and ISO 11443.
These equations provide apparent values for the shear rate and shear stress. Just as
there was a correction factor for the stress of non-Newtonian fluids with the parallel
plate geometry, there is a corresponding correction to the shear rate for non-Newtonian fluids in capillary flow. This is called the Rabinowitsch correction. It is shown in
Equation 7.23 [4]:
corr
app
(3n 1)
where n
4
d log
d log
(7.23)
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Viscosity at 170 C (Pa-sec)
1.0E+04
1.0E+03
LLDPE
LDPE
1.0E+02
1
10
100
Shear rate (1/sec)
1000
Figure 7.37 Flow curve for LLDPE and LDPE at 170 °C from capillary rheometry [17].
At the lowest shear rates where the material is Newtonian, n = d(log )/d(log ) = 1,
and thus corr
In the high shear rate limit, n may approach 0, so the relation
app .
would be corr
and thus the largest correction will be a 75% decrease in the
app /4,
apparent shear rate.
With the stress, as mentioned above, there is a correction to account for entrance
pressure losses. This is the Bagley correction which has been mentioned earlier. Some
operators have done tests with dies of various L/D ratios and performed a linear regression of pressure as a function of L/D ratio for each shear rate [12]. Sometimes, the
correction is simply done by using the dual bore approach with one die being a conventional die with L/D ≥ 20 and the other die having L/D ≈ 0. The corrected stress can be
expressed as Equation 7.24:
corr
(P P0 )
(4 L /D)
(7.24)
where P0 is the entrance pressure loss.
In many cases, only a comparison is required for two or more materials, and there
is no concern about correction factors for the shear rate or entrance pressure losses for
the stress. Apparent values for the shear rate, the shear stress, and the viscosity may be
acceptable in these cases.
Here is an example of a flow curve from capillary rheometry on two PE resins, each
with an MI of about 1. The shear viscosity data are shown in Figure 7.37. Both materials have the Rabinowitsch and Bagley corrections applied [17]. There are a few things
to point out here. One is that, although both resins have a similar MI, the LLDPE has a
higher viscosity when run on the capillary rheometer. Note that the entrance pressure
drop is taken into account. This graph indicates that much of the MI measurement
result of the LDPE, with its long chain branching, is due to entrance pressure loss.
The more linear LLDPE does not have the same long chain branching, so it does not
have the same effect from the entrance pressure loss, and the MI measurement is due
more to viscosity than it is to entrance pressure drop. Also, note that the LDPE exhibits
greater shear sensitivity than does the LLDPE.
Rheology of Polyethylene 273
Table 7.6 Relation of melt flow rate parameters to fundamental rheology parameters.
Correlations are to be performed with discretion.
Method
Shear rate, 1/s
Shear stress, Pa
Viscosity, Pa∙s
A
1.85*MFRA/ Melt
9035*W
4880*W* Melt/MFRA
B
2010/tB
9035*W
4.495*W*tB
Even with all the caveats regarding the correlation of melt index data to genuine
viscosity data from capillary rheometry, some operators have wanted to calculate shear
rates, shear stresses, and viscosities from melt index values, just to get some idea of the
more fundamental rheological parameters. Table 7.6 has the relationships that can be
used to try to correlate MI data with other measurements. Method A is the cut and
weigh method. Method B is the automatic timed descent method. These methods were
discussed in Section 7.2.1.
In Table 7.6, MFRA is the melt flow rate (dg/min) by Method A; Melt is the melt
density (g/cm3); W is the weight of the piston plus the added weight, e.g., 2.16 kg; tB is
the time in seconds for the plunger to descend 25.4 mm (1 inch), as measured by an
electronic device for Method B.
Sometimes the best that can be achieved is to estimate the viscosity to within an
order of magnitude for materials that have high entrance pressure losses. For materials
other than polyolefins and materials that have low entrance pressure losses, the agreement between the rheological parameters calculated from Table 7.6 and either rotational or capillary rheometry can be very good.
7.2.4 Time Temperature Superposition with Capillary Data
Time temperature superposition for capillary data is essentially the same as that performed for dynamic frequency sweep data and as described in Section 7.2.2.3. This is
achieved by plotting log stress versus log shear rate at multiple temperatures, selecting
a reference temperature, and then shifting the data to get superposition, as shown in
Figure 7.38 [21].
7.3 Dynamic Mechanical Testing on Solids
and Solid-Like Materials
7.3.1 Dynamic Testing with Rotational Deformation
Most of the testing discussed has been on polymer melts. These are performed isothermally or in a series of isothermal steps. The dynamic temperature ramp method using
a torsion fixture is a method of determining the temperature range over which a material can contribute mechanically to the part it is in and determining glass transition
temperatures. For dynamic mechanical analysis (DMA) testing, the most common test
is the dynamic temperature ramp. These are usually done with a heating rate of 3 °C/
min and at a frequency of 1 Hz. Some operators have used a heating rate of as low as
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Handbook of Industrial Polyethylene and Technology
1.0E+06
Shear stress (Pa)
1.0E+06
1.0E+05
1.0E+05
170
190
210
1.0E+04
1.0E+04
Individual temperatures
Master curve at 190 C
1.0E+03
1.0E+03
1
10
100
Shear rate (1/sec)
1
10
100
Shear rate (1/sec)
Figure 7.38 LLDPE capillary rheometer data [20]. Data at each temperature is shown on the left. The
master curve at 190 °C is shown on the right.
1 °C/min, while others have gone as high as 5 °C/min. The faster the heating rate the
shorter the test. However, there can be a thermal lag in the test specimen, so transition temperatures will seem higher than they should be. If a very low heating rate is
used the time it takes to perform the test will increase, and the test time may become
prohibitively long. For the frequency, some operators have chosen 10 rad/s, which is
1.6 Hz. Some operators will go to 10 Hz for their frequency. For temperature ramps, it
is strongly recommended to use only one frequency in a test.
As mentioned earlier in dynamic testing, the first step is to determine the LVR of
the material. This should be done at the lowest temperature of the test. Temperature
ramps are usually performed from low temperature to high temperature, so the lowest
temperature is the initial temperature. Most of the time, the operator will select a strain
that is in the LVR at the lowest temperature and then use that throughout the temperature ramp. Often, the operator will set a minimum oscillatory torque or use an autostrain feature that increases the strain once a certain threshold has been reached so that
meaningful data can be acquired even after the specimen has weakened substantially.
At the start of a temperature ramp the strain is typically about 0.1%. As the temperature
increases and the specimen goes through a transition, the 0.1% strain may be insufficient to generate a meaningful torque. In this case, the instrument software can call for
a greater strain to give a reasonable resultant torque.
For rotational rheometers, the geometry that is used the most for this type of test is
the torsion rectangular geometry. It provides a good grip on the specimen, and the fact
that there is a large gap between the fixtures (~30 mm) means that the specimen will
not be overly stiff for the rheometer. For some materials such as adhesives the parallel plate geometry is preferred. With the torsion rectangular geometry, an axial tensile
force is imposed prior to the temperature ramp and the instrument maintains at least
some axial force throughout the test. A typical axial force at the beginning of the test
is 2 N. It is common to set a tolerance of 1.75 N, so that the upper geometry does not
move all that much, even when the material goes through a transition.
For softer materials, the parallel plate is useful. At times, it is helpful to impose a
compressive force to maintain good contact between the plates and the specimen. In
Rheology of Polyethylene 275
Torsion
rectangular
Parallel
plate
Figure 7.39 Diagram of geometries that can be used for temperature ramps with rotational rheometers.
setting up the experiment, the operator will include a maximum amount of distance
that the upper plate can descend so that not all the test material will be pushed out from
between the plates.
Maintaining good contact between the specimen and the plates can be a challenge
for the parallel plate geometry. For some materials, sufficient contact can be achieved by
using cross-hatched or sandblasted geometries to improve the grip. Another technique
used by some operators is to glue sandpaper to their plates. Finally, some operators will
glue the specimens to the upper and lower plates with the hope of achieving good contact. If there is a way to get the torsion rectangular fixture to do the testing, this usually
works better than using the gluing methods that have just been mentioned for the parallel plate geometry. This author has seen the rheological properties of elastic materials
increase as a function of applied compressive force with the parallel plates, but even the
highest values with the highest compressive forces did not match the values obtained
with torsion rectangular geometry. For this type of test, cone-and-plate geometries are
not used because of the non-uniformity of the gap. Figure 7.39 shows diagrams of these
geometries that are used with rotational rheometers. An example of a dynamic temperature ramp with the torsion rectangular fixture has been given in Figure 7.15.
7.3.2 Dynamic Testing in Linear Deformation
For solids and solid-like materials, it is actually more common to use linear deformation
rather than shear deformation. This testing is performed on a dedicated DMA instrument. The common clamps that are used for DMA testing are shown in Figure 7.40.
For thin films, the film clamp is the obvious choice. For thicker and stiffer materials, flexural deformation, which includes single and dual cantilever and the 3-point
bend clamp, is the better choice. There are some situations where the shear sandwich
is the best geometry, but, if the operator also has a rotational rheometer, the rheometer
is more suitable than the shear sandwich. If a material goes through a transition and
becomes molten, it may drip from the shear sandwich and fall onto the stage of the
instrument. For some samples, like foamed materials, compression testing is a good
option.
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Handbook of Industrial Polyethylene and Technology
S/D cantilever
Tension-film
Shear-sandwich
3-Point bending
Compression
Figure 7.40 Common clamps for DMA testing [21].
In flexural testing, there are operators who have preferences for the cantilever clamp
over the 3-point bend. The cantilever has a good deformation that goes up and down
the same amount. The issue with some materials is clamping effects. More operators
prefer the single cantilever over the dual cantilever. The dual cantilever essentially doubles the force that can be obtained with the single cantilever clamp.
The 3-point bend is considered a purer type of flexural deformation since there is no
clamping. However, a pre-load static force must be applied to ensure that the fulcrum
always stays in contact with the specimen. As the material goes through the transition,
there can be some deformation of the specimen due to the pre-load static force, and the
specimen can become warped, which will affect the quality of the data.
Most of the instruments that do the linear deformation shown in this section are dedicated to linear testing. They cannot do rotational deformation. In recent years, though,
some rotational rheometers have added linear deformation capabilities. Examples of
these are the TA Instruments ARES-G2 and the TA Instruments DHR rheometers.
7.3.3 Dynamic Temperature Ramps
As mentioned previously, dynamic temperature ramps indicate over what temperature
range a given material may be used and still contribute mechanically. DMA testing
by temperature ramps is also useful for identifying the primary glass transition, often
called the transition, and secondary transitions, if there are any, which are the and
transitions. The transition is associated with cooperative motion among a large number of chain segments. Secondary transitions are associated with intramolecular rotational motion of a few atoms, side group motion, and internal motion within the side
group [23].
The presence of a transition is often associated with good impact properties of a
material, as long as the origin of the transition is from the main chain or from a separate phase. Materials like polycarbonate (PC), polytetrafluoroethylene (PTFE), PE, and
polyvinyl chloride (PVC) have peaks and have good impact properties. A material like
polystyrene (PS) does not have a peak and does not have good impact properties [23].
Rheology of Polyethylene 277
4000
160
–107.27 °C
HDPE
–119.26 °C
140
48.62 °C
–103.96 °C
120
0.15
G’
–88.79 °C
2000
80
13.33 °C
60
39.61 °C(I)
–103.11 °C
1000
100
G”
[––––] Loss modulus (MPa)
3000
Tan delta
Storage modulus (MPa)
0.20
0.05
40
tan
0
–150
–100
–50
74.29 °C
0
50
Temperature (°C)
20
150
100
Universal V4.1D TA instruments
Figure 7.41 DMA scan of HDPE [24].
10000
10000
G’
1000
96.33 °C
0.25
tan
0.20
100
1000
0.15
0.10
100
–118.12 °C
–10.55 °C
-Relaxation
-Relaxation
An amorphous phase relaxation
A local-mode, simple, non-cooperative Originates in amorphous phase
Related to glass transition
relaxation process
–150
–100
–50
0
Temperature (°C)
50
G’’
[ ] Loss modulus (MPa)
Tm
Large scale cooperative motion:
Disruption of crystalline structure
[ + ] Tan delta
[ ] Storage modulus (MPa)
LDPE film
0.05
100
10
150
Universal V2.5D TA instruments
Figure 7.42 DMA scan of LDPE film [22].
Temperature ramps are also useful for indicating the crystallinity of semi-crystalline
materials, degrees of cross-linking, whether blend components are compatible, and the
effects of aging and humidity. For a material like PE, increasing the crystallinity will
increase the transition temperature; it will decrease the intensity of the glass transition; and it will broaden the transition temperature range. There is still discussion
regarding exactly what is occurring on a molecular level when it comes to the transitions of PE. The interested reader can review some articles on that topic.
DMA scans of three types of PE are shown next. They include HDPE, LDPE, and
LLDPE in Figures 7.41, 7.42, and 7.43, respectively. With the HDPE, a limiting storage modulus of about 3300 MPa is observed at the lowest temperatures. Moreover,
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Handbook of Industrial Polyethylene and Technology
300
5
LLDPE
200
3
0.15
E’
tan
Storage Modulus (GPa)
tan
2
0.10
100
’
1
0
–125
E’’
–25
Loss Modulus (MPa)
0.20
4
0.05
–25
25
Temperature (°C)
75
125
0
Figure 7.43 DMA scan of LLDPE [25].
and transitions are observed in the storage modulus and loss modulus curves. Tan
showed a good peak but not much of an peak. The LDPE DMA scan exhibited
and transitions in the storage modulus, the loss modulus, and tan . As expected, the
temperature at which the storage modulus abruptly decreases is lower than that of the
HDPE resin. With LLDPE, transitions at characteristic temperatures are also observed.
In this case, the author proposed that they were observing not only an transition, but
also an ’ transition [25]. It was proposed that there was a transition just below the ’
transition and a low temperature transition that he called the transition. Some of the
assignment of transitions can be subjective.
In all cases, the useful temperature range for a material can be assessed. Also secondary transitions that indicate that the PE materials should have good impact strength
are observed. The transitions would not be observed with a DSC test, but can only be
observed by rheological testing.
In terms of plotting the data, the most common way to display the data is with the
storage and loss moduli plotted on logarithmic scales. Tan is often plotted on a log
scale to be consistent with the other variables, but sometimes the peaks are easier to
discern when tan is plotted on a linear scale.
One thing to mention is that often data from DMA temperature ramps obtained
with film clamps or with flexural clamps can be merged with rheometer data obtained
in parallel plate geometry. The DMA instrument provides access to low temperature
testing. At some point, the specimen will lose its mechanical integrity and not lend
itself to further testing by this method. However, the parallel plate geometry is excellent for testing material that is in the molten state. If the temperature is then decreased
the rheological properties (storage modulus, loss modulus, complex viscosity) will
increase. At some point for this geometry, the specimen will become too stiff for further testing. After both experiments are completed, the data can be merged by taking
the low temperature DMA data and the high temperature parallel plate data and then
one can obtain a complete temperature scan on the material. If the DMA testing is
done in the tensile or flexural mode, the DMA data must be converted to shear data by
Rheology of Polyethylene 279
10
8.944%
Good
Poor
Excellent
Strain (%)
8
5.520%
6
Jer = 2.89E–8
DMA creep conditions
Clamp type: Tension -Film
Temperature = 85°C
Stress level = 1 Mpa
3.466%
4
Jer = 2.43E–8
Jer = 8.86E–9
2
0
0
5
10
15
20
Time (min)
25
30
35
Unversal V2.5H
Figure 7.44 Creep testing on a DMA [27].
400
350
Dimension change ( m)
300
250
206.20°C
188.1 m
200
150
100
50
0
–50
–100
0
50
100
150
Temperature (°C)
200
250
Universal V3.2A TA instruments
Figure 7.45 Controlled force-temperature ramp to simulate heat deflection under load test [29].
using Poisson’s ratio (or the shear data must be converted to DMA data). However, the
tan values should agree even without correction for the differences in geometry. If a
DMA type of scan is done with the torsion rectangular geometry, the storage and loss
modulus values should be fairly close and should match up well with parallel plate data.
7.3.4 Other Tests on a DMA
In addition to temperature scans, DMA instruments can be used for other tests. These
include creep, stress relaxation, controlled force-temperature ramp, controlled straintemperature ramp, and stress-strain types of tests [26]. An example of creep testing is
given in Figure 7.44.
Another test shown here is the controlled force-temperature ramp. This is meant to
simulate the heat deflection temperature under load test (ASTM D648) [28], as shown
by Figure 7.45.
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Handbook of Industrial Polyethylene and Technology
These are representative tests that can be done on a DMA. There are other accessories and options that are available and new accessories and methods that are being
developed. Some additional accessories that are currently available include submersion
clamps and humidity control.
7.4 Conclusions
It is hoped that the description of the rheological testing, on both the melts and solids,
has provided the reader with information on how to characterize PE resins to ensure
that the resins process well and the products meet the performance expectations in the
field.
References
1. Dealy, J.M., and Wissbrun, K.F., Melt Rheology and Its Role in Plastics Processing: Theory and
Applications, Van Nostrand Reinhold, New York, NY, 1990.
2. Dealy, J.M., and Wang, J., Melt Rheology and Its Role in the Plastics Industry, Springer
Publishers, New York, NY, 2013.
3. Ferry, J.D., Viscoelastic Properties of Polymers, 3rd ed., J. Wiley, New York, NY, 1980.
4. Macosko, C.W., Rheology: Principles, Measurements, & Applications, Wiley-VCH, New
York, NY, 1994.
5. Ahirwal, D., Filipe, S., Neuhaus, I., Busch, M., Schlatter, G., and Wilhelm. M., Large
Amplitude Oscillatory Shear and Uniaxial Extensional Rheology of Blends from Linear and
Long-Chain Branched Polyethylene and Polypropylene, J. Rheol., 58, 635 (2014).
6. Hoyle, D.M., Auhl, D., Harlen, O.G., Barroso, V.C., Wilhelm, M., and McLeish. T.C.B.,
Large Amplitude Oscillatory Shear and Fourier Transform Rheology Analysis of Branched
Polymer Melts, J. Rheol., 58, 969 (2014).
7. Habas-Ulloa, A., Moraes D’Almeida, J.R., and Habas, J.P., Creep Behavior of High Density
Polyethylene After Aging in Contact with Different Oil Derivates, Polym. Eng. Sci., 50, 2122
(2010).
8. ASTM D1238–13, Melt Flow Rates of Thermoplastics by Extrusion Plastometer, ASTM
Standard, ASTM International, West Conshohocken, PA, 2013.
9. Lue, C.-T., Estimate of ASTM Melt Flow Rates from Oscillatory Shear Rheology, SPEANTEC Tech. Papers, 61, 87 (2015).
10. ASTM D4440–15, Standard Test Method for Plastics: Dynamic Mechanical Properties Melt
Rheology, ASTM Standard, ASTM International, West Conshohocken, PA, 2015.
11. Garofalo, E., Incarnato, L., and Di Maio, L., Effect of Short-Chain Branching on Melt
Fracture Behavior of Metallocene and Conventional Poly(ethylene/α-olefin) Copolymers,
Polym. Eng. Sci., 52, 1968 (2012).
12. Thimm, W.B., Friedrich, C., Marth, M., and Honerkamp, J., An Analytical Relation Between
Relaxation Time Spectrum and Molecular Weight Distribution, J. Rheol., 43, 1663 (1999).
13. Mavridis, H., and Shroff, R., Appraisal of a Molecular Weight Distribution-to-Rheology
Conversion Scheme for Linear Polyethylenes, J. Appl. Polym. Sci., 49, 299 (1993).
14. Teh, J., Lin, W., and Popuri, P., Determination of the Zero Shear Viscosity of Polyethylene,
SPE-ANTEC Tech. Papers, 61, (2015)
15. Xpansion Instruments, http://www.xinst.com/results_rheology.htm, 2015
Rheology of Polyethylene 281
16. TA Instruments, Extensional Viscosity of Polyolefins, TA Instruments Application Brief
AAN020, 2004.
17. Ansari, M., Zisis, T., Hatziriakos, S.G., and Mitsoulis, E., Capillary Flow of Low-Density
Polyethylene, Polym. Eng. Sci., 52, 649 (2012).
18. ASTM D3835–08, Standard Test Method for Determination of Properties of Polymeric
Materials by Means of a Capillary Rheometer, ASTM Standard, ASTM International, West
Conshohocken, PA, 2015.
19. Goettfert, http://www.goettfert.com/images/stories/download/produkte/Rheotens_71–97en.pdf
20. Mavridis, H., and Shroff, R.N., Temperature Dependence of Polyolefin Melt Rheology,
Polym. Eng. Sci., 32, 1778 (1992).
21. TA Instruments, http://www.tainstruments.com/pdf/brochure/dma.pdf.
22. Turi, E., Thermal Characterization of Polymer Materials, 2nd ed., vol. 1, Academic Press,
Brooklyn, NY, 1997.
23. Murayama, T., Dynamic Mechanical Analysis of Polymeric Materials, Elsevier, New York,
NY, 1978.
24. Sewda, K., and Maiti, S.N., Dynamic Mechanical Properties of High Density Polyethylene
and Teak Wood Flour Composites, Polymer Bulletin, 70, 2657 (2013).
25. Liu, T.M., and Baker, W.E., The Effect of the Length of the Short Chain Branch on the
Impact Properties of Linear Low Density Polyethylene, Polym. Eng. Sci., 32, 944 (1992).
26. TA Instruments, DMA for Heat Deflection Temperature, TA Instruments Application Brief
TA 307, 2003.
27. ASTM D638–14, Standard Test Method for Tensile Properties of Plastics, ASTM Standard,
ASTM International, West Conshohocken, PA, 2014.
28. ASTM D882–12, Standard Test Method for Tensile Properties of Thin Plastic Sheeting,
ASTM Standard, ASTM International, West Conshohocken, PA, 2014.
29. TA Instruments, TA Instruments 2009 DMA Training Course, 2009.
8
Processing-Structure-Property
Relationships in Polyethylene
Rajen M. Patel
The Dow Chemical Company, Freeport, Texas, USA
Contents
8.1 Introduction ............................................................................................................................... 283
8.2 Processing-Structure-Properties Relationship in PE Blown Films..................................... 284
8.3 Processing-Structure-Properties Relationship in PE Cast Films ........................................ 295
8.4 Processing-Structure-Properties Relationship in PE Injection Molding ........................... 296
8.5 Processing-Structure-Properties Relationship in PE Blow Molding .................................. 298
8.6 Processing-Structure-Properties Relationship in PE Fibers and Nonwovens................... 302
8.7 Summary .................................................................................................................................... 304
Acknowledgments .............................................................................................................................. 305
References ............................................................................................................................................ 305
Abstract
Polyethylene (PE) is fabricated into various final end-use parts such as blown film, cast film,
injection molded articles, blow molded bottles, pipes, fibers and nonwovens, and many others. In addition to five primary molecular parameters, the molecular orientation and morphology imparted during fabrication also have a major effect on the mechanical properties of the
final fabricated parts. This chapter will cover processing-structure-properties relationships in
PE, especially focusing on PE blown and cast films, injection and blow molded articles, and
nonwoven fibers.
Keywords: Polyethylene, structure-property relationships, blown film, cast film, LLDPE,
metallocene LLDPE, orientation, WAXD, FTIR, Herman’s orientation factor, draw resonance,
dart impact, Elmendorf tear
8.1 Introduction
Polyethylene, like other semicrystalline polymers, has three basic parameters that
are typically used to characterize its solid state; namely, degree of crystallinity, glass
transition temperature, and melting peak temperature. The melt properties of PE are
Corresponding author: rmpatel@dow.com
Mark A. Spalding and Ananda M. Chatterjee (eds.) Handbook of Industrial Polyethylene and Technology, (283–308)
© 2018 Scrivener Publishing LLC
283
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Handbook of Industrial Polyethylene and Technology
LCB
MWD
SCB
Inter-molecular
SCBD
Mw
+
Additives
Figure 8.1 The five primary molecular parameters of PE.
Primary molecular structure
Analytical techniques
Molecular weights (Mw)
Molecular weight distribution (MWD)
Long chain branching (LCB)
GPC, Rheology, l2
GPC, Rheology, l10/l2
Rheology, l10/l2, GPC-DV, NMR
Short chain branching (SCB) level
Density, DSC, FTIR, NMR
SCB distribution (SCBD)
DSC, ATREF, CRYSTAF
Figure 8.2 Analytical techniques to characterize the five primary molecular parameters of PE.
characterized by shear and elongation flow behavior. However, on a molecular level,
solid and molten state properties of PEs depend on only five primary molecular parameters. These parameters are shown in Figure 8.1. Out of these five primary molecular parameters, molecular weight (Mw), molecular weight distribution (MWD), and
long-chain branching (LCB) affect melt state (rheological) properties. The degree of
short-chain branching (SCB) and short-chain branch distribution (SCBD) do not affect
the melt properties except at high levels of short-chain branching (above about 10 to
12 mole% comonomer). In contrast, all five primary molecular parameters affect solid
state mechanical and thermal properties. Analytical techniques to characterize these
five primary molecular parameters are shown in Figure 8.2
PE is fabricated into various final end-use parts such as blown film, cast film, injection molded articles, blow molded bottles, pipes, fibers and nonwovens, and many others. In addition to the five primary molecular parameters, the molecular orientation
and morphology imparted during fabrication also has a major effect on the mechanical
properties of the final fabricated parts. This chapter will focus on the processing-structure-properties relationships in PE, especially focusing on blown and cast films.
8.2 Processing-Structure-Properties
Relationship in PE Blown Films
The primary process variables in a blown film process include: 1) temperature of the
melt at die exit, 2) die gap, 3) blow up ratio (BUR) defined as the ratio of final film tube
Processing-Structure-Property Relationships in Polyethylene 285
PE blown
films
Die gap, melt temp.
Blow up ratio (BUR)
Draw down ratio (DDR)
Frost line height
Output rate
Processing
Primary
structure
Film
structure
MW, MWD, SCB,
SCBD, LCB.
Properties
Morphology
Orientation
Tensile properties
Dart impact
MD/TD tear
MD/TD shrinkage
Haze & gloss
Figure 8.3 Schematic of processing-structure-properties in PE blown films.
diameter to die diameter, 4) draw down ratio (DDR) defined as the ratio of film velocity at the nip roll to average melt velocity at die exit, 5) frost-line height and cooling
conditions, and 6) rate. DDR is approximately equal to die gap / (film thickness x BUR).
Figure 8.3 shows the schematic of effects of primary molecular structure and processing induced secondary film structure (morphology and molecular orientation) on the
key mechanical and thermal properties of PE blown films.
Note that the deformation and flow of a polymer melt lead to the orientation (alignment due to stretching) of the molecules. Some types of deformation have a greater
ability to orient molecules than others. Deformations that can generate a high degree of
molecular orientation (chain stretching) are called “strong flows,” while those that cannot are called “weak flows.” Extensional (elongational) flow is a strong flow in that it has
a strong tendency to orient (stretch) the molecules in the direction of flow [1, 2]. Simple
shear flow (e.g., flow in the land-length of a die) is a weak flow in that it does not orient
molecules much. Also, whatever little molecular orientation is induced due to shear
flow in a die tends to quickly relax as the melt exits the die (causing die swell). Thus, it is
the post-die extensional flow in blown film, cast film, extrusion coating, melt spinning,
and many other processes that determines the induced final molecular orientation and
subsequent properties. Almost all the orientation induced in a blown film occurs due to
the elongational flow between the die and the frost-line height. Molecular orientation
developed in a film is primarily a function of molecular weight (Mw), MWD, the LCB
level of PE, and the fabrication conditions. The orientation in a machine direction (MD)
is affected by draw down ratio (DDR), strain rate (affected by the time to reach frostline height), and relaxation time (affected by polymer as well as cooling conditions).
In PE blown (and cast) films, the molecular orientation in the crystalline and amorphous phases are usually different and are determined by stresses in the molten polymer just before solidification; i.e., crystallization near the frost-line height. Hence, one
needs to measure orientation in each phase and the total orientation in the final film.
Also, it is necessary to understand molecular orientation in which phase (amorphous
286
Handbook of Industrial Polyethylene and Technology
or crystalline or both) affects film properties. There are various methods to measure
molecular orientation in a PE film. These methods include wide angle x-ray diffraction
(WAXD), Fourier transform infrared (FTIR) spectroscopy, birefringence, shrinkage,
and shrink tension. WAXD measures molecular orientation in the crystalline phase
while shrinkage and shrink tension measure molecular orientation in the amorphous
phase. FTIR can measure molecular orientation in both the amorphous and crystalline phase depending upon peak selections. Birefringence measures total (crystalline
and amorphous phase) molecular orientation in PE films. Film properties are affected
by primary molecular structure and the orientation-morphology developed as a result
of fabrication. Two different structural models have been proposed for PE extruded
films: row-nucleated structure [3] and a-axis structure [4]. The row-nucleated model
of Keller and Machin has been widely adopted to understand the orientation of PE
blown films [3, 5, 6]. Figure 8.4 shows normalized azimuthal intensity curves of (110),
(200), and (020) reflections indicating typical oriented WAXD patterns of shish-kebab
row-nucleated structures. According to the Keller and Machin row-nucleated model,
twisted lamellae are observed under low stress conditions and untwisted lamellae are
observed under high stress conditions.
Patel et al. used a design of experiment (DOE) to determine the effect of blown film
fabrication variables on the structure and properties of Ziegler-Natta (Z-N) catalyzed
LLDPE films [7]. A 3-variable Box-Behnken DOE was conducted to study effects of
die gap, die land length, and blow up ratio (BUR) at constant final film thickness on
key LLDPE blown film properties. Differences in molecular orientation in the films
were studied using optical birefringence and shrinkage methods. The measured key
film properties were correlated to the processing conditions and to measured molecular
orientation. It was shown that the die land length had no effect on film structure and
Figure 8.4 Normalized azimuthal intensity curves of (110), (200), and (020) reflections (a,d) indicating
typical oriented WAXD patterns (b,c) of shish-kebab row-nucleated structures with twisted lamellae –
under low stress conditions (c) and untwisted lamellae – under high stress conditions (f), according to
Keller and Machin [5].
Processing-Structure-Property Relationships in Polyethylene 287
properties. This is due to the fact that shear flow in the die does not orient molecules
much as mentioned above. It was shown that MD shrinkage of LLDPE films is mainly
related to draw down ratio (DDR) and higher DDR gave higher MD shrinkage. All
LLDPE films exhibited about 70 to 80% shrinkage in the MD but expanded in the crossdirection (CD). The expansion was caused by the very low orientation in CD due to
lack of long-chain branching in the Z-N LLDPE. Most films exhibited negative in-plane
birefringence. MD Elmendorf tear strength was found to be inversely related to DDR
and MD shrinkage, suggesting that MD tear strength is dependent primarily on amorphous chain extension and hence amorphous segments orientation in LLDPE blown
films. Dart impact strength of the films was shown to be related to MD shrinkage and
to the induced varying surface roughness due to varying die gaps.
Patel et al. also studied blown film molecular orientation and mechanical properties
(MD/TD tear and dart impact strength) of various LLDPE resins (both Z-N and metallocene catalyzed) under low and high MD orientation conditions [8]. Blown films under
the low MD orientation conditions were made at resin melt temperatures between 248
and 259 °C depending on resin type, a rate of 52 kg/h (15.2 cm die), 1.02 mm die gap,
and a 3.5 BUR at 25 m thickness. Blown films under the high MD orientation conditions were made at melt temperatures between 204 and 231 °C depending on resin type,
a rate of 85 kg/h (15.2 cm die), 2.79 mm die gap, and a 2.0 BUR at 25 mm thickness.
Frost-line height was kept at 70 cm in both cases using 10 °C external cooling air without internal bubble cooling. The film molecular orientation was characterized using
WAXD, FTIR, shrinkage, and shrink tension methods. The absorbances of the –CH2
rock bands at 719 cm 1 and 730 cm 1 were used to estimate the orientation of the crystal
b-axis and a-axis, respectively. The amorphous band at 1368 cm 1 was used to estimate
the orientation of the amorphous phase. Shrinkage was measured using ASTM D2732
(hot-oil bath) using an immersion time of about 20 s. Shrink tension was measured
using ASTM D2838 at 140 °C using four ply film to obtain measurable shrink force.
Table 8.1 shows the polymers used in the study. LLDPE resins A through D were
Ziegler-Natta (Z-N) catalyzed ethylene-octene copolymers. Metallocene LLDPE
Table 8.1 Polymers used in the study [8].
Sample
Melt index*, dg/min
Density, g/cm3
MWD
LLDPE A
1
0.921
3.8
LLDPE B
0.83
0.917
3.8
LLDPE C
0.63
0.917
4.1
LLDPE D
1.1
0.940
3.6
mLLDPE E
0.84
0.921
3.2
mLLDPE F
1.1
0.918
2.5
mLLDPE G
0.94
0.917
3.8
mLLDPE H
1.2
0.940
3.2
*190 °C and 2.16 kg
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Handbook of Industrial Polyethylene and Technology
(mLLDPE) resins E through H comprised single-site catalyzed ethylene-alpha-olefin
copolymers. Table 8.2 summarizes the characterization of molecular orientation via
WAXD, FTIR, shrinkage, and shrink tension methods. Measured Elmendorf tear
strength and dart impact strength properties of high and low MD orientation blown
films are shown in Table 8.2.
From the WAXD scattering patterns, the PE film crystalline structure and degree of
crystalline orientation can be measured. Representative (200), (110) and (020) reflections of the orthorhombic PE crystal structure are shown in Figure 8.5. The degree of
crystalline orientation imparted during the film blowing process, under high and low
orientation conditions, can be seen in Figure 8.6 for the LLDPE A resin. The crystal
a-axis orientation in the MD is clearly seen in LLDPE A resin at the high MD orientation conditions, representative of low stress row-nucleated twisted lamellae via the
Keller-Machin morphology. Figure 8.7 shows WAXD 2-D pattern for mLLDPE F blown
film at the high and low MD orientation conditions. The blown film mLLDPE F resin
shows qualitatively a very low degree of a-axis orientation even under the high MD orientation conditions, representative of low stress row-nucleated twisted lamellae via the
Keller-Machin morphology. This is due to its very narrow MWD and lack of long-chain
branching, leading to fast melt relaxation characteristics during film blowing resulting
in a low stress level in the melt just before crystallization. Blown films of all the resins
exhibited b-axis (020) (crystal growth axis) oriented preferentially towards the crossdirection as expected. This is shown by the negative Herman’s orientation factor (factor
of – 0.5 indicates total molecular alignment in the cross-direction).
Figure 8.8 shows a plot of the a-axis (200) Herman’s orientation factor measured
using FTIR versus WAXD. It can be seen that excellent correlation was obtained
between the two techniques for measuring a-axis orientation in PE blown films. Note
that a-axis orientation factors measured using FTIR were always lower than that measured using WAXD. This could be due to the two techniques probing slightly different
structural moieties for measuring the a-axis orientation. Measurement (or quantification) of the amorphous orientation in the blown films was attempted using FTIR.
However, as shown in Table 8.2, the amorphous Herman’s orientation factor measured
using 1368 cm-1 band was very close to zero due to a low orientation level. Traditional
FTIR is not sensitive enough to measure and differentiate such low levels of amorphous
orientation.
Amorphous orientation in blown and cast films can be measured using shrinkage and
shrink tension. The MD shrink tension for the high orientation blown films of LLDPE A
through C (all about 0.92 g/cm3 density but with varying melt index) increased significantly with decreasing melt index (higher weight average molecular weight), as shown
in Table 8.2 and Figure 8.9. This is because amorphous orientation in LLDPE blown
films would be directly proportional to relaxation time which is related to the weight
average molecular weight. The measured MD shrink tension was also dependent upon
MWD of the resin at a given melt index. For example, the mLLDPE F high orientation
blown film exhibited lowest MD shrink tension due to its very narrow molecular weight
distribution. For blown film made under the low orientation conditions, the MD shrink
force was too low to measure even using four plies of film. For LLDPE blown films, CD
shrink force was too low to measure in general. The MD shrink tension results suggested that shrink tension differentiated amorphous orientation in the blown films very
High
Low
High
Low
High
Low
High
Low
High
Low
High
Low
High
Low
High
Low
Sample
LLDPE A
LLDPE A
LLDPE B
LLDPE B
LLDPE C
LLDPE C
LLDPE D
LLDPE D
mLLDPE E
mLLDPE E
mLLDPE F
mLLDPE F
mLLDPE G
mLLDPE G
mLLDPE H
mLLDPE H
NMF – No measurable force
MD
orientation
0.13
0.28
0.09
0.36
0.01
0.16
0.18
0.34
0.24
0.29
0.18
0.36
0.15
0.30
0.19
0.28
WAXS
(200)
0.24
0.38
0.13
0.40
0.04
0.22
0.33
0.42
0.33
0.41
0.34
0.43
0.35
0.42
0.34
0.40
WAXS
(020)
0.14
0.16
0.11
0.17
0.22
0.03
0.07
0.21
0.15
0.00
0.10
0.14
0.18
0.01
0.10
0.12
0.21
0.18
0.19
0.20
0.1
0.16
0.25
0.11
0.14
0.18
0.20
0.09
0.13
0.19
FTIR
b-axis
0.12
0.20
FTIR
a-axis
0.02
0.01
0.03
0.01
0.01
0.01
0.00
0.01
0.01
0.00
FTIR
amorphous
54
71
26
43
48
75
62
77
54
75
51
72
Shrink
MD,
%
NMF
30.5
NMF
9.7
NMF
17.6
6.5
52.4
NMF
32.6
NMF
19.5
Shrink tension
MD, psi
64
45
416
388
312
285
377
229
40
31
444
505
380
484
300
Tear, MD,
g/mil
227
1014
483
656
360
445
522
716
131
627
529
785
599
810
594
818
Tear, CD,
g/mil
Table 8.2 Molecular orientation, Elmendorf tear, and dart impact (method A) of the high and low MD orientation blown films [8].
120
60
835
276
835
717
835
312
65
60
835
504
513
267
261
144
Dart,
g
Processing-Structure-Property Relationships in Polyethylene 289
290
Handbook of Industrial Polyethylene and Technology
200
110
MD
020
Figure 8.5 Representative WAXD pattern for high orientation LLDPE with (200), (110) and (020)
reflections marked [8].
High
Low
MD
Figure 8.6 WAXD patterns for LLDPE A under high and low MD orientation conditions [8].
High
Low
MD
Figure 8.7 WAXD pattern for mLLDPE F under high and low MD orientation conditions [8].
effectively compared to the FTIR method discussed above. The measured MD shrink
tension was a function of melt index and MWD for the LLDPE resins studies.
Shrinkage of blown film can also be used to qualitatively understand amorphous orientation. MD shrinkage was lower for the low orientation blown films compared to the
high orientation blown films, as expected. MD shrinkage of the LLDPE blown films at
a-axis Hermans orientation factor (FTIR)
Processing-Structure-Property Relationships in Polyethylene 291
0.4
0.35
0.3
0.25
0.2
0.15
0.1
0.05
0
0
0.05 0.1 0.15 0.2 0.25 0.3 0.35
a-axis Hermans orientation factor (WAXD)
0.4
MD shrinkage (%) and shrink tension (psi)
Figure 8.8 Comparison of the a-axis Herman’s orientation factor measured using FTIR and WAXD [8].
90
MD shrinkage
80
70
60
50
MD shrink tension
40
30
20
10
0.5
0.6
0.7
0.8
0.9
Melt Index, dg/min
1
1.1
Figure 8.9 MD shrinkage and shrink tension as a function of melt index of 0.920 g/cm3 Z-N LLDPE A
through C blown films (high MD orientation) [8].
the high orientation condition, however, was not a strong function of melt index compared to MD shrink tension as seen from Table 8.2 and Figure 8.9. MD shrinkage was
influenced strongly by MWD. The mLLDPE F resin having narrow MWD exhibited the
lowest MD shrinkage.
Table 8.2 shows that resin comprised of single-site catalyzed ethylene/alpha-olefin
copolymer (E through H), in general, exhibited lower Elmendorf tear strength but
significantly higher dart impact strength compared to standard Z-N LLDPE resins
(A through D).
Overall, measurement of film molecular orientation helps to understand mechanical properties, especially tear strength and dart impact strength of blown films as a
function of processing conditions. However, correlating orientation measurements to
film properties of different PE resins is always challenging as resin intrinsic properties would be different along with differences in orientation (due to differences in melt
relaxation times) in the final blown films.
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Handbook of Industrial Polyethylene and Technology
There have been many other studies done to understand the effect of intrinsic resin
molecular structure and resulting molecular orientation in blown films on the film
properties [9]. Sukhadia has studied the blown films of resins made using Z-N, chromium, and metallocene catalysts [10, 11]. He showed that LLDPE resins made from
these catalysts are dramatically different in nature, especially in MWD, thereby resulting
in different trade-offs in their processing-structure-properties relationships in blown
films. The LLDPE resins made from chromium catalyst had very broad MWD and high
zero shear viscosities, melt relaxation times, and elongational viscosities. The LLDPE
resin from Ziegler-Natta (Z-N) catalyst was intermediate while the LLDPE resin from
the metallocene catalyst was the lowest in terms of the same attributes. The chromium
catalyzed resin exhibited excellent bubble stability and rates. However, the resulting
blown films exhibited low dart impact strengths and MD tear strengths, especially at
lower film gauges, as well as poor optical properties (high % haze). Sukhadia concluded
that chromium catalyzed resins are most suitable for non-clarity thicker film gauge
applications where high rates can be achieved due to relatively higher melt strength
resulting from a very broad MWD. Metallocene catalyzed LLDPE resins exhibited lowest bubble stability and rates at high motor loads. Blown film from the resin, however,
exhibited best optics as evidenced by very low % haze and the best impact strength. The
impact and tear strengths of metallocene catalyzed blown LLDPE resins were shown to
depend primarily upon resin density, increasing as resin density decreased. Z-N catalyzed LLDPE resin exhibited intermediate processability and rates, and the blown film
exhibited intermediate optics, impact strength but best MD and transverse direction
(TD) Elmendorf tear strength among the three resins compared.
Krishnaswamy and Sukhadia [12] have reported on the influence of lamellar morphology and orientation on dart impact of LLDPE blown films. They showed that a
lower degree of lamellar orientation (more random orientation) resulted in higher
impact strengths for blown films. Krishnaswamy and Sukhadia also studied morphological origin of Elmendorf tear of blown films. The orientation features of the blown
films were described in terms of Keller-Machin “row” structures [3]. Investigation of
the crystalline phase orientation in the LLDPE blown films indicated preferential orientation of the unit cell a-axis along the film MD (Keller–Machin-I “row” structure)
with the lamellar long axis located along the TD-normal plane in all films. The authors
reported orientation in the non-crystalline phase to be substantial, although the degree
of orientation was not as high as in the crystalline phase. MD Elmendorf tear of LLDPE
blown films was observed to be dependent on the process extension rate imposed along
the MD. Lower MD extension rates (low Deborah numbers) were observed to favor
MD tear performance, while higher MD extension rates (high Deborah numbers) were
observed to favor TD tear performance.
Guichon et al. studied the blown film structure and properties of three types of PE
resins; namely high pressure LDPE, Ziegler-Natta catalyzed LLDPE, and metallocene
catalyzed LLDPE (mLLDPE) [13]. The blown films displayed a decreasing puncture
and impact resistance in the order mLLDPE, Z-N LLDPE, and LDPE. In parallel, the
tear resistance of the films became increasingly unbalanced in the same order of the
polymers. The morphological study showed an increased anisotropy of the films in the
same polymer order, the crystalline lamellae being increasingly oriented normal to the
take-up direction. This was attributed to the kinetics of chain relaxation in the melt that
Processing-Structure-Property Relationships in Polyethylene 293
governed the orientation in the amorphous phase due to the elongational flow-induced
stretching before crystallization. They concluded that the average relaxation time in the
melt is a major factor for the development of morphology, orientation, and mechanical
anisotropy.
Zhang et al. studied blown film properties of LDPE and LLDPE (Z-N and singlesite catalyzed) resins, especially focusing on MD/TD tear strength ratios [14]. They
postulated that the different rheological response of LDPE versus LLDPE, especially in
terms of relaxation time may play an important role in the establishment of the crystalline morphology. They showed that LDPE blown film developed a row-nucleated
structure, while a spherulite-like and/or random lamellar superstructure was observed
for LLDPE blown film. These morphological differences were shown to translate into
different ratios of machine and transverse direction tear strengths in the blown films.
Zhang et al. also showed that the morphologies and orientation for LDPE, LLDPE and
HDPE blown films were related to the stress applied (related to DDR) and their relaxations times in the flow induced crystallization process, which determine the amount of
fibrillar nuclei available at the time of crystallization and therefore the final crystalline
morphology [15]. These structure differences were shown to translate into different
ratios of machine and transverse direction tear and tensile strengths.
Savargaonkar et al. studied abuse performance fundamentals of blown films made
from Z-N LLDPE/LDPE blends [16]. They showed melt strength synergy [17] with
melt strength going through a maximum, as shown by Figure 8.10. MD Elmendorf
tear strengths blown films from LLDPE/LDPE blends provides a minimum, as shown
by Figure 8.11. The minimum in MD Elmendorf tear strength was observed at a LDPE
level between 40 and 80%. This observation was found to be related to the MD orientation effects in a blown film and seemed to be independent of the type of LDPE used in
these blends. The intrinsic tear strength of the LLDPE/LDPE blends showed a continuous decrease with increasing LDPE concentration for all three LDPE resins. This was
hypothesized to be due to decreasing tie-chain concentration with increasing LDPE
content of the blends. Dart impact strength of the blown films showed a continuous
decrease with increasing LDPE level for all three LDPE resins studied. Similar trends
were observed for the decrease in intrinsic dart impact strength with increasing LDPE
25
LDPE-B
LDPE-A
Melt strength (cN)
20
LDPE-C
15
10
5
Resin
MI, dg/min
Density, g/cm3
LLDPE
0.95
0.920
LDPE-A
0.31
0.925
LDPE-B
0.64
0.921
LDPE-C
1.63
0.919
0
0
10
20
30
40
50
60
70
80
90 100
% LDPE
Figure 8.10 Melt strength synergy of LDPE and LLDPE resins at 190 °C [16].
294
Handbook of Industrial Polyethylene and Technology
1000
900
MD tear strength (g)
800
700
LDPE-C
600
Density, g/cm3
Resin
MI, dg/min
LLDPE
0.95
0.920
LDPE-A
0.31
0.925
LDPE-B
0.64
0.921
LDPE-C
1.63
0.919
500
400
300
LDPE-B
200
LDPE-A
100
0
0
10
20
30
40
50
60
70
80
90 100
% LDPE
Figure 8.11 MD Elmendorf tear strength of 50 m blown films of LLDPE/LDPE blends [16].
content of the blends. The decrease in the dart impact and MD Elmendorf tear of the
LLDPE/LDPE films with increasing LDPE content was thought to be primarily occurring due to increased MD orientation in the films and secondarily due to decreased
tie-chain concentration.
Lu et al. studied the morphological features in the LDPE, Z-N LLDPE, and a blend
of LDPE-LLDPE (50/50) in blown films [18]. They showed that under similar processing conditions, the LLDPE film had a relatively random crystal orientation. The LDPE/
LLDPE blend film exhibited the highest crystal orientation. The crystal orientation of
the LDPE film was shown to be intermediate between the two. The LLDPE film exhibited a random amorphous phase orientation while the LDPE film exhibited the highest
amorphous phase orientation. Co-crystallization between LDPE and LLDPE was also
observed.
Lee et al. examined LDPE, HDPE, and LLDPE blown films with similar crystal orientations, as verified using through-film X-ray scattering measurements [19]. With
these common orientations, LDPE and HDPE films still followed the usual preferred
tear directions (TD for LDPE and MD for HDPE), whereas LLDPE exhibited isotropic
tear despite an oriented crystal structure. These differences were attributed to the number densities of the tie chains, especially along MD, which are considerably greater for
linear-architecture polymers.
Elkoun et al. [20] have investigated the effect of processing parameters on the enduse properties of mono – and five-layer coextruded PE blown films using three different LLDPE resins (Z-N 1-butene LLDPE, Z-N 1-octene LLDPE and single-site
catalyzed 1-octene LLDPE). It was found that within the investigated range of processing conditions, tear strength increased in the direction perpendicular to the highest
orientation, impact and puncture strength increased with the overall orientation, and
the effect on orientation due to shear stresses in the die was negligible because of rapid
macromolecular relaxation before crystallization. Note that this is expected as shear
flow is a weak flow as mentioned earlier. Finally, it was also shown that due to changes
in the size of the crystallites, haze increased with die gap and frost-line height (FLH),
and decreased as take-up speed increased. Elkoun et al. [21] have also investigated the
Processing-Structure-Property Relationships in Polyethylene 295
performance of multilayer coextruded LLDPE blown films. The co-extruded structures
composed of the Z-N butene-LLDPE and the single-site catalyzed octene LLDPE resin
yielded improved end-use properties relative to the monolayer blended films. This
result was ascribed to the presence of interfacial trans-crystalline layers. Also, blends of
the single-site LLDPE and the advanced Z-N LLDPE octene resins within selected layers of coextruded films showed slightly enhanced tear resistance. It was also found that
haze was significantly reduced when the outside layers were composed of the single-site
catalyzed LLDPE resin.
Gupta et al. [12] studied the effect of short-chain branch length (comonomer type) on
the properties of blown and compression molded films of ethylene/alpha-olefin copolymers (1-butene, 1-hexene or 1-octene comonomer based) at a solid density of about
0.918 g/cm3. The tensile properties of the films did not show any significant difference
at slow deformation rates (up to 510 mm/min). However, at higher deformation rates
(about 1 m/s), the tensile strength of these films was found to increase with increasing short-chain branch length. In addition, the Spencer impact and Elmendorf tear
strength of the blown films were also observed to increase with increasing short-chain
branch length. Dart impact strength and high-speed puncture resistance (5.1 m/s) of
1-octene – and 1-hexene-based films was also observed to be higher than that based on
1-butene comonomer. The authors concluded that while the origin of these differences
in mechanical properties with increasing short-chain branch length is not fully understood, their investigation confirmed this effect to be pronounced at high deformation
rates for both the blown and compression molded quenched films.
8.3 Processing-Structure-Properties
Relationship in PE Cast Films
DeGroot et al. [22] used a design of experiment (DOE) to determine the effect of cast
film fabrication variables on the performance of LLDPE stretch films. A three variable
Box-Behnken designed experiment was conducted to study the effects of air gap, melt
temperature, and line speed on the key cast stretch film properties. In addition, the differences in molecular orientation in the films were studied using optical birefringence
and shrinkage methods. The key film properties were correlated with the fabrication
conditions using a statistical analysis program. A higher melt temperature, larger air
gap, and slower line speed resulted in greater film extensibility (ultimate stretch) and
lower load retention. MD shrinkage and optical birefringence measurements correlated very well with the ultimate stretch and the load retention of the film samples. The
results of this study were explained in terms of web tensile stresses before solidification
and the degree of amorphous orientation developed in the film due to the stresses.
Dohrer et al. [23] demonstrated that at constant molecular weight, the onset of draw
resonance in the cast extrusion of LLDPE is substantially influenced by MWD. The
greatest haul-off line speeds were shown to be achieved using LLDPE copolymers with
the narrowest MWDs. The film tensile strengths of the LLDPE were shown to be superior to those of LDPE. They concluded that the use of narrow-MWD LLDPE in conjunction with improved extrusion practices, such as a draw resonance eliminator [24],
should result in even greater line speeds.
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Handbook of Industrial Polyethylene and Technology
8.4 Processing-Structure-Properties
Relationship in PE Injection Molding
Injection molding is a process where a solid thermoplastic material is heated and
melted, then injected (under pressure) into a mold which is the inverse of the product’s
shape. The material is cooled until it solidifies, conforming to the cavity shape. The
injection molding process is detailed in Chapter 15. During the process, the polymer
experiences a very complex thermo-mechanical history [25]. This history is affected by
molding conditions such as melt temperature, mold temperature, injection rate, cavity
dimensions, and material properties, and it introduces into the injected parts molecular
orientation, shrinkage, and residual stresses resulting in warpage and highly anisotropic
mechanical properties. An injection molding machine has two basic functions: first, it
must melt the polymer, and inject the molten plastic into a closed mold using pressure;
and second, it must generate enough clamping force to hold the mold closed against the
amount of injection pressure being used. An injection molding cycle consists of injecting and filling about 95% of the mold as quickly as possible (injecting), and then slowly
adding more molten material in the mold to compensate for shrinkage (packing), keep
applying pressure until the gate freezes (hold), cool the polymer until the part can be
ejected (cool), and prepare molten material for the next shot (plasticize). Key properties required by molders include processability, stiffness, flexibility, tensile properties,
environmental stress crack resistance (ESCR), impact resistance, and tear resistance.
Key factors affecting mold cycle include PE melt index (most important), molecular
weight distribution, density, melt temperature, mold temperature, injection pressure,
and injection speed. PE with a higher melt index flows easier for better mold fill-out,
leading to reduced cycle times and higher gloss for the final injection molded part. PE
resins having a broader molecular weight distribution flow easier (lower viscosity) at
high injection speeds (shear rates), leading to reduced cycle times. Higher PE melt temperature leads to a better mold fill-out, improved part toughness, and improved ESCR.
However, it leads to a longer molding cycle due to an increased cooling time. Higher
mold temperature increases cycle time and reduces part toughness and ESCR properties. Higher injection pressure leads to more packing thereby reducing part shrinkage.
Injection pressure is only effective until the gate freezes.
During injection molding, the PE melt is subjected to a high shear flow prior to crystallization (solidification) of the final part, leading to molecular orientation (stretching)
in the direction of flow along the flow path. The resulting part morphology is quite different from what is observed under quiescent crystallization conditions. Typically, oriented crystalline structures like “shish-kebabs” or lamellae row structures are formed
if flow and cooling rates are large enough. This is generally the case near the cold walls
of the injection mold.
Injection-molded samples of semicrystalline polymers are known to show an inhomogeneous structure over both the thickness as well as the length of the sample. The
molecular orientation in the skin layer of the injection-molded products is due to the
fountain-like flow at the melt front which orients the polymer molecules in the flow
direction parallel to the mold surface [25]. In the center of the molding, however, it has
been found that the shear flow dominates in orienting the polymer molecules in the
Processing-Structure-Property Relationships in Polyethylene 297
flow direction. Therefore, during the filling stage, both shear and extensional stresses
are present in the polymer melt [25]. The variation of the intensity of the shear and
elongational stresses will influence the degree of molecular orientation in the injected
part. Generally, a higher injection flow rate results in a higher molecular orientation
while a higher melt temperature reduces the molecular orientation due to lower viscosity and the molecular relaxation after the cessation of the flow. Morphology and thickness of the skin layers vary with the molding conditions applied: initial flow rate, melt,
and mold temperature. The morphology of these skin layers varies with polymer melt
properties and molding conditions, but also along the flow path of the molded parts.
The inhomogeneity of such parts affects final physical properties such as dimension stability, Young’s modulus, and tensile strength. A low amount of orientation, often present
in thick molded samples, does not have a significant effect on the mechanical behavior
[5]. However, if oriented structures percolate through the sample, like observed in thinwalled injection-molded parts, the anisotropic nature dominates mechanical properties; e.g., tensile and impact behavior.
Schrauwen et al. studied the flow-induced oriented structures in injection molded
HDPE samples using optical microscopy, X-ray scattering techniques, and transmission electron microscopy [5]. The optical micrographs showed layers of different levels of orientation. Layer thicknesses in the injection molded samples were found to
decrease along the flow path and increase for lower melt temperatures and longer filling
times. Wide-angle X-ray diffraction patterns taken along the thickness of the sample
indicated the existence of several layers composed of orientation varying from lamellae
row structures to highly oriented shish-kebabs. From tensile tests, performed in and
perpendicular to the flow direction, the yield stress was found to increase in the flow
direction due to the amount of oriented extended chain crystals (shish), whereas strain
hardening was assumed to be mainly increased by chain orientation forming oriented
lamellae (kebabs).
Figure 8.12 shows optical micrographs of the complete cross-section of an injection
molded HDPE (8 MI, 0.963 g/cc, Mw=70,000, Mn=11,000) part taken from a position
close to the injection molding gate a), in the middle of the sample (b), and far from the
A
B
C
F
D
E
(a)
(b)
(c)
200 m
Figure 8.12 Optical micrographs of cross-sections of an injection molded HDPE part (8 dg/min MI,
0.963 g/cm3) at different locations: (a) a position close to the gate, (b) middle of the part, and c) far from
the gate viewed between cross-polars at 45° with the flow direction, showing difference in thickness of
oriented skin layer. The layers are “skin layer” (A), “transition layer” (B), “shear layer” (C), “fine-grained
layer” (D), “isotropic core” (E), and “fiberlike layer” (F) [5].
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Handbook of Industrial Polyethylene and Technology
gate (c). At all positions, several structural layers can be distinguished: a thin “skin layer”
(A), a “transition layer” (B), a “shear layer” (C), a “fine-grained layer” (D), and an “isotropic core” (E). At the positions near the gate (a) and in the middle (b) of the part the “shear
layer” shows a very highly oriented “fiber like layer” (F) which is partly delaminated
upon microtoming. The total thickness of the oriented layers (A,B,C, and F) is largest at
the position close to the gate (a) and on the order of 350 m, decreases a to approximately
300 m in the middle of the sample at (b), and is clearly smallest at the end position (c)
at about 200 m. Reduction of the oriented layer thicknesses at increasing distances from
the injection gate was observed for all injection-molded samples. However, the thickness
of the oriented layers strongly depended on molecular weight and molding conditions. A
higher melt temperature and higher melt index led to a lower melt viscosity, decreasing
the thickness of the oriented layers. The higher flow rate (injection speed) increased the
shear stresses in the polymer, resulting in higher orientation level.
Schrauwen et al. studied the relation between the impact toughness and flow induced
crystalline orientation of injection molded HDPE (8 dg/min MI, 0.963 g/cm3) with and
without calcium carbonate filler [26]. The amount of orientation was controlled via
processing conditions and part thickness. Impact toughness was found to be strongly
dependent upon the level of crystalline orientation and loading direction (highest
impact properties in flow direction). Sousa et al. investigated structure-properties relationships in shear-controlled orientation in injection molding (SCORIM) of HDPE
[27]. SCORIM was used to deliberately induce a strong anisotropic character in the
HDPE microstructure.
8.5 Processing-Structure-Properties
Relationship in PE Blow Molding
Blow Molding is a process suited to fabricate hollow parts with the process dating
back to the earliest days of glass blowing. Blow molded HDPE containers gained full
acceptance in the market place with unprecedented growth in the 1960s after commercial introduction of HDPE resins in the late 1950s. Coextrusion and the manufacture of large parts have driven the continued growth of the blow molding process.
Blow molding process trends include multilayer blow molding for improved properties
such as barrier constructions for food and beverage long shelf-life packaging, agricultural chemical packaging, “soft-touch” outer layers, and enhanced cooling technologies for shorter cycle times. Blow molding market trends include increasing the use of
post-consumer recycled content in a bottle. Common blow molding processes include
continuous wheel, continuous shuttle, intermittent reciprocating, accumulator head,
injection blow, and injection stretch blow molding. Blow molding applications include
household and industrial chemicals (HIC) bottles (detergent, bleach, fabric softeners,
household products, antifreeze, motor oils, and agricultural chemicals such as pesticides and herbicides), food bottles (dairy, water, juice, condiments, and oils), pharmaceuticals and health care bottles (personal care, cosmetics, vitamins, and medicines),
and large parts and drums (automotive fuel tanks, 55 gallon drums, and industrial bulk
containers in excess of 300 gallons). HDPE resins used in the blow molding process
are made using slurry and gas phase or hybrid processes. HDPE homopolymers have
Processing-Structure-Property Relationships in Polyethylene 299
superior performance (stiffness, taste, and odor) for packaging milk, water, and other
non-aggressive contents. HDPE copolymers perform well in packaging applications
where an optimized balance of environmental stress crack resistance (ESCR), toughness, and stiffness are required (for chemicals, detergents, and shampoo packaging).
LLDPE/MDPE (0.92 to 0.94 g/cm3) resins are used typically to make small squeezable
bottles due to their lower stiffness. In North America, the majority of HDPE blow molding volume is produced via chromium oxide catalyst systems (unimodal resins). Such
HDPE resins have a very broad MWD, and blow molding dies are currently designed
to match the rheology of chromium oxide resins.
Key processing requirements for blow molded bottles include thickness and diameter swell control, melt strength for parison sag resistance to maintain good part thickness distribution, processability for high extrusion rates and low cycle times, and no
melt fracture at high shear rates. The thickness swell is the increase in parison thickness
immediately after leaving the die, so that parison wall thickness is greater than the
actual die gap. The diameter swell is the increase in parison diameter immediately after
leaving the die, so that parison diameter is greater than the actual die diameter. The
most critical stage in the extrusion blow-molding process is the parison formation as
the dimensions of the blow-molded part are directly related to the parison dimensions.
The swelling due to stress relaxation and sagging due to gravity are strongly influenced
by the resin characteristics, die geometry, and operating conditions. These factors significantly affect the parison dimensions. This could lead to a considerable amount of
time and cost through trial and error experiments to get the desired parison dimensions based upon variations in the resin characteristics, die geometry, and operating
conditions [28, 29]. Eggen and Sommerfeldt developed an image analysis technique to
measure diameter and thickness distribution of a parison during the extrusion stage
in blow molding [30]. They showed that swell and sag related properties during blow
molding could not be inferred from standard laboratory die swell measurements.
Modeling approaches have been used to predict blow molding of HDPE parts
[28–30]. Yousefi et al. demonstrated that a modeling approach capable of creating the
link between the rheological properties of the resin and parison swell and sag could be
the key to a better material design and a better control over the thickness distribution
for the final blow molded part [29]. Ariawan et al. studied the influence of molecular
structure on the rheology and processability of twenty four commercial blow-molding
grade HDPE resins in terms of their shear and extensional flow properties, extrudate
swell characteristics, and melt strength [32]. They found that extrudate swell behavior
and melt strength are important parameters to be considered during parison formation, as observed during blow-molding experiments.
Key performance requirements for blow molded bottles include stiffness to minimize
bottle weight, ESCR for resistance to the content of the bottle, impact strength for drop
height resistance, taste-odor for food applications, optics for consumer appeal, chemical resistance, and heat resistance for hot fill applications. In any blow molding process
the key is to balance the physical properties of the blow molded article. These include
stiffness measured as top load, ESCR to that of the packaged content, and impact as
measured in the height from which the blow molded article can be drop without failure. The desire is to make the lightest weight blow molded article without failure due
to top load, ESCR, and toughness. The physical properties of the blow molded article
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Handbook of Industrial Polyethylene and Technology
are dependent upon the processability of the resin. If the polymer does not process well
then the desired performance is unattainable. Proper parison thickness and diameter
swell are key to obtaining the desired physical performance of the blow molded article
as described previously. This is shown in Figure 8.13.
Environmental stress cracking (ESC) is defined as the surface-initiated failure of
polyaxially stressed polymers in the presence of surface-active substances such as
alcohols, soaps, and wetting agents [33]. ESC is a physical phenomenon; the surfaceactive substances do not cause chemical degradation of the polymer, but they accelerate
the process of macroscopic brittle-crack formation. High ESCR is one of the primary
requirements of PE bottles used for household and industrial chemical packaging. In
the presence of the aggressive environment which they contain, these bottles must
withstand low stresses for long periods without cracking. The effect of resin properties on ESCR has been the subject of many investigations. In general, lowering density,
MI (melt index), and HLMI (high load melt index) increase ESCR. Copolymers have
higher ESCR than homopolymers. The formation of tie molecules is essential for high
ESCR. By traversing the amorphous domains and linking adjacent lamellae, these tie
chains retard the growth of inter-lamellar stress cracks. Soares et al. analyzed nineteen
commercial HDPE resins, made with different polymerization processes and catalyst
types, using high-temperature size exclusion chromatography (SEC) and crystallization analysis fractionation (CRYSTAF) to understand their ESCR properties [33]. They
found that ESCR increased when the molecular weight and concentration of polymer
chains that crystallized in trichlorobenzene (TCB) between 75 and 85 °C increased.
They hypothesized that polymer chains present in this crystallization range act as tie
molecules between crystal lamellae. Strebel and Benson studied the effects of blowmolding process variables on the ESCR of HDPE bottles [34]. Experimental design was
employed to test the effect of die temperature, mold temperature, molding time, and
drop time. They showed that increasing die temperature, mold temperature, and molding time decreased ESCR. However, increasing drop time increased ESCR.
Unimodal HDPE blow molding resins are produced using one catalyst (chromium
oxide) in one reactor. The result of this process is a polymer with a reasonably broad
molecular weight distribution and a very small level of long-chain branching. This broad
range of polymer chain sizes includes both smaller molecules, which affect processability (e.g. extrusion flow rates) and stiffness, and much larger molecules, which influence
physical properties such as ESCR and drop impact strength. Density (or crystallinity) is
Impact
Processing
Thickness & diameter swell
Melt strength
Top load
Figure 8.13 Key performance requirements for blow molding.
ESCR
Processing-Structure-Property Relationships in Polyethylene 301
a critical attribute for PE blow molding resins. For a given PE material, reducing density improves many important physical properties related to ductility such as ESCR and
impact strength. Density is controlled by the incorporation of comonomers into the polymer at relatively small levels during polymerization. These comonomers create short side
chain branches, which act to disrupt the crystalline structure and lower density. However,
this process is not completely efficient in unimodal chromium oxide catalyzed resins,
because the comonomer preferentially goes into the smaller, lower molecular weight
chains, which are less effective (than incorporating comonomer in the longer polymer
chains) at the formation of tie-chains and hence at improving physical properties. This
tendency for comonomer incorporation into shorter polymer chains limits the ESCR and
drop impact properties for a unimodal resin at a given density. This limitation is particularly important for blow molding, where top load strength (favored by higher density)
must be balanced with ESCR and drop impact resistance (favored by lower density).
In North America, most HDPE blow molding resins are based on chromium oxide
catalyst using unimodal (single reactor) process technologies and have performed well
for many decades. However, sustainability (light weight containers) and the emergence
of new applications with more severe performance requirements (improved ESCR)
have pushed the performance demands on blow molding resins to new levels. Although
unimodal chromium oxide catalyzed HDPE blow molding resins generally perform
very well in terms of the balance of top load, ESCR, and drop impact strength, bimodal
Ziegler-Natta (Z-N) catalyzed blow molding resins represent a significant improvement in the balance of all of these properties, enabling light weighting of the containers
or incorporation of more post consumer recycle (PCR).
Bimodal Z-N HDPE resins give superior balance of stiffness for top load, impact
strength, and ESCR in the final blow molded part compared to unimodal chromium
oxide resins. This is due to higher tie-chain concentration in the bi-modal resins compared to the unimodal resins at a given density and melt index (processability). Higher
tie-chain concentration is achieved by incorporating alpha-olefin comonomer into very
high molecular weight fraction in one reactor (making a very high molecular weight,
lower density fraction) and incorporating low or preferably no comonomer into the
very low molecular weight fraction in the other reactor (making a very low molecular
weight, high density fraction). This is shown schematically in Figure 8.14. Long molecules having short-chain branches due to comonomer incorporation have a greater
High MW
copolymer
fraction for tie
chains to give
improved ESCR
and impact,
strength
Weight percent
Low MW
homopolymer
fraction for
stiffness and
processability
Comonomer
content (wt%)
Molecular weight
Figure 8.14 Bimodal HDPE resins design principle for enhanced stiffness (top load), impact, and ESCR
balance via optimizing tie-chain concentration.
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Handbook of Industrial Polyethylene and Technology
Bi-modal HDPE materials
High ESCR unimodal product
Standard unimodal HIC products
NCLS, h
100
Bi-Modal HDPE samples at 0.9575
g/cm3 density have ESCR exceeding
that for unimodal resins with
density of 0.949 g/cm3 and lower
10
1
0.948
0.950
0.952
0.954
0.956
0.958
0.960
Density, g/cm3
Figure 8.15 ESCR as a function of density for bimodal and unimodal HDPE resins. ESCR was measured
using the notched constant ligament stress (NCLS) method (ASTM F2136).
chance of becoming crystallized in separate lamellae, and hence forming a tie chain. In
addition, these long molecules have a higher probability of becoming entangled with
other chains, and hence increasing tie-chain concentration. Note that tie-chain concentration is a very useful concept to understand and rationalize mechanical properties
of PEs. However, the tie-chain concentration cannot be quantified experimentally and
theoretical models have been developed to predict relative tie-chain concentration as a
function of molecular weight and alpha-olefin incorporation [35, 36]. Key drawbacks
of bimodal HDPE resins are their lower die swell (affecting bottle thickness, weight,
and diameter) and melt strength (affecting parison sag and bottle thickness variation)
compared to unimodal chromium oxide catalyzed HDPE resins. This makes it quite
challenging to process bimodal HDPE resins on blow molding machines having dies
designed to process unimodal chromium oxide catalyzed resins.
Superior stiffness (controlled by PE density) and ESCR balance of bimodal Z-N
HDPE resin compared to unimodal chromium oxide resins is illustrated in Figure 8.15
below. This can lead to a reduction in bottle weight of 10 to 15% while retaining ESCR
and bottle drop resistance. Attempts have been undertaken to produce such bimodal
resins in a single reactor, using a bimodal catalyst system (e.g., a dual site catalyst or
bimetallic catalyst) or catalyst blends [37].
8.6 Processing-Structure-Properties Relationship
in PE Fibers and Nonwovens
Polyolefin resins are widely used in various hygiene and medical applications such
as baby diapers, training pants, sanitary napkins, adult incontinence diapers, medical gowns, and many others. Polyolefin-based thermal bonded spunbond fabrics, melt
blown fabrics, and elastic laminates have been employed in these applications for
strength, comfort and fit, containment, and barrier properties.
Polypropylene (PP) comprises a significant portion of the materials used for spunbond and melt blown fabrics in hygiene and medical applications [38–43]. Ziegler-Natta
Processing-Structure-Property Relationships in Polyethylene 303
catalyzed (vis-broken) or metallocene catalyzed homopolymer PP (hPP) resins are
widely used to make thermal bonded spunbond fabric due to their excellent spinnability, broad thermal bonding window, high abrasion resistance, and high tensile strength.
However, hPP-based nonwovens exhibit low drapeability (stiff fabric). To meet the
market trend for increased softness, The Dow Chemical Company commercialized
ASPUNTM fiber grade PE resins for hygiene and nonwovens applications in 1986 to
provide fabrics with improved softness and drapeability compared to hPP-based spunbond. The applications of ASPUN resins include monocomponent and bi-component
spunbond fabric and binder fibers.
Bicomponent spunbond fabrics are extensively used in order to combine the
properties of the hPP and PE resins in the final application. Typically, PE is used in the
sheath layer to impart desired softness and hand feel and hPP is used in the core layer
for improved spinnability and to help maintain the physical properties such as tensile
strength and abrasion resistance.
There is a market trend to provide hygiene and medical products that exhibit
improved comfort and fit by means of increased softness, more cloth-like feel, and
increased drapeability. Despite the improved softness and haptics, PE resins are not
widely used in monocomponent spunbond nonwoven applications due to their inferior spinnability, narrower thermal bonding window, lower tensile strength, and lower
abrasion resistance of the final spunbond fabric. This was stated by Quantrille et al.
[44]: “The bonding of LLDPE filaments into a spunbond web with acceptable abrasion
resistance has proven to be very difficult, since acceptable fiber tie down is observed at a
temperature just below the point that the filaments begin to melt and stick to the calender. Because of this very narrow bonding window and the resulting abrasion resistance
and fuzz properties, spun bonded LLDPE nonwovens have not found wide commercial
acceptance.”
Patel et al. hypothesized that resins exhibiting slower crystallization rates and
broader melting distribution lead to a broader thermal bonding window and improved
bond strength (abrasion) in thermal bonded spunbond fabrics [45]. The absolute and
relative rates of crystallization are of basic importance in polymer fabrication processes
such as the spunbond process. Khanna [46] has proposed the crystallization rate coefficient (CRC) parameter to compare crystallization rates of various semicrystalline polymers. CRC is defined as the change in cooling rate in degrees Celsius per hour (°C/h)
required to lower the crystallization temperature by 1 °C. Patel et al. used Perkin-Elmer
Diamond DSC to measure crystallization onset temperatures of hPP and ASPUN resins
at high cooling rates [45]. Figure 8.16 compares the crystallization onset temperatures
of the hPP and ASPUNTM resins at high cooling rates using a semi-log plot. The CRC
for each resin was calculated using the slope of the linear fit from 10 °C/min to 50 °C/
min. Patel et al. measured CRC of about 308 °C/h for ASPUN PE fiber grade resin
versus about 120 °C/h for hPP resin suggesting the ASPUN resin crystallized about 2.3
times faster than the hPP resin [45]. This supported the hypothesis that narrow thermal
bonding window of PE resins is due to its very fast crystallization rates compared to
hPP resin. There are no practical ways to decrease the crystallization rates of neat PE
resins used in spunbond fabric other than to reduce stress during spinning to minimize stress-induced enhancement in rate of crystallization. Patel et al. showed that the
bonding window and abrasion resistance of monocomponent spunbond fabrics made
Handbook of Industrial Polyethylene and Technology
Crystallization onset temperature, °C
304
125
120
ASPUN
115
110
105
PP H502-25RG
100
95
5
10
20
50
100
Cooling rate, °C/min
200
500
Figure 8.16 Crystallization onset temperature of the ASPUNTM PE (30 dg/min MI, 0.955 g/cm3) and hPP
(25 dg/min MFR) resins as a function of cooling rate up to 400 °C/min [45].
from ASPUN fiber grade PE resins can be significantly improved by blending with a
polyolefin plastomer (POP) resin (30 dg/min MI, 0.913 g/cm3) [45]. This was mainly
due to a broader melting distribution of the resultant blend composition. Such blends
also exhibited improved spinnability compared to the ASPUN resin. Spunbond fabrics
made from the ASPUN/POP blend exhibited significantly higher softness rating compared to the hPP spunbond and carded fabrics.
Polyolefins are also extensively used to make sheath-core bi-component staple binder
fibers. PE is used as a sheath layer to give a lower melting sheath for flow and bonding to other heterogeneous fibers such as cellulose and glass [47, 48]. A high melting
point polymer like PP or a polyester is used in the core of a bi-component staple fiber.
A maleic anhydride grafted PE (MAH-g-PE) is typically blended in the PE sheath layer
to provide improved adhesion to polar fibers such as cellulose and glass fibers. Patel
et al. showed that bi-component binder fibers having PE sheaths made from a blend of
ASPUN PE resin and AMPLIFYTM MAH-g-HDPE when blended with cellulose fibers
gave excellent tensile strength in bonded air-laid pads [45].
Smith et al. investigated effects of bond roll pattern and bonding temperature on
the microstructure, tensile properties, flexural rigidity, and abrasion of PE nonwoven
(carded) fabric [49]. They showed that the nonwoven fabric generally failed in destructive testing by disintegration of the bond point; i.e., fibers pulling out the bond area
rather than breaking at the bond perimeter. An increase in bond roll temperature
increased the breaking (tensile) strength, initial modulus, and abrasion resistance, but
also increased flexural rigidity (undesired). A decrease in bond point concentration
significantly increased the elongation to failure and decreased flexural rigidity of PE
nonwoven fabrics.
8.7 Summary
This chapter covered various studies conclusively demonstrating that in addition to
the five primary molecular parameters, the molecular orientation and morphology
Processing-Structure-Property Relationships in Polyethylene 305
imparted during fabrication processes can have a major effect on the mechanical properties of the final fabricated PE parts. The post-die extensional flow in blown film, cast
film, extrusion coating, blow molding, melt spinning, and other processes determines
the induced final molecular orientation and subsequent properties. Overall, the measurement of film molecular orientation helps to understand mechanical properties,
especially tear and dart impact strength of blown films as a function of processing conditions. However, correlating orientation measurements to the film properties of different PE resins is always challenging as resin intrinsic properties would be different along
with differences in orientation (due to differences in melt relaxation times) in the final
blown films.
Acknowledgments
The author would like to thank his colleagues Herbert Bongartz, Shaun Parkinson, and
John Sugden for their critical review and suggestions.
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to the Effect of Resin Characteristics on Parison Formation in Extrusion Blow Molding,
Polym. Eng. Sci., 49, 251, 2009.
30. Eggen, S., and Sommerfeldt, A., Online Measurement of Parison Geometry During Blow
Molding: Parison Swelling for Three High-Density Polyethylenes with Different Molecular
Weights and Molecular Weight Distributions, Polym. Eng. Sci., 36, 336, 1996.
31. Debbaut, B., Homerin, O., and Jivraj, N., A Comparison between Experiments and
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32. Ariawan, A.B., Hatzikiriakos, S., Goyal, S.K., and Hay, H., Effects of Molecular Structure on
the Rheology and Processability of Blow-Molding High-Density Polyethylene Resins, Adv.
Polym. Technol., 20, 1, 2001.
33. Soares, J.B.P., Abbott, R.F., and Kim, J.D., Environmental Stress Cracking Resistance of
Polyethylene: The Use of CRYSTAF and SEC to Establish Structure-Property Relationships,
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34. Strebel, J.J., and Benson, M., The Effect of Processing Variables on the Environmental Stress
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36. Seguela, R., Critical Review of the Molecular Topology of Semicrystalline Polymers:
The Origin and Assessment of Intercrystalline Tie Molecules and Chain Entanglements,
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to Univation Technologies LLC, 2011.
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78(6), 185, 1995.
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Properties of Polypropylene Spunbond Fabrics, J. Appl. Polym. Sci., 98, 2355, 2005.
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and Properties of Polypropylene Copolymer and Homopolymer Filaments during a
Spunbonding Process, J. Text. Inst., Part 1, 89(2), 289, 1998.
42. Bhat, G.S., Nanjundappa, R., and Kotra, R., Development of Structure and Properties
During Spunbonding of Propylene Polymers, Thermochim. Acta, 392, 323, 2002.
43. Bhat, G.S., Jangala, P.K., and Spruiell, J.E., Thermal Bonding of Polypropylene Nonwovens:
Effect of Bonding Variables on the Structure and Properties of the Fabrics, J. Appl. Polym.
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9
Mechanical Properties of Polyethylene:
Deformation and Fracture Behavior
Alexander Chudnovsky1* and Kalyan Sehanobish2
1
2
University of Illinois at Chicago, Chicago, Illinois, USA
The Dow Chemical Company, Midland, Michigan, USA
Contents
9.1 Introduction ............................................................................................................................... 310
9.2 Stress-Strain Relations for PE .................................................................................................. 313
9.3 True Stress-Strain-Temperature Diagrams ............................................................................ 315
9.4 Time Dependency of Necking in PE ...................................................................................... 320
9.5 Accelerated Testing for PE Lifetime in Durable Applications ............................................. 322
9.6 Temperature Acceleration of SCG in PE ................................................................................ 326
9.7 Conclusions................................................................................................................................ 333
References ............................................................................................................................................ 333
Abstract
Polyethylene (PE) [– (CH2–CH2) n–] is suitable for applications ranging from adhesives to engineering structures such as pipes and tanks. In addition to homopolymers, a fairly large portion
of the polyolefin market employs the branched polyolefins linear low density PE (LLDPE) and
low density PE (LDPE). Polymers like PE are covalently bonded long-chain structures in melt
or solution which influence how the crystals are formed during solidification. Molecular weight
distribution (MWD) and branch distribution are two very important characteristics that control
the morphological organization during solidification that ultimately influence the thermomechanical properties. Application of PE is too broad to cover every type of thermomechanical
and field performance tests used, but this chapter will list some of the tests done in the industry.
Strong time dependency manifested in dimensional change over time and also a strong dependency of yield strength, draw stress, elongation at break, stiffness on temperature, and applied
strain rate of PE are discussed. A significant portion of this chapter is focused on PE necking and
its time dependency, since it is a main contributor to the ductile PE failure by large deformation
as well as to the brittle fracture by crack growth that limits functional service time of PE. Next
the focus is on accelerated testing methods for PE lifetime analysis followed by mechanism and
kinetics of brittle fracture. It is demonstrated that PE resistance to slow crack growth (SCG) is
directly related to a microscale cold drawing in the vicinity of a crack tip. Finally, the chapter
*Corresponding author: achudnov@uic.edu
Mark A. Spalding and Ananda M. Chatterjee (eds.) Handbook of Industrial Polyethylene and Technology, (309–336)
© 2018 Scrivener Publishing LLC
309
310
Handbook of Industrial Polyethylene and Technology
ends with a warning that due to a transition in mechanism of SCG, presently used standards may
lead to an overestimation of time to failure that is contrary to conservative design philosophy.
Keywords: Polyethylene, mechanical properties, necking, slow crack growth, lifetime
prediction
9.1 Introduction
Polyethylene has its origins in gaseous CH4 (methane). It can also start from a colorless gas called ethylene (CH2=CH2; C2H4) and thus has found its modern name polyethylene [1]. A gas like C2H4 (a by-product) can be polymerized first to convert into a
liquid oligomer (CH3 CH2-----~~-----CH2 CH2 CH3). On further polymerization it
becomes a soft waxy solid (roughly 40 carbon atoms). Even further, it will become a
hard solid homopolymer (in short polymer) called PE [–(CH2–CH2)n–], suitable for
applications ranging from adhesives to engineering structures such as pipes and tanks.
Upon degradation, solid (and molten) homopolymer can often undergo chain scission back to the oligomeric state, together with cross-linking. In the long-chain state,
polymers like PE are covalently bonded, random coil structures in a melt or a solution. These chain structures in the solution or melt state influence how the crystals are
formed during solidification. For example, in highly linear PE (no branches) homopolymers, only all-trans planar zigzag sequences in the chain crystallize.
Other than linear homopolymers of PE like high density PE (HDPE), a fairly large
portion of the polyolefin-based thermoplastics market is serviced by the branched
polyolefins (LLDPE and LDPE). Short branches are introduced in the PE backbone
in a random to semi-random fashion by comonomers, leading to LLDPE [2, 3]. If we
continue to increase the number of carbons in the (C2H4) monomer we will end up with
C5H10 (pentene), C6H12 (hexene), C7H14 (heptene), C8H16 (octene), and so on. During
polymerization two of the carbon units enter the main chain leaving the rest to form
short branches. Figure 9.1 shows the molecular structure of hexyl short branches in
PE resulting from C8H16 comonomer used during polymerization. Molecular weight
distribution (MWD) and branch distribution are two very important characteristics
that lead to morphological organization during solidification, ultimately influencing
the thermomechanical properties. Catalyst choices (Ziegler-Natta, metallocene, etc.)
and polymerization process variables/types (solution, gas and slurry phases) ultimately
decide the molecular arrangements in LLDPE. However, if polymerization occurs
under high pressure branches can be formed simply by a reaction mechanism called
backbiting [1]. These branches are called long-chain branches and the resulting polymers are known as LDPE. More unique branch structures can be created in this process
using free-radical polymerization.
The next level of organization of these long-chain structures occurs if some portions
of the polymer chains can organize into crystal units. All grades of PE (HDPE, LLDPE
and LDPE) are semicrystalline in microstructure. When a polymer of sufficiently regular structure is crystallized from a dilute solution, single crystals can be obtained in the
form of regular lamellae of uniform thickness [4–8]. In commercial polymers that crystallize from melt and form large crystalline aggregates, it is common to see only nonuniform lamellar structures. Lamellae thickness is of the order of 100–200 Å whereas
Deformation and Fracture Behavior of Polyethylene 311
H2
C
C
H2
H2
C
n
H
C
C
H2
H2
C
Abstract hydrogen
H
C8H16
Octene
C6H13
H2
C
C
H2
H2
C
n
C
H2
Hexyl branch
CH
C
H2
Figure 9.1 Example of branched PE with hexyl branches.
the chain length itself can be as high as 105 Å long. While the lamellae thickness is fixed
for a polymer with a certain melting point, other dimensions of the lamellae are controlled by many competing parameters related to the crystallization process and molecular weight. It was found that the lamellae have regularity in thickness dimension only
[3, 9, 10]. In melt crystallized PE, flat structures of lamellae are hardly possible and give
rise to what is known as curved lamellae fibers [9]. These curved lamellae often form
bundles that aggregate to form spherical morphology also known as spherulites [2, 11].
Spherulites are not pure crystals and comprise large amounts of amorphous material
in the intercrystalline spaces [9, 10, 12]. Molecular shape and the way molecules are
arranged in a polymeric solid lead to a morphological hierarchy which finally determines the properties of PE. Semicrystalline polymers like PE have the largest variation
of diverse features based on the scale of observation such as single crystal unit, spherulites, lamellae, and fiber bundles.
The amorphous or glass-like structures within the semicrystalline PE show no
long range order and they derive their properties through relatively less mobile chain
entanglements. Anchors like entanglements can slowly disentangle while melting or
the breakdown of crystals is necessary for destabilization of crystalline structures.
Tie chains are less well-characterized structures that constitute the amorphous layer
between crystalline lamellae. Tie chains have an important influence on the resistance
to relative deformation between adjacent crystalline lamellae at higher strains [13].
Comonomers like 1-octene or 1-hexene play a key role in formation of tie chains by
disrupting crystal lamellae. This disruption also leads to decreasing the crystallinity.
Another aspect of polymers including semicrystalline polymers is the state of orientation attained during melt or solvent flow needed to fabricate these polymers. The orientation state dictates the three-dimensional mechanical properties of PE. For example, in
the direction of orientation, crack propagation is easier. Extent of orientation under the
same state of flow is dictated by the shear thinning behavior of the polymer. In general,
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Handbook of Industrial Polyethylene and Technology
LDPE shows higher shear thinning behavior than the linear polymers like LLDPE. Due
to poor mechanical properties, LDPE is generally used as a blend component to impart
melt strength and processing ease needed in fabricating blown films, extrusion coated
films, cast films, and others.
The applications of PE are too broad to cover every type of thermomechanical testing needed. For simplicity, we will focus on flexible films used in monolayer, multilayer
industrial, consumer packaging and food, and specialty packaging applications. We will
also cover rigid engineering applications such as extruded pipes and blow molded tanks.
In flexible films, although they are highly deformable, engineering stress-strain curves
are commonly employed. Young’s modulus or 2% secant modulus, yield point, postyield stress drop at a specified rate, and elongation at break are the common parameters
extracted from this characterization for comparative purpose employing ASTM D882.
Modulus and yield stress are primarily affected by the total crystallinity. There are secondary effects of spatial dimensions of the crystallite and the state of orientation [14].
For tear resistance, the Elmendorf tear test based on ASTM D1922 is commonly
employed. This test is done in a pendulum type apparatus. The work done in tearing a
specimen is read from the pointer at the tip of the pendulum arm as a measure of tear
strength. Elmendorf tear is the most popular test used in the industry for the basic
characterization of flexible films in monolayer and multilayer configuration. Other
tests, like tensile tear, are also employed at certain thicknesses (especially thin sheet)
based on ASTM D1938. While the tear test is believed to give a measure of resistance
to preexisting cuts or slits, it is not adequate for a thin film or sheet characterization.
Experts in the industry believe that when a package drops from hand the resistance to
such impact could be characterized in individual films using the free falling dart test
according to ASTM D1709.
The dart impact machine utilizes an enclosed dart well containing a single unobstructed vent or sometimes no venting with a specified minimum area. A two-piece
annular specimen clamp mechanism is used to hold the film sample firmly in place
during the test. The apparatus drops the dart from specified heights for Test Method
A and for Test Method B. Darts for Test Methods A and B have hemispherical heads
of specified dimensions and long shafts of specified length to accommodate removable
incremental weights. The dart head surfaces must be free of nicks, scratches, or other
irregularities. Although the dart head size is specified, material of construction can
be chosen depending on the case. Data have shown to be sensitive to the material of
construction and finish of the dart head used. This usually is an important part of the
characterization since samples have frictional interactions with the dart head.
It is believed that sharp objects from outside or inside of packaging could be responsible for the puncture of films. Moreover, films can also be punctured and snagged,
leading to a tear. Industry has designed many types of analytical tests to address these
concerns with many of the tests depending on the application. Some of the more common ASTM tests include: 1) puncture propagation tear resistance (ASTM D2582), 2)
high speed puncture (D3763/D7192), 3) hydrostatic puncture testing (D5514), and 4)
static puncture (D6241). D2582 is actually a puncture-propagation of a tear test that
measures the resistance to dynamic puncture and propagation of that puncture resulting in a tear. D3763/D7192 provide load versus deformation response of plastics under
essentially multi-axial deformation conditions at varying impact velocities. This test
Deformation and Fracture Behavior of Polyethylene 313
method further provides a measure of the rate sensitivity of the material resistance
to impact. This test, similar to the dart test, is affected by the frictional response of
the probe to the film and needs to be taken into account. D5514 is a puncture testing
method applied to materials for geomembranes. Hydrostatic puncture testing is essentially done inside a pressure vessel. D6241 is also used to measure the force required to
puncture a geotextile and geotextile-related products.
Some or all of the film tests described above are only used in qualifying for a specific application, and each value chain in the industry has its own qualifying methods.
However, once the package is made using the PE film, it has to be further tested in the
final form. In many cases PE film is simply a component in a multilayered structure.
There are too many brand owner specific test methods on the final filled package, and
this chapter is not going to address them all. One of them includes dropping filled
packages from different heights and also dropping the package on various sides (such
as flat side, edge with the seal, etc.) to understand the performance. Some brand owners actually do some of the tear tests described above on the final filled package by
puncturing it and dragging by a probe. Static load tests are also done on filled packages stacked on each other for long-term creep at warehouse temperatures. Shipping
the packages involves vibration with stacked filled packages, which are studied under
simulated vibrations.
In rigid applications like fuel tanks, bottles, and drums, tensile tests based on ASTM
D638 are utilized. Several other tests, however, are specified depending on the application. Flexural modulus based on ASTM D790, notched Izod impact, forced and falling
dart impact, and creep tests are the most common. For applications like pipe, resistance
to continuous increase in pressure is characterized by a burst test and hoop strength,
slow crack propagation is characterized by PENT test based on ASTM F1473, and rapid
crack propagation (RCP) in pipe observed in field is simulated by a test developed at the
Imperial College of Science and Technology in the UK called S4 [15]. In pipes there is a
very involved lifetime prediction methodology called hydrostatic pressure testing, generally following ISO 1167. Environmental stress crack resistance (ESCR) tests are very
common in any application intended for a long service time and could be subjected to
environmental degradation. There are many types of such tests that either tend to follow some standard or each specific company specifies their own methods.
9.2 Stress-Strain Relations for PE
Basic characterization of mechanical properties of PE starts with stress versus strain
relations. Researchers skilled in the art can extract most of the fundamental parameters necessary from such relations to predict performance in any real-world application situation. The following paragraphs will discuss the fundamental parameters of PE
instead of focusing on the test specific parameters described earlier. Most commonly a
simple uniaxial stretching of a standard tensile bar (ASTM D638) is conducted for rigid
PE applications. The terms that are important here include: engineering tensile stress,
(en)
, in the direction of tension acting over an initial cross-sectional area, A0 , perpendicular to the tension is defined as the applied force, F, divided by A0. It is also called
“normal,” i.e., orthogonal to the cross-section stress. The engineering strain, (en), in the
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Handbook of Industrial Polyethylene and Technology
tension direction is the change of the test bar length l = l – l0 divided by the initial bar
length, l0, (l is the current length of the deformed bar). Thus, the engineering stress,
(en)
, and engineering strain, (en), in uniaxial tension is expressed as follows:
F
;
A0
( en )
l l0
l0
( en )
l
(9.1)
l0
According to the above definitions the stress has dimension of force per unit area; for
example, Newton per square meter or a Pascal (Pa) in the International System (SI) of
units, or lbf/in2 (psi) in the Imperial units system. The strain is a dimensionless quantity.
In many engineering applications of traditional materials where only a small strain
is considered (on the order of a fraction of a percent), the above definitions work well.
However, PE experiences large deformations (up to a few hundred percent) in most
applications ranging from very thin film (<20 m) to a few cm thick parts. Simply
replacing ASTM D638 by ASTM D882 does not help the situation when measuring
thin specimens with extensively large deformation. Instead, the engineering stress
needs to be changed to true stress for all PE applications. The current cross-sectional
area A (that varies as the bar is stretched) rather than the initial area should be used in
defining the true stress.
(tr )
F
A
(9.2)
Assuming that the material is incompressible, i.e., the volume of the bar does not
change in the process of deformation ( A l A0 l0 ), the true stress is related to the
engineering stress as follows:
(tr )
F
A
F l
A0 l0
( en )
( en )
(1
( en )
),
(9.3)
where is a stretch parameter defined as
l /lo (1 (en) ) .
The measure of deformation such as the engineering strain introduced in Equation 9.1
becomes unacceptable for PE when the bar length increment l is large in comparison to
the initial length l0. The problem with engineering strain as a measure of deformation
is that it does not possess additivity. A proper measure for large deformation is introduced as the true strain (tr), defined as the sum of the infinitesimal engineering strains;
i.e., infinitesimal length increments dl normalized by the current length l (In a limit of
the length increment dl 0 the sum converts to a definite integral):
l
(tr )
dl
l
l0
ln
l
l0
ln
( en )
ln(1
In Equation 9.4 a current draw ratio or a stretch
For small engineering strains (en) 1, ln(1 (en) )
)
(9.4)
l /lo (1 (en) ) is introduced.
due to the properties of a
( en )
Deformation and Fracture Behavior of Polyethylene 315
natural logarithm. Thus, for a small (en), the engineering strain is a good approximation
of the true strain: (en) (tr). This justifies the common use of engineering strain in the
mechanics of traditional materials, such as metals, rocks and concrete, which experience small deformation in load-bearing applications. At the same time, a plastics engineer should be aware of the limitations of the engineering strain when dealing with PE.
We can also decompose the force acting on a selected cross-section into normal and
tangential components. The tangential components of force per unit cross-sectional
area define the shear stresses. There are three normal stresses and three shear stresses
acting at a point of a solid. Together they constitute a more complex object: the stress
tensor. The normal and shear stresses are the components of the stress tensor that
characterize the stress state at a given point [16, 17]. The values of stress components
depend on the orientation of the 3D coordinate system. For example, there is a special
orientation of coordinates, called principle directions, for which all three shear stress
components vanish. Then there are invariants of the stress state that do not depend
on the coordinates used. For instance, the three normal stresses in the principle directions, called maximum, minimum, and intermediate principle stresses, are such invariants [16, 17]. In the uniaxial tension test the engineering stress (en) and true stress
(tr)
used in presenting the test results are the maximal engineering and true principle
stresses, respectively. Another important invariant is the hydrostatic stress that strongly
affects the response of PE to loading: a high hydrostatic compression enhances ductile
behavior similar to that by high temperature or low strain rate, and hydrostatic tension
promotes a brittle response to the loading resembling the effect of low temperature or
high strain rate. For example, the presence of hydrostatic tension in front of a crack
is responsible for cavitation and crazing. Other invariants of stress tensor can also be
involved in formulation of yielding and fracture criteria.
A strain tensor that represents deformation also has three normal and three shear
components similar to the stress tensor [16, 17]. The relations between stress and strain
tensors present constitutive equations of the material. For a small deformation, a linear elastic stress versus strain relation (Hooke’s law) is commonly used. Unfortunately,
Hooke’s law has a very limited applicability for PE that exhibit highly nonlinear timedependent (viscoelastic) behavior. There are two origins of PE nonlinear behavior:
1) a physical nonlinearity, such as nonlinear stress versus strain relation, and 2) a geometrical nonlinearity associated with the large deformation. Temperature, strain rate,
and stress state also affect the response of PE to loading, as discussed above.
9.3 True Stress-Strain-Temperature Diagrams
Schematics of a) engineering stress versus engineering strain and b) true stress versus true strain, both presenting a PE uniaxial tension ramp test data, are shown in
Figure 9.2 [16–18]. At the very beginning of the curves (2% ~ 5% strain) the process of
deformation is almost linear elastic (elastic deformation implies that the process is fully
reversible and unloading follows the same path as the active loading). As the deformation increases the nonlinearity of the curve it becomes more and more pronounced,
eventually approaching the yield point, as shown in Figure 9.2a. If the loading stops
before reaching the yield point and the specimen is unloaded, the unloading path may
Handbook of Industrial Polyethylene and Technology
316
300
250
20
Necking
15
(MPa)
25
Strain
hardening
Engineering draw stress ( dr(en))
10
Draw ratio:
= n(en) + 1
n
5
0
(a)
True stress (tr)
Engineering yield stress
Engineering yield stress ( y(en))
Stress
Engineering strain
(en)
(MPa)
30
0
2
200
150
True yield stress ( y(tr))
4
6
Engineering strain (en)
(en)
b
8
0
0.0
10
(b)
Strain
hardening
(tr) )
s ( dr
tres
s
raw
ed
Tru
ng
ck i
Ne
100
50
(en)
n
(en)
0.5
1.0
1.5
True strain (en)
(tr)
n
2.0
(tr)
b
2.5
Figure 9.2 Schematics of stress and strain; (a) engineering stress and engineering strain curve, and
(b) true stress and engineering stress stretch curves.
return to the origin, but departs from the initial loading path, thus forming a so-called
hysteresis loop. The area of the hysteresis loop in stress-strain coordinates presents the
energy dissipation and suggests that the loading-unloading process may be mechanically reversible, but thermodynamically it is an irreversible process; i.e., during the hysteresis some irreversible changes of PE molecular architecture and/or morphology take
place.
With further increase in the engineering strain, just after yield stress y there is a sharp
decrease in stress from y to a drawing stress dr. After that the stress remains almost
constant with continuing increase of strain that depends on the specimen length. At
the end of the plateau the stress starts to increase with increasing strain until ultimate
fracture takes place. The true stress versus true strain diagram is reconstructed from
the above engineering stress-strain curve using the true stress and true strain expressions; i.e., Equations 9.3 and 9.4. True stress and true strain are shown in Figure 9.2b.
At small strains the true stress and engineering stress versus strain curves are very close.
However, with increasing strain the true stress departs significantly from the engineering stresses, as shown in Figure 9.2b.
The stress versus strain curves strongly depend on temperature and strain rate. A set
of engineering stress-strain curves for a pipe grade HDPE tested at various temperatures is shown in Figure 9.3a. At the room temperature 23 °C there is a relatively sharp
decrease from y to a level slightly below the plateau and then increasing to the plateau level dr. The difference ( y
) decreases with increasing temperature. The yield
dr
and draw stress dependency on temperature observed in these tests are depicted in
Figure 9.3b. The difference between the yield and draw stresses disappears as test temperature approaches the melting temperature Tm of PE.
Another set of engineering stress-engineering strain curves for the same HDPE
tested at 23 °C and various strain rates is shown in Figure 9.4a. This set of curves resembles the one shown in Figure 9.3. At the highest strain rate ( 0.139 s 1 ) there is
a relatively sharp decrease from y to dr, and then the plateau continues until strain
hardening takes place. With decreasing strain rate the stress decrease becomes gentler,
similar to that observed with an increasing temperature in Figure 9.3a. The difference
( y
) also becomes smaller with a reduction of the strain rate similar to that with
dr
increasing temperature. The yield and draw stresses dependency on strain rate is shown
Deformation and Fracture Behavior of Polyethylene 317
30
23 C
40 C
60 C
80 C
25
15
20
10
5
5
0.5
1.0
1.5
2.0
2.5
Engineering strain (mm/mm)
0
3.0
dr
15
10
0
0.0
= –0.29x + 32.3
R2 = 0.98
y
20
(a)
y
25
Stress (MPa)
Engineering stress (MPa)
30
dr
= –0.16x + 20.2
R2 = 0.99
0
20
40
60
Temperature (°C)
(b)
80
100
Figure 9.3 Stress versus strain curves strongly depend on temperature and strain rate for a commercial
HDPE: (a) engineering stress-strain curves at a strain rate of
0.028 1/s for different temperatures,
and (b) yield strength and draw stress temperature dependency.
35
30
20
15
10
(a)
25
y
20
= 3.71 log + 19.77
R2 = 0.95
dr
15
dr
c
10
5
0
0.0
y
30
Engineering stress MPa
Engineering stress (MPa)
25
0.139/sec
0.0278/sec
0.0167/sec
0.00333/sec
0.000667/sec
= 1.34 log + 13.95
R2 = 0.80
5
0.5
1.0
1.5
2.0
2.5
Engineering strain (mm/mm)
0
3.0
(b)
–5
–4
–3
1
2
–1
0
Log (strain rate, hour–1)
c –2
3
4
5
Figure 9.4 Engineering stress curves for HDPE at 23 °C [19]: (a) engineering stress-strain curves for
various strain rates including
0.028 s 1 used in Figure 9.3a, and (b) yield strength and draw stressstrain rate dependency.
in Figure 9.4b. The similarity of the decreasing strain rate and increasing temperature
effects on yield and draw stresses is one of the manifestations of the well-known timetemperature correspondence.
It will soon be clear that both engineering stress vs. engineering strain curve
(Figure 9.2a) as well as true stress versus true strain relation (Figure 9.2b) are deficient
since both curves reflect a mixture of material properties and the test specimen geometry. Careful observations of the specimen deformation during the test and analysis of
the data are required to extract the true material properties from the recorded data. It
should be emphasized that the properties that can be attributed to a small (but macroscopic) domain of a material can be reconstructed from a test of a large specimen
made of the same material only if the specimen is a homogeneous one (all points of the
specimen are the same). Keeping that in mind, we reexamine the standard uniaxial tension test data of the same commercial pipe grade HDPE used for generating Figures 9.3
and 9.4. Figure 9.5 explains the origin of the problem arising from interpretation of
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Handbook of Industrial Polyethylene and Technology
Drawn
material
Stress
Original
material
n
= n+1
y
A
C
dr
B
O
A
O
n
Strain
B
C
Figure 9.5 Schematics of stress-strain curve for uniaxial tension test and specimen appearance at various
stages of the test (adopted from [20]).
PE stress-strain curve as material property [20]. It presents a schematic illustration of
stress-strain curve of PE together with the test specimen configurations at the various
stages of the test. On the left-hand side, there is a sketch of the original homogeneous
specimen at the beginning of the test. The horizontal marks have been printed on the
specimen to monitor the relative motion of material points.
As the deformation increases and approaches the yield point, the stress-strain curve
departs from a straight line. This corresponds to the orientation of lamellas and amorphous chains, crystal rotation, and sliding as a precursor of yielding. At the yield stress,
the lamellae structure is broken (fragmented), allowing rearrangements. Randomly
coiled molecular chains between the lamellae as well as some of the macromolecules
initially incorporated into lamellae are first released and then become highly stretched
and oriented in the direction of the applied tension. When the stretch (orientation)
of macromolecules reaches a critical level a recrystallization takes place. Some of the
oriented molecules are incorporated into newly formed crystals and “freeze” the newly
formed oriented state. The stretching in axial direction is accompanied by thinning
(reduction of the cross-sectional area perpendicular to tension) since the volumetric change in PE is very small. Such large strain localization called cold drawing (or
necking) takes place at a random location within the specimen. The onset of necking is depicted in the second from the left sketch of the specimen in Figure 9.5. The
localized necking accompanied by a large displacement at the grips occurs much faster
than the testing machine cross-head speed. Since it is a displacement control test, an
unloading is recorded with the stress reduction from y to dr. The fully developed neck
within the specimen domain between two neighboring marks in the original specimen configuration is depicted in the left sketch of the specimen on the right side from
the stress-strain curve of Figure 9.5. Notice the difference in the distance between the
same two marks in the initial and drawn state. There are narrow boundaries between
the drawn (oriented) material between the two marks and original (isotropic) material
outside of that domain. The ratio of the distances between the same marks before and
after necking is called the “natural draw ratio” n. The oriented (drawn) state formed
in the necking process is preserved upon unloading of the specimen, except for a very
small elastic strain recovery of the drawn material. After initiation of necking, a plateau
with a practically constant level of stress dr is observed in the engineering stress-strain
Deformation and Fracture Behavior of Polyethylene 319
curve. This phenomenon resembles metal plasticity but the mechanisms of PE cold
drawing is very different from that in the plastic deformation of metals. In PE there
are two narrow boundary layers separating the drawn material from the rest of the
specimen. During the plateau stage, these boundaries of the neck are spreading in both
directions. The domain of drawn material grows by transforming the initial PE into
drawn (oriented) state. In that process no changes in stresses and strains within the
initial and drawn domains takes place. Mechanical and phase equilibrium across the
boundaries are maintained. The recorded increase of the gage length results exclusively
from the increase in n times the drawn domain length in comparison to the length of
the corresponding original PE. Thus, the movement of the neck boundaries is ( n 1)
times smaller than the displacement at the grips. When the specimen is fully drawn,
the so-called “strain hardening” effect is observed. It represents the properties of oriented material (see the right-hand sketch of the specimen in Figure 9.5). Thus, the
cold drawing can be viewed as consecutive PE transformations: 1) fragmentation and
partial melting of lamellas of the initial semicrystalline PE, 2) stretching of the molecular chains, rotating, and orienting the remaining lamella fragments, and 3) crystallization of the oriented molecules and formation of a new highly oriented semicrystalline
state [21, 22]. The first transformation takes place when the yield strength of material
is reached, and the second one occurs when the entire gaged part of the specimen
y
is transformed into oriented state and both the drawing stress dr and critical level of
orientation (stretch = n) are achieved at the boundaries of the necked region. The
drawn (oriented) PE possesses noticeably different mechanical properties (e.g., rigidity, strength or creep behavior) than the original (isotropic) one. Upon unloading the
drawn material returns to a stress free, but oriented state. This results in a large irreversible deformation when compared with the initial state.
Thus, during the entire plateau stage of the necking process, the specimen is composed of two regions with different properties and variable sizes. Therefore, the plateau part of the engineering stress-strain curve (Figure 9.2a) and linearly increasing
portion of true stress-true strain curve marked as “neck propagation” in Figure 9.2b
represent a mixture of material properties with a composite specimen behavior dependent on a variable composition of the original and oriented materials. Therefore, the
PE properties independent of specimen geometry are reflected by the two parts of the
stress-strain curve: one is before yielding and another one is after the cold drawing
process is completed; i.e., the strain hardening region. The true PE stress-strain relations before and after cold drawing for the same commercial HDPE used in generating
Figure 9.3 are presented in a true stress-strain-temperature diagram in 2-dimensional
and 3-dimensional forms in Figure 9.6 [20–22]. The elastic behavior of the original and
oriented (drawn) PE at different temperatures is presented by region I and region III.
The isotherms of the region II connecting regions I and III correspond to a metastable
transition from the initial to the drawn state. The length of the isotherms of region II
reflects the natural draw ratio dependency on temperature. The transition takes place
within the thin boundaries of partially drawn PE and occurs much faster than the
boundaries propagation controlled by the cross-head speed of the tensile apparatus.
Thus, the neck boundaries in the ramp test propagate in a quasi-equilibrium manner,
maintaining mechanical (force) and phase equilibrium between the initial and oriented
states of PE [21, 22]. The PE cold drawing phenomenon discussed in this section has
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Handbook of Industrial Polyethylene and Technology
296 K (T/Tm 0.72)
313 K (T/Tm 0.77)
333 K (T/Tm 0.81)
10
353 K (T/Tm 0.86)
1
0.0
(a)
True stress (tr) (MPa)
True stress (tr) (MPa)
335
100
0.5
1.0
1.5
True strain (tr)
2.0
(b)
III
100
II
10
1
I
N0.72
or
m 0.77
.t
em 0.81
p.
T/ 0.86 0
T
m
1.5
1
(tr)
0.5
train
s
e
Tru
2
Figure 9.6 True stress-strain-temperature (TSST) diagrams based on stress-strain behavior shown in
Figure 9.3a: (a) 2-dimensional schematic, and (b) 3-dimensional schematic.
a direct relation to various mechanisms of PE fracture. Thus, the ductile failure of PE
is actually the material instability resulting in large macroscale deformation (necking)
under load. PE brittle fracture by cracking is, to a large extent, also controlled by cold
drawing on a microscale.
9.4 Time Dependency of Necking in PE
The mechanical behavior of original (isotropic), drawn (oriented) PE as well as the cold
drawing phenomenon itself exhibit a strong time dependency [23] that is not reflected
in the discussion above. As we have seen in Figure 9.4b, both yield and draw stresses
in the ramp test increase with increasing strain rate. The draw stress dr corresponds
to quasi-equilibrium neck propagation along the bar in the direction of tension and
reflects the mechanical equilibrium between the initial (isotropic) and oriented states
of PE. In contrast, the yield stress y corresponds to the nucleation of the oriented state
(onset of necking). Nucleation of a new state commonly occurs with an overshooting
of the conditions of the equilibrium coexistence of two states. For the PE ramp test, the
higher the strain rate the larger is the overshooting ( y
) at the onset of necking, as
dr
shown in Figure 9.4b. Extrapolating such trends to low strain rates can identify a strain
rate at which the yield stress coincides with the draw stress. Thus, the point of intersection between the two lines ( y vs. log ) and ( dr vs. log ) defines a “characteristic
stress c” and “characteristic strain rate c ” in Figure 9.4b [19]. The combined effects
of temperature and strain rate on the yield and draw stresses are depicted in Figure 9.7.
Figure 9.7a shows the yield and draw stresses dependency on the strain rate at various
temperatures; Figure 9.7b displays the characteristic stress c and characteristic strain
rate c dependency on temperature. Clearly, c decreases and c increases with temperature. These data are complimented by creep tests. Typical PE uniaxial creep strain
dependence on time at various stress levels is shown in Figure 9.8a [24]. The creep strain
monotonically increases with time, approaching a vertical asymptote (strain rate
).
A reduction of the cross-sectional area of the specimen and therefore an increase of true
stress accompanies the increasing axial creep strain. A uniform distribution of creep
Deformation and Fracture Behavior of Polyethylene 321
4.0×10–4
14
(MPa)
y
dr
c
c
12
3.0×10–4
c
25
.
Characteristic stress
(MPa)
20
23 °C
=11.44 MPa
15
c
10
5
c
0
–7
40 °C
= 8.32 MPa 60 °C
= 7.04 MPa
c
–6
(a)
80 °C
= 6.07 MPa
c
–2
–5
–4
–3
Log (strain rate, sec–1)
10
2.0×10–4
8
1.0×10–4
6
0.0
–1.0×10–4
100
4
–1
0
0
20
(b)
40
60
T (°C)
Characteristic strain rate (sec–1)
30
80
Figure 9.7 Yield stress and draw stresses: (a) yield stress and draw stress dependency on strain rate at
various temperatures (based on the data presented in Figures 9.3 and 9.4), and (b) characteristic stress
and characteristic strain rate dependency on temperature [20].
35
14.5 MPa
14.0 MPa
13.5 MPa
13.0 MPa
Engineering stress (MPa)
Strain(%)
150
12.0 MPa
100
n
50
tn
0
(a)
0
1
10
Creep time (hour)
25
20
(b)
dr
c
15
10
Creep
5
0
–5
100
y
30
–4
–3
–2
Ramp
–1
0
1
2
3
4
5
Log (Engineering strain rate, hour–1)
Figure 9.8 Engineering strains for the commercial HDPE as used in Figures 9.3 and 9.4: (a) engineering
strain (axial) vs. time in creep at 23 °C and various stress levels (the same), and (b) determination of
characteristic stress and strain rate using ramp (Figure 9.4b) and creep tests data [24].
strain over the entire specimen length is observed until a critical level of yield true
stress y* is achieved. At that time a spontaneous strain localization (onset of necking)
takes place. The critical level of strain and the duration of creep time prior to the onset
of necking are marked as n and tn accordingly. After initiation, the neck region spreads
relatively fast through the specimen until it is fully transformed into oriented material.
In Figure 9.8a this process at various stress levels is depicted by the vertical lines, since
the length of PE specimen increases 5 to 6 times over a relatively short duration of time
in comparison with creep time tn prior to necking. Necking that takes place after a certain time under constant applied stress that is significantly lower than yield strength is
referred to as “delayed necking” [22, 24]. Considering the average creep rate during the
steady creep stage and constancy of the applied stress we can add the creep test data to
data of the ramp test to more accurately determine the characteristic stress c and strain
rate c . The combined creep and ramp tests data used for determination of c and c
are presented in Figure 9.8b (compare with Figure 9.4b). The yield strength y vs log
and draw stress dr vs log can be mathematically expressed as:
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Handbook of Industrial Polyethylene and Technology
y
c
k y log / c
(9.5)
dr
c
kdr log / c
(9.6)
These equations are a basis for prediction of long-term PE ductile failure using shortterm testing [19, 25].
9.5 Accelerated Testing for PE Lifetime
in Durable Applications
In durable applications of PE, there is a well-recognized quest for accelerated testing, since the life expectancy of engineering structures made of PE varies from about
10 years for automotive applications to 50 and even 100 years for building and construction, water distribution, and natural gas transmission systems. Commonly accepted
failure accelerating factors are high stress levels, loading frequency (fatigue testing),
temperature, and aggressive environments. There are two main challenges in development of an accelerated testing for lifetime: 1) to reproduce in accelerated test the failure mechanism(s) that may take place in service, and 2) to establish a correspondence
between the time to failure in the accelerated testing and the failure time under service
conditions (extrapolation technique) [20].
Let’s consider failure acceleration by high stress that, as a rule, leads to an earlier
failure in comparison with intermediate or low stresses. However, the fracture under
high stresses is commonly a ductile one associated with large deformation (stretching)
of PE. An example of ductile failure of PE pressure pipes under a sustained pressure
test is shown in Figure 9.9. In creep under sustained pressure the pipe wall thins over
time and the true hoop stress increases until the onset of delayed necking at the maximum hoop stress location that coincides with a random minimum pipe wall thickness
position. Occasionally it may happen at two or more locations, as shown in Figure 9.9
Figure 9.9 Cross-section of a HDPE pressure pipe that failed by “ballooning” in a sustained pressure
test. Two regions of the thinned wall are the delayed necking regions blown into two unequal bulges by
internal pressure (commercial HDPE).
Deformation and Fracture Behavior of Polyethylene 323
(a multiple onset of necking is also observed in the uniaxial tension). This ductile failure of a pressure pipe is called “ballooning.” Ballooning can occur in any application
where a cylindrical part is under internal pressure. A biaxial stress state is the major
difference between ballooning in the pressure pipe and delayed necking in the uniaxial
tension creep. A thinned (drawn) pipe wall within the neck region in comparison with
the pipe wall outside of that region is well visible on the cross-section through the center of the bulge, as shown by Figure 9.9. The time tn of the delayed necking gives a good
approximation of the pipe failure time tf, since the duration of neck propagation (the
balloon growth) is small in comparison to tn. Thus, by running a number of sustained
pressure tests the ductile failure time logtf vs. log relation can be constructed, where
represents hoop stress. A large body of sustained hydrostatic pressure test data
summarized in the international standards ISO 9080 suggests the following empirical
relations between the ductile failure time, hoop stress and test temperature T:
logt f
A1
B1
T
C1
D1
T
log
(9.7)
where A1, B1, C1, & D1 are adjustable parameters determined from experimental data.
A similar expression is used in the North American standard ASTM D2837, but the
right-hand side of Equation 9.7 is truncated and only three adjustable parameters are
employed. The linear relation between logtf and log implies a power law form of failure time dependency on stress. The form of Equation 9.7 is not convenient for designers, since all four parameters have different dimensions and change their values with
changing basic units, e.g., {cm, N, K} vs. {in, Lb, °F}. Thus, it is more convenient to
use the dimensionless form of the above equation. Introducing a dimensionless stress
=
/ y we can rewrite Equation 9.7 as:
logt f /t 0
m log
(9.7a)
where
log t 0
A1
B1
T
m log
y
and the power m (c2 D1 /T ). The power m is determined by the slope of the straight
line approximation of the experimental data in log
and log tf coordinates.
Alternative modeling of ductile failure time follows from i) the recognition of ductile
failure as delayed necking, ii) a fundamental symmetry between logarithm of strain rate
in steady state creep vs. stress on one side and logarithm of ductile failure time vs. stress
on the other that have been proven more than half a century ago by Hoff [26], and iii)
the Equation 9.5 for yield strength and draw stress strain rate dependency. The time to
ductile failure tf tn stress dependency is determine from material instability condition
and on the basis of the above-mentioned symmetry (ii) that is expressed by simple relations between the coefficients of Equations 9.5 and 9.6 as follows [20, 23, 24]:
Handbook of Industrial Polyethylene and Technology
324
log t f /t *
kf
/ c, kf
k y 1 , if
c
and k f
kdr1 , for
c
(9.8)
= c.
where t* is the ductile failure time in creep under
Equation 9.8 gives a very close long-term ductile failure time prediction to that of
ISO 9080 and the ASTM D2837 empirical power law, but better captures the nonlinearity of tf stress dependency for short-term failure and, what is most important, requires
a much shorter testing program. In addition, the ductile failure time of geometry of
application can be scaled up from a simple uniaxial creep test by determining t* for the
geometry of application. This is a more accurate predictive tool than the approach used
in the standards.
It is tempting to extrapolate the linear relations log t f /t 0
m log of Equation 9.7a
or log t n /t *
kf
/ c of Equation 9.8 from high to intermediate and to low (in
comparison to the yield strength) stress levels. Unfortunately, such an approach does
not work, since there is a change in the mechanism of PE failure from ductile to brittle
(cracking) with a reduction of stresses, as illustrated in Figure 9.10 [20]. The brittle fracture of PE is mainly attributed to crack growth driven by either mechanical stresses at
intermediate stress levels or by a combination of low stresses and a chemically aggressive environment. Figure 9.10a illustrates the changes in the mechanisms of PE failure
with a variation of stress level.
The solid lines in Figure 9.10a schematically present the hoop stress dependency of
failure time tf in log hoop stress ~ log tf coordinates. The sketches and photos next to the
solid lines depict the typical appearances of the corresponding failure mode (I, II and
III). Quantitative diagrams of that type have been widely used for plastic components
in durable applications. A schematic illustration of such a diagram for PE pressure
pipes at various temperatures is shown in Figure 9.10b. Figure 9.10b displays combined log hoop stress ~ log tf relations for ductile and brittle modes of failure (modes
I and II). It does not show the degradation driven brittle facture (mode III) depicted in
Figure 9.10a. The stress sensitivity of the time to failure is manifested in the slope of the
lines: ductile failure is highly sensitive to stress variation (a very small change in stress
level results in a large variation of time to failure); degradation driven brittle failure
(mode III) in Figure 9.10a weakly depends on stress level, and the stress driven brittle
Ductile failure
(a)
Log stress
Log (Hoop stress)
20 °C
40 °C
I
II
60 °C
80 °C
Brittle failure
III
Log (time)
(b)
Log time
Figure 9.10 Lifetime stress dependency: (a) schematic illustration of lifetime-stress dependency, and
(b) schematics stress-time to failure-temperature relations for PE (adopted from [20]).
Deformation and Fracture Behavior of Polyethylene 325
fracture (mode II) shows an intermediate time to failure stress dependency, as can be
seen from Figure 9.10a. The change in failure mechanism with variation of the stress
is a direct indication of a strong limitation of failure acceleration by high stress level.
Temperature is widely used as an alternative accelerating factor for various physical
and chemical processes. Indeed, the pairs of straight lines corresponding to ductile and
brittle failure modes at various temperatures shown in Figure 9.10b closely resemble
one another. It suggests an existence of temperature-dependent shifts (double shift)
along log
and log tf directions that bring together the corresponding ductile and
brittle failure lines, as well as the so-called “knee” points representing the transition
from ductile to brittle mode of failure, and forms one master plot. Such empirical methods were proposed independently by Popelar et al. [27] and Lu and Brown [28, 29]
25 years ago. Both empirical propositions are practically identical in application to
the temperature range between 80 °C (353 K) and 20 °C (293 K). The double shift, if
proven, allows extrapolation of high temperature and short time data into low temperature service time prediction. However, the temperature-dependent double shift implies
a self-similarity of the stress-time dependency at different temperatures. Specifically,
the “knee” at various temperatures must belong to a straight line in log
~ log tf coordinates and the ductile and brittle failure lines for one temperature should be parallel
to the corresponding lines at other temperatures. Analysis of a large number of longterm tests with durations of 1,000 to 10,000 hours and longer conducted in the process
of ISO 9080 development, however, proved that log
vs. log tf relations at different
temperatures are not self-similar and a double shift, in general, is not applicable. Thus,
many years of PE pressure pipes testing and analysis of large amounts of data resulted
in two empirical equations for stress and temperature dependency of time to failure
(see Eq. 9.7) for ductile and brittle modes of failure recommended by ISO 9080 and
ASTM D2837 standards. In ISO 9080 the time to failure in brittle mode is expressed
in the same form as Equation 9.7, but with a different set of coefficients A2, B2, C2 &
D2. Thus, there are two equations for time to failure for contemporary standards: one
is for ductile and another one is for a brittle mode of failure with a total of eight (two
in ISO 9080) and (six for ASTM D2837) adjustable parameters. There is also a statistical method suggested for data analysis. However, there is no solid procedure for
identification of the “knee” point. It leads to a high uncertainty in determination of the
parameters for the equations. The shortcoming of empirical equations for PE lifetime
is obvious: too many adjustable parameters and a high uncertainty of the prediction.
Moreover, when a few years of test data are extrapolated to 50 or 100 years there is a
silent assumption of no material aging during the extrapolation period (50 or even
100 years). There is no experimental support for such an assumption. To the contrary,
there are anecdotal evidences of PE aging (changing properties) over long times in
service under low stress levels. It should be mentioned that ductile failure of PE parts
rarely happens in service except for some accidental overloading, since the parts are
designed for the service stresses that are significantly lower (about factor of 2 or lower)
than the PE yield strength. Therefore, ductile failure is primarily observed in laboratory
testing. The majority of field failures are associated with slow crack growth (SCG). In
the next section we focus on a brittle mode of failure and analysis of SCG acceleration
by elevated temperatures. Unfortunately there is a limitation of SCG acceleration by
temperature similar to the limitation of failure acceleration by stresses discussed above.
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Handbook of Industrial Polyethylene and Technology
Indeed, there are significant changes in the SCG mechanisms and kinetics resulting
from a variation of load and temperature [20, 30]. It is important to keep in mind that
the lifetime of a PE part in service (e.g., pipe) also depends on the chemical makeup
and concentration of antioxidants and other additives used in the formulation of PE.
After the stabilizers are functionally depleted, in general only a limited useful lifetime
is left for the part.
9.6 Temperature Acceleration of SCG in PE
Under selected normalized stress levels and an elevated temperature, cracks grow faster
and shorten PE lifetime in brittle fracture. However, at a certain combination of load
and temperature, a transition from a continuous SCG to discontinuous, stepwise propagation has been reported [20, 30]. The change of SCG mechanisms is related to PE’s
ability to form a stable process zone (PZ) consisting of cold drawn PE microfibers and
micro-membranes in front of the crack. A high hydrostatic tension component of the
crack-tip stress field and PE propensity to cold drawing plays a major role in PZ formation and growth. There are three main stages of PE brittle fracture: 1) crack initiation, 2) slow crack growth, and 3) crack instability and transition to either rapid crack
propagation (RCP) or ductile rupture. Note, the third stage occurs relatively fast, and
the lifetime of the PE component mainly consists of the time prior to crack initiation
and the duration of stable, slow crack growth.
Crack initiation is the most uncertain and the least studied stage of PE fracture.
Under fatigue and creep loading, crack initiation is an ultimate outcome of a slow process of damage accumulation on sub-micro and micro scales. The process progresses
faster in a vicinity of preexisting micro defects such as cavities, gels and/or foreign particles, agglomerates of additives resulting from poor mixing, and others. For an existing
field of material imperfections, the crack initiation time strongly depends on service
(or test) conditions such as the stress level, strain rate, temperature, and potentially an
aggressive environment. A large scatter, commonly observed in PE lifetime, is primarily associated with the random geometry, chemical composition, and the location of
the defect triggering the crack initiation. In contrast, the crack initiation in degradation driven brittle fracture (Mode III failure) does not require the presence of a defect
[31–34]. A thin surface layer of PE component that is in direct contact with a chemically aggressive environment undergoes degradation. Chemical degradation of semicrystalline PE is manifested in a reduction of molecular weight (MW) that often leads
to an increase of crystallinity and therefore PE density. The densification of the surface
layer attached to the interior unchanged material results in buildup of residual stress:
tensile stress in the degraded surface layer and compression in the interior domain.
The degradation of the surface layer also leads to material embrittlement; i.e., dramatic
reduction of PE toughness. As the degradation progresses, a combination of increasing
tensile stress and decreasing toughness of degraded PE leads to multiple crack initiations within the degrading surface layer [31–34]. The presence of a defect within the
surface layer may accelerate crack initiation.
The formation of a crack dramatically changes the stress field around it. According
to the linear elastic solution, there is a singular stress field in the vicinity of a crack tip
Deformation and Fracture Behavior of Polyethylene 327
[35, 36]. It is expressed in polar coordinates {r,q }with the origin at the crack tip and
polar axis as an extension of the crack line:
K
(r , )
r
f( )
(9.9)
The wave above a symbol like indicates that it is a stress tensor, a complex object
characterized by six independent stress components; r is the distance from the crack
tip and is the angular coordinate. In Equation 9.9, when r 0 all stress tensor components become unlimited; i.e., the stress tensor
. The coefficient K is called
stress intensity factor (SIF) and f ( ) is a universal angular distribution of stress tensor
components that is independent of crack size, load, and specimen geometry. Only SIF
contains the information of loading, crack, and specimen geometry. Equation 9.9 is an
idealization that is based on a few assumptions such as: a crack as an infinitely sharp
cut, the material follows the linear elastic stress-strain relations (material linearity),
and the displacements and strain are small (geometrical linearity). The assumptions
are quite restrictive in application to PE. However, the idealization is justified when
it is applied for evaluation of energy variation due to small crack extension, the socalled energy release rate (ERR). Indeed, in a linear elastic solid ERR, G1, is expressed
in terms of SIF: G1 = K2/E (E¢ is Young’s modulus for either plane stress or plane strain
conditions). In a nonlinear material like PE, ERR is commonly evaluated using a pathinvariant integral representation, J1. Direct calculations suggest that for many practical
problems G1 serves as a reasonably good approximation of J1.
An equation of SCG relates the crack growth rate cr with a crack driving force ( cr
stands for the crack length, a unit specimen thickness is assumed). Following conventional thermodynamic considerations a crack driving force (CDF) Xcr is introduced as
the derivative of Gibbs’ potential with respect to crack length cr : X cr
G / cr [37–
39]. Thus, the CDF is expressed as the difference between the elastic energy release rate
per unit crack area J1 (or G1 in the linear approximation) and the energy 2g required to
form that unit crack area (a specific fracture energy): J1 – 2 [38, 39]. The condition of
vanishing CDF is the celebrated A.A. Griffith’s crack equilibrium condition for linear
elastic solid [40, 41]:
G1
2
G1C
(9.10)
In linear elastic fracture mechanics (LEFM) the critical value G1C of ERR is widely
used as the fracture toughness parameter [36, 41]. Another commonly used measure of
fracture toughness is KC, with K C
E G1C . In LEFM, the fracture toughness parameter G1C (or KIC) is assumed to be a material constant. For a material with nonlinear elastic properties J1C should be used instead of G1C However, the assumption that J1C = const.
does not apply to a nonlinear material like PE, since J1C is observed to be dependent on
specimen geometry, loading rate, and most importantly crack size [36, 39]. It is related
to the fact that in PE a process zone (PZ) is formed around and in front of the propagating crack. A part of the PZ in front of the crack is referred to as the active zone (AZ),
since there are active processes of cold drawing and creep of oriented material within it.
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Handbook of Industrial Polyethylene and Technology
lcr
lAZ
Figure 9.11 SEM micrograph of crack and a PZ side view in PE, 100x magnification [39].
The AZ has a narrow wedge shape and extends along the crack plane. It consists of thin
microfibers and membranes drawn from the original homogeneous material. Creep
and ultimate rupture of drawn microfibers within the AZ results in crack extension
through the mid plane of the AZ and form a crack with “hairy” surfaces. The integral
effect of energy absorption by cold drawing of microfibers and membranes within the
AZ and by creep of the drawn material prior to its ultimate rupture is the origin of PE
SCG resistance; i.e., fracture toughness. Thus, the crack growth is closely coupled with
formation and evolution of the AZ in front of the crack. A coupled system of the crack
and the AZ is referred to as the crack layer (CL) [38, 39]. A typical appearance of a CL
in PE is shown in Figure 9.11 [39]. In Figure 9.11, a relatively light (white and gray)
domain of length cr with well visible ruptured fibers shows a crack (displacement discontinuity) with “hairy” upper and lower “surfaces.” There are sharp boundaries that
separate the conically shaped cold drawn PE fibers from the original material (the dark
domains) above and below the crack. The crack front is easily identified by a sharp
increase of the crack opening in comparison to that within the AZ, where oriented PE
fibers bridge the upper and lower faces along the AZ length AZ . The total CL length is
AZ
L cr
. The AZ width is not an independent parameter, since it is determined by
a) PE natural draw ratio n and b) by the displacement field computed along the boundary of the entire CL domain cutoff [20]. Thus, cr and AZ are the only geometrical
characteristics of CL. Rupture of the oriented PE within the AZ results in the crack
advance into AZ and therefore reduction of AZ length. In response, the AZ growth
away from the crack increases the AZ length AZ . Thus, CL propagation takes the form
of the AZ and crack-tip movements in material coordinates. The driving forces for
such processes are defined following the thermodynamic formalism mentioned above
[37], and they are expressed as the derivative of Gibbs potential with respect to AZ
and cr respectively: the AZ driving force X AZ
G / AZ , and the crack driving force
cr
cr
X
G / [38, 39]. Both driving forces have similar expressions:
tr
X
AZ
AZ
J
AZ
1
1
(x ) |x cr ,
where J1 is the elastic energy release rate associated with the AZ growth into the original PE domain; stands for displacement discontinuity within CL and (x )/ 1 |x cr
is the thickness of original PE from which the AZ was drawn at the tip of the crack. The
crack driving force has a similar expression X cr J1cr 2 . Despite the similarity of the
AZ
two expressions, XAZ and Xcr are essentially different. J1 is the elastic energy release
Deformation and Fracture Behavior of Polyethylene 329
due to AZ growth presented by a path independent contour integral with the entire
cr
integration path outside of CL domain, whereas J1 is the elastic energy release due to
the unit crack advancement into the AZ also presented by a path independent integral
with an integration path strictly confined within AZ. The tr in XAZ stands for the specific energy of the original PE transformation into the drawn, oriented PE, in contrast
with 2 , which represents the work done on oriented PE rupture under creep.
Conventionally the SCG kinetic equation is formulated as an empirical relation
cr
between crack growth rate cr and SIF KI, or in a more general case, ERR G1 or J1
[39, 42–44]. However, in the case of the CL system in PE we have two dynamic variAZ
. Therefore, the kinetic equations of fracables: the crack and AZ lengths cr and
ture propagation imply functional relations between the crack and AZ growth rates
cr
and AZ on one side and corresponding driving forces Xcr and XAZ on the other
side. Adopting classical Onsager’s arguments for proportionality between thermodynamic fluxes and corresponding forces for a relatively small deviation from equilibrium
and considering the irreversibility of crack and AZ growth, the following system of CL
kinetic equations has been studied [39, 45]:
cr
AZ
k1 X CR ( cr , AZ ), if X CR
AZ
k2 X (
cr
,
AZ
), if X
AZ
0, and
cr
0, and
AZ
0, if X CR
0, if X
AZ
0
0
(9.11)
The kinetic equations (9.11) are supplemented by initial conditions
and a law of fiber creep that determines decay with time of the
(0) cro ; AZ (0) AZ
o
specific fracture energy 2 (t). The CL driving forces XCR and XAZ are nonlinear functions of crack and active zone lengths and depend on material parameters and tr. To
evaluate such material parameters it is convenient to use a specimen geometry where
ERR does not depend on the variable crack length. It also helps to study CL growth
mechanisms and kinetics by reducing the number of variables. Tapered double cantilever beam (TDCB) specimens have been used for crack growth studies in metals
[46]. However, a TDCB specimen made of PE displays a very large deformation that
compromises the accuracy of basic fracture parameters evaluation. Another specimen
geometry, a stiff constant K (SCK) specimen, more rigid than TDCB and thus more
suitable for studies of SCG in PE, has been developed [20]. In the SCK specimen, SIF is
independent of the crack length in a certain range of crack length.
Different mechanisms of SCG in PE, specifically continuous and discontinuous
or stepwise SCG, have been consistently observed under constant load and temperature on various specimen geometries [20, 47–50]. The different mechanisms of SCG
are illustrated in Figure 9.12 by CL side views observed in the creep (constant load)
tests conducted on SCK specimens at the same temperature (80 °C), but different load
levels indicated by constant SIF values shown on the micrographs [39]. A continuous CL growth (upper micrograph) with a wedge-shaped PZ takes place at a low load
(KI = 10N/mm3/2), whereas discontinuous, pulsating PZ is observed at a load 1.8 times
higher (KI = 18N/mm3/2). The corresponding fracture surfaces observed in the same
tests are shown in Figure 9.13 [39]. In continuous SCG (at low load) there is a uniformity of fracture surface in the direction perpendicular to the crack growth direction
and graduate reduction of the random pit size in the direction of crack growth (toward
cr
330
Handbook of Industrial Polyethylene and Technology
KI = 10 N/mm3/2
KI = 18 N/mm3/2
Figure 9.12 SEM micrographs of typical CL side views in continuous (upper) and discontinuous (lower)
SCG in SCK specimens at different loads presented by SIFs (adopted from [39]).
Figure 9.13 SEM micrographs of typical fracture surfaces in continuous (left) and discontinuous (right)
SCG at low (KI = 10N/mm3/2) and high (KI = 18N/mm3/2) applied loads; both tests conducted at 80 °C
(adopted from [39]).
the crack front). A discontinuous crack growth depicted by the “pulsating” CL side view
(lower micrograph in Figure 9.12) is manifested by striations on fracture surface (right
micrograph in Figure 9.13). Monotonic and stepwise load point displacement (LPD)
versus creep time curves confirm the continuous and discontinuous SCG mechanisms
independently on side views and fracture surface observations [39].
A continuous SCG under creep conditions is expected, since there is no variation
of load and/or temperature. In contrast, the discontinuous (stepwise) growth pattern
is a surprise, since there are no external factors that may cause such behavior. The discontinuous mode of fracture growth is generated by the crack and the AZ interaction. A simple mathematical model of such interaction expressed as the system of CL
growth equations (9.11) captures the essence of such interaction. The solution of the
nonlinear system of equations (9.11) is not unique and suggests various scenarios of CL
evolution such as continuous, discontinuous, and transient from continuous to discontinuous growth. The model predictions are confirmed by experiments and field observations. A set of SCG tests conducted with identical SCK specimens led to formulation
of a fracture mechanism map (FMM) in terms of normalized SIF and temperature, as
Deformation and Fracture Behavior of Polyethylene 331
K/( cW1/2)
1.5
RCP and/or
Ductile failure
1.0
Discontinuous SCG
0.5
Continuous SCG
0.0
20
40
60
Temperature (°C)
80
Figure 9.14 Fracture mechanism map for pipe grade PE in normalized SIF and temperature (°C)
coordinates (adopted from [39]).
shown in Figure 9.14 [30, 49, 50]. The FMM summarizes the observations illustrated
by Figures 9.12 and 9.13. Specifically, there are domains in SIF-temperature coordinates with different SCG patterns: continuous SCG for all temperature ranges at low
SIF (normalized KI around 0.5 and below), rapid crack propagation (RCP) or ductile
rupture at high SIF (normalized KI is about 1 and above) also for all temperature ranges,
and a discontinuous SCG at normalized SIF values in between, as shown in Figure 9.14.
FMM is equipped with photos that illustrate SCG pattern. A typical sinusoidal RCP
trajectory for pressure pipes is illustrated in the upper photo.
Continuous and discontinuous SCG mechanisms are closely associated with different SCG kinetics: continuous propagation is much more “efficient” than a discontinuous one. For example, to compare the crack growth rate one can correlate an average
cr
cr
crack rate
with SIF K using for simplicity a power law
AK n [43, 44].
Then, the power “n” for continuous SCG in PE studied is found to be 2 to 3 times larger
than that for a discontinuous growth [49, 50]. Larger power n means higher acceleration of crack growth and the shortening of the PE lifetime. Thus, the application of
existing standards, such as ASTM F1473, ASTM D2837, and ISO 9080, for estimation
of PE lifetime at ambient temperature on the basis of elevated temperature testing may
lead to an overestimation of the safe service time with corresponding consequences.
The different mechanisms and kinetics of SCG is mainly controlled at the microscale
by creep deformation and fracture of the cold drawn fibers within the AZ. Continuous
SCG results from the rupture of micro fibers/membranes of AZ during or shortly after
completion of necking. Rupture of AZ material is the mechanism of crack growth in
PE. In contrast with that, when the fibers/membranes survive the cold drawing process
and reach a quasi-equilibrium coexistence with the surrounding original PE, a creep of
drawn material may take a sufficiently long time before fracture (the creep strain rate
of drawn PE is significantly lower than that of the original). During creep of drawn
material the PZ size remains almost the same. Then, after a certain time, micro fibers/
membranes break, leading to a spontaneous crack advance almost to the end of the AZ.
Thus, the creep of drawn material within the AZ is the origin of the stepwise, discontinuous SCG in PE.
332
Handbook of Industrial Polyethylene and Technology
Recognizing that continuous versus discontinuous SCG patterns are determined by
brittle versus ductile failure of microfibers within the PZ, the continuous-discontinuous
SCG transition has been referred to as ductile-brittle transition of the second kind, in
short DBT2, by analogy with conventional ductile-brittle transition (DBT) in impact
toughness [30, 39]. The corresponding DBT2 transition temperature (TT2) depends on
SIF. The relation between SIF (representing applied load) and TT2 is shown by the lower
boundary of the discontinuous SCG domain on the fracture mechanism maps [49, 50].
In summary, a transition in the mechanism and kinetics of SCG, DBT2, presents a
limitation for employment of an elevated temperature as an accelerating parameter for
brittle fracture. An extrapolation of SCG data from one domain of FMM into another
across the DBT2 boundary is not acceptable, since it violates the similarity of SCG processes. Therefore, the commonly used acceleration technique based on extrapolation of
1 to 1.5 year long test data at elevated temperatures into 50-year or 100-year life span
in service condition is suspect. It is called an “empirical” method. However, there is no
experience of using PE for 100 years, since there was no engineering PE a century ago.
Moreover, the newest grades of PE possess superior properties in comparison with the
old ones, but have a relatively short service record. In addition, there is a silent assumption that after decades of service the PE retains the same properties as that when tested
for a year or two. The material aging process (chemical degradation and/or physical
aging) commonly takes place over time and should be taken into consideration. It
seems the “empirical” extrapolation method works well as a marketing tool rather than
a solid engineering method of lifetime assessment.
An alternative approach to PE lifetime assessment is based on a) a fundamental understanding of underlying mechanisms of PE deformation and fracture, b) the
experimental evaluation of basic material parameters determining PE behavior, and c) a
sound physical model for the time to fracture (lifetime). For example, understanding
of the relation between delayed necking and PE pressure pipe ductile failure combined
with a symmetry between yield stress and draw stress dependency on the logarithm of
strain rate on one side and the logarithm of failure time dependency of applied stress
on the other, leads to a simple predictive model for PE lifetime in ductile failure using a
relatively short testing program. Similarly, a fundamental understanding of the connection of creep of original and drawn PE with mechanisms and kinetics of brittle fracture
propagation in PE presented in the form of the system of kinetic equations for crack
and AZ growth (9.11) leads to numerical simulation of brittle fracture propagation and
prediction of PE lifetime in brittle failure [45]. Thus, a set of tests for experimental
determination of basic PE properties combined with a numerical modeling of ductile
and brittle fracture substitute for the empirical procedure of ISO 9080 and similar standards [45]. It is important to keep in mind that PE properties may change over time due
to the aging process. A large amount of information on PE aging has been accumulated
in the vast literature on the subject. The goal is to incorporate the PE aging data into a
constitutive model of PE behavior and thus include the aging process into the computational method of PE lifetime assessment in brittle fracture. Parametric analysis readily
available in numerical modeling allows consideration of the effects of various service
conditions on PE life expectancy. The above suggestion is based on a large amount of
experimental data mainly related to HDPE. However, the proposed approach is well
applicable to various PE types, including LLDPE, LDPE, and metallocene PE.
Deformation and Fracture Behavior of Polyethylene 333
9.7 Conclusions
Polyethylene represents a family of polymers that can be very tough and sufficiently
strong to cover a wide range of applications from packaging to very demanding durable
structures. Design engineers for PE applications should also be aware of some specifics
of PE behavior, such as a very strong time dependency manifested in creep. There is
also a strong dependency of yield strength, draw stress, elongation at break and stiffness on temperature and applied strain rate. These dependencies are associated with
polymer chain mobility within amorphous phase as well as regions adjacent to lamellae
with very different time scales. Qualitatively it is well understood, but practical characterization is done on the phenomenological basis. In this chapter we focused on PE
necking, since it is a main contributor to ductile PE failure by large deformation as well
as to the brittle fracture by crack growth that limits the service life of PE. Necking is a
manifestation of instability of the original PE that leads to a transition to a more stable
oriented (drawn) state. It explains a direct connection of necking with PE ductile failure. The time dependency of PE necking and therefore ductile failure is associated with
delayed necking phenomenon. Next, we focused on accelerated testing methods for PE
lifetime. There was a brief discussion of commonly used failure accelerating methods
and their limitations. An alternative approach for prediction of PE lifetime in ductile
mode of failure that agrees well with existing empirical predictions but requires a much
shorter testing program was also discussed. Next, the mechanism and kinetics of brittle
fracture were discussed. It is demonstrated that PE resistance to SCG is directly related
to a microscale cold drawing in a vicinity of the crack tip. Thus, resistance to crack
growth in PE is linked to PE necking. Then it is shown that there are two different
mechanisms of SCG, continuous and discontinuous, controlled by the ductile-brittle
behavior of drawn microfibers in front of the crack (ductile-brittle transition of second
kind (DBT2)). This is captured by the formulation of a fracture mechanism map. There
is a very large difference in continuous and discontinuous crack growth kinetics. The
existence of DBT2 limits application of commonly used temperature accelerated SCG
studies for lifetime prediction. It presents a warning that presently used standards may
lead to an overestimation of time to failure that is contrary to a conservative design
philosophy.
References
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Handbook of Industrial Polyethylene and Technology
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19. Zhou, Z., Chudnovsky, A., and Sehanobish, K., Evaluation of Creep Property and Lifetime
in Ductile Failure of PEs from Short-Term Tests, SPE-ANTEC Tech. Papers, 51, 784, 2005.
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21. Chudnovsky, A., Kim, A., Chen, T.-J., Sehanobish, K., Bosnyak, C.P., and Kao, C.I.,
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Necking Process in Polyethylene, SPE-ANTEC Tech. Papers, 45, 3399, 1999.
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Part 2
PROCESSING AND FABRICATION
OF POLYETHYLENE
10
Single-Screw Extrusion of Polyethylene Resins
Mark A. Spalding
The Dow Chemical Company, Midland, Michigan, USA
Contents
10.1 Introduction............................................................................................................................. 339
10.2 Screw Sections and Processes ................................................................................................ 342
10.3 Common Problems ................................................................................................................. 348
10.3.1 Gels............................................................................................................................. 348
10.3.2 Rate Restriction at the Entry of a Barrier Flighted Melting Section ................. 352
10.3.3 Nitrogen Inerting ..................................................................................................... 352
10.4 Process Assessments ............................................................................................................... 353
References ............................................................................................................................................ 356
Abstract
Single-screw extruders are the machine of choice for most polyethylene (PE) resin conversion
processes. Their basic operation is relatively simple, and they provide a low-cost process for
plasticating and pumping molten resin to downstream shaping operations. The design and
troubleshooting of process lines, however, can be very complicated. This chapter provides an
overview of the single-screw extrusion process, screw channel design basics, and provides
troubleshooting remedies for the most common problems that a converter would encounter.
Keywords: Single-screw extrusion, screw, plasticating, extruder
10.1 Introduction
Single-screw extruders are one of the most common plastics processing equipment
used in the industry. They are capable of melting unfilled PE resins and pumping them
at rates up to about 2200 kg/h. The device is shown in Figure 10.1 for a relatively small
63.5 mm diameter extruder. The extruder is configured with a barrel and a tight fitting
screw (not shown) positioned inside the barrel. The rotation of the screw conveys the
solid PE feedstock forward from the hopper. The feedstock is then melted, pumped,
sometimes mixed, and then discharged. As shown by Figure 10.1b, a motor, gearbox,
Corresponding author: maspalding@dow.com
Mark A. Spalding and Ananda M. Chatterjee (eds.) Handbook of Industrial Polyethylene and Technology, (339–356)
© 2018 Scrivener Publishing LLC
339
340
Handbook of Industrial Polyethylene and Technology
Hopper
Control
panel
Pressure
sensor
Barrel
Gearbox
Belt
sheaves
Feed
casing
Motor
(a)
Heaters
Air cooling
fans
(b)
Figure 10.1 Photograph (a) and a schematic (b) for a 63.5 mm diameter single-screw extruder.
(Courtesy of William Kramer of American Kuhne)
and belt sheave system is typically used to convert the high-speed, low-torque power
from the motor to the low-speed, high-torque power needed to rotate the screw. Barrel
heaters and coolers are used to maintain the temperature of the barrel. Cooling water
is typically used to maintain the feed casing to temperatures less than about 50 °C.
A pressure sensor should always be positioned in the discharge stream to monitor process performance including process stability and to alert the operators for high-pressure
safety issues. Downstream shaping equipment is attached to the exit barrel flange.
The barrel is covered with heating and cooling units such that the temperature of
the barrel can be controlled to a set point value. Typically the heating is via electrical
resistance heaters. Cooling can be from forced air fans or from water. The heaters are
typically on full during the initial heating phase prior to startup. During operation, the
heaters (or coolers) typically cycle on-off to maintain the temperature as essentially all
energy for the melting and pumping processes come from the motor via the rotating
screw. The operating temperatures used depend on the application and the resin.
Plasticating single-screw extruders are currently built with barrel diameters between
20 mm and 203 mm; smaller and larger machines have been built. The term plasticating is used here to indicate that the feedstock is a solid resin, typically pellets, and the
plasticating process converts the solids to a molten material. Other feedstock forms
are commonly used including powders, ground film recycle, and ground sheet. Typical
operating rates for plasticating single-screw extruders running unfilled PE pellets are
provided in Table 10.1. Actual rates obtained on a machine, however, depend on many
factors including screw geometry, barrel length, discharge pressure, the motor drive
system, fillers incorporated into the resin, and the downstream process. At rates higher
than about 2200 kg/h, a single-screw extruder is no longer the machine of choice for
plasticating PE resins. For these higher rate applications, twin-rotor machines are often
preferred including twin-screw extruders and continuous mixers.
The screw and process are designed to meet the rate and discharge temperature
requirements for the downstream application, as shown in Table 10.2. For example,
thick wall plastic pipe is often produced using a high viscosity high density PE (HDPE)
resin. Here the process is designed such that the extruder is running at maximum rate
and typically at a discharge temperature no higher than about 230 °C. In order to meet
Single-Screw Extrusion of Polyethylene Resins
341
Table 10.1 Typical rates for unfilled PE resins at a screw speed of 100 rpm as a function of
barrel diameter.
Barrel diameter, mm
Typical rate, kg/h
Specific rate, kg/(h rpm)
20
3
0.03
60
60
0.6
80
140
1.4
100
280
2.8
120
500
5.0
150
1000
10.0
200
2200
22.0
Table 10.2 Application requirements for a 100 mm diameter barrel plasticator for unfilled
PE resins.
Application
Discharge
temperature, Rate, Screw speed, Meter channel Specific rate,
°C
kg/h
rpm
depth, mm
kg/(h rpm)
Pipe
230
270
71
7.0
3.8
Blown Film
230
250
74
6.4
3.4
Cast Film
250
240
90
5.0
2.7
Injection Molding
250
250
100
4.5
2.5
Paper Coating
300
300
190
3.0
1.6
these requirements, the screw channels are relatively deep and the screw rotation speed
is maintained low such that viscous energy dissipation is minimized. For a paper or foil
coating process, the PE resin must be applied very hot and in the range of 280 to 310 °C.
Here the screw channels are relatively shallow and the screw speed is maintained very
high such that the level of viscous dissipation is high. The specific rate is defined as the
rate divided by the screw speed, and it is a good indicator for the relative discharge
temperature for machines with the same diameter and running the same resin. That
is, a screw with deeper channels will operate at a higher specific rate and a lower discharge temperature. Many other processes are designed with process temperatures and
specific rates that are intermediate to these extreme applications. For example, cast film
processes often operate at intermediate specific rates and with discharge temperatures
near 250 °C.
As discussed next, the metering channel geometry controls the specific rate of a
screw design for smooth bore extruders. The actual design of the metering channel is
beyond the scope of this writing, but the main geometry component for setting the rate
and energy dissipation is the depth of the channel. As shown in Table 10.2, the specific
rate decreases and the discharge temperature increases as the metering channel depth
decreases. For machines with screw diameters of 100 mm, a blown film process would
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Handbook of Industrial Polyethylene and Technology
typically be designed with a metering channel depth of about 6.4 mm while a screw for
paper coating would have a meter channel depth that is considerably shallower at about
3.0 mm.
10.2 Screw Sections and Processes
The extruder must perform at least three different processes, and often several more
depending on the application. For example, at a minimum the screw must convey the
PE solids from the hopper to the melting section of the screw. The melting section
must be able to convert the solid pellets to a molten liquid. The molten PE resin is then
pumped or metered to the downstream sections of the process. The metering section is
where pressure can be generated if needed. Dispersive mixers are often used to homogenize the discharge before it exits the extruder. These mixers are typically positioned in
the middle of the metering section although other locations are practiced commercially.
There are two main classifications of plasticating, single-screw extruders: smoothbore feed machines and grooved-bore feed machines. For grooved-bore extruders, the
grooves are placed in the feed section of the barrel where they are the deepest at the
feed opening and taper in depth over the next 3 to 4 diameters. The barrel section
downstream of the grooved feed section is smooth. Smooth-bore extruders are not
designed with grooves. The designs of the screws are very different for the machines.
For smooth bore extruders, the metering section of the screw controls the rate while the
solids conveying section in the grooved part of the barrel controls the rate for groovedbore extruders. Rate calculation methods and the fundamentals for the processes were
detailed by Campbell and Spalding [1].
The single-screw plasticating process starts with the mixing of the feedstock materials. Typically, several different feedstocks are added to the hopper such as fresh resin
pellets, recycle material, additives, and a color concentrate. The recycle material typically comes from the grinding of edge trim, web material from a thermoforming processes, or off-specification film and sheet. Next, the feedstock flows via gravity from the
hopper through the feed throat of the feed casing and into the solids conveying section
of the screw. Typically this feed casing is cooled using water. The feed section of the
screw is typically designed with a constant depth and is about 4 to 8 barrel diameters
in axial length. Directly after the solids conveying section in a smooth bore extruder
is a section where the channel depth tapers to a shallower depth metering section. The
tapered depth section is commonly referred to as the transition or melting section.
Most of the melting process and the primary mixing occur in the transition section. In
general, the metering section is also a constant depth, but many variations exist where
the channels oscillate in depth. The metering section pumps and pressurizes the material for the downstream unit operations including static mixers, screen filtering devices,
gear pumps, secondary extruders, and dies. A conventional style screw for a smooth
bore extruder is shown in Figure 10.2. The total length of the extruder screw and barrel
is typically measured in barrel diameters or as a length-to-diameter (L/D) ratio. Section
lengths are often specified in barrel diameters or simply diameters.
A conventional single-flighted screw is shown in Figure 10.2. This screw has a single
helix wound around the screw root or core. Multiple flighted screws with two or more
Single-Screw Extrusion of Polyethylene Resins
Flight
Shank
Solids conveying
section
Flight
tip
Transition or
melting section
Screw
root
343
Tip
Metering
section
Pocket or
flight start
Figure 10.2 Schematic of a typical single-flighted screw. (Courtesy of Jeff A. Myers of Robert Barr, Inc.)
helixes started on the core are very common on high-performance screws. Screws with
barrier-flighted melting sections are very common in film processes because they provide high rates with lower discharge temperatures. Barrier melting sections have a secondary barrier flight that is located a fraction of a turn downstream from the primary
flight, creating two flow channels in the transition section; a solids melting channel and
a melt conveying channel. Barrier flighted sections will be discussed in more detail later.
Many high-performance screws [1] have two or more flights in the metering section of
the screw. The screw is rotated by the shank using either specially designed splines or
by keys with rectangular cross-sections. The mathematical zero position of the screw
is set at the pocket where the screw helix starts. Most extruder manufactures rotate the
screw in a counterclockwise direction for viewers positioned on the shank and looking
towards the tip. This rotation convention, however, is not standard.
The flight is a helical structure that is machined into the screw and extends from the
flight tip to the screw core or root. The flight has a width at the flight tip called the flight
land. The small clearance between the flight land and the barrel wall minimizes the flow
of resin over the land. This clearance is typically 0.001 times the diameter. The polymer
that does flow between the clearances supports the screw and centers it in the barrel.
The radial distance between the flight tip and the screw root is referred to as the local
flight height or channel depth. The feed section usually has the largest channel depth,
and provides the largest cross-sectional volume in the screw. The deep channel conveys
the relatively low bulk density feedstock pellets into the machine via the motion of the
helix. The feedstock is conveyed forward into the transition section or melting section of the screw. The transition section increases in root diameter in the downstream
direction and thus the channel depth is decreasing. Here, the feedstock is subjected to
higher pressures and temperatures, causing the feedstock to compact. As the material
compacts, its bulk density can increase by a factor of nearly 2 or more. As the feedstock
compacts, the entrained air between the pellets is forced backwards and out through
the hopper. For example, a pellet feedstock such as low density PE (LDPE) resin can
have a bulk density at ambient conditions of 0.58 g/cm3 while as a full compacted solid
bed in the transition section the density will approach 0.92 g/cm3 before melting starts.
Thus for every unit volume of resin that enters the extruder, about 0.4 unit volumes of
air must be expelled through the voids in the solid bed and then discharged through the
hopper. The transition section is where most of the polymer is converted from a solid
to a fluid. The fluid is then conveyed to the metering section where the molten resin is
pumped to the discharge opening of the extruder. In general, the metering section of
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Handbook of Industrial Polyethylene and Technology
a conventional screw has a constant root diameter and it has a much smaller channel
depth than the feed section.
The compression ratio sets the depth of the feed section channel of a smooth-bore
extruder. The compression ratio for most screws is simply the depth of the feed channel
divided by the depth of the metering channel. In general, the feed section must be able
to convey resin at a rate high enough to maintain the downstream sections of the screw
full of resin and pressurized. For PE pellet feedstocks, the compression ratio should be
between 2.4 and 2.8. If a high level of low bulk density recycle material is added with
the pellets, the compression ratio may need to be increased [1]. The compression rate
is the rate of change of the channel depth in the transition section of the screw. Since
the compression ratio sets the feed channel depth, the compression rate sets the axial
length of the transition section. The compression rate for PE resins is generally between
0.0025 and 0.0045; 0.0035 is typical. The compression ratio and compression rate are
calculated as follows for conventional flighted transition sections:
C
R
(H f
Hf
(10.1)
H
H )sin b
ML
tan b
L
Db
(10.2)
(10.3)
where C is the compression ratio, Hf is the channel depth of the feed section, H is the
depth of the metering channel, R is the compression rate in the transition section, M
is the number of turns in the transition section, b is the helix angle at the barrel wall,
Db is the diameter of the barrel, and L is the lead length; ML is the axial length of the
transition section. The lead length is the axial distance for a flight to make a complete
revolution. The lead length is often equal to the diameter of the barrel; i.e., a square
pitched screw.
As previously discussed, the majority of the solids are melted in the transition section of the screw. The melting mechanism is fairly complex with the largest portion of
the resin melted between the solid bed formed in the channel and the barrel wall, as
shown by Figure 10.3. The red and white colored cross-sectional view of Figure 10.3
was obtained from a Maddock solidification experiment. The Maddock solidification
experiment [2] was performed by adding a low level of colored pellets to the feedstock
resin. When the color appeared at the die, screw rotation was stopped, full cooling was
applied to the barrel, and the resin was solidified in the screw channels. Next, the screw
and solidified resin were pushed out of the barrel and the solidified resin was removed
from the screw. The white colored section of the photograph is the solid bed while
the red sections were molten when the screw was stopped. As shown by the view, a
relatively large level of mixing occurs during the melting process [3]. The large portion
of the molten resin is referred to as the melt pool. Details of the melting process are
described elsewhere [1].
Single-Screw Extrusion of Polyethylene Resins
Melted and
partially mixed
resin (melt pool)
Compacted natural pellets
(solid bed)
345
Melt film
Color concentrate
(solid)
Hopper
Discharge
Solids
conveying
Melting
Metering
Figure 10.3 Schematic of the melting and mixing process that occurs during the plastication of natural
acrylonitrile butadiene styrene ABS (white) resin with a red color concentrate resin using a conventional,
single-flighted screw.
Maddock
mixer
Solids conveying
section
Barrier flighted
melting section
Metering
section
Figure 10.4 Schematic of a Steward barrier flighted screw with a downstream dispersive Maddock mixer.
(Courtesy of Jeff A. Myers of Robert Barr, Inc.)
The transition section shown in Figure 10.2 is a conventional single-flighted design.
These designs are still used for film operations, but barrier flighted melting sections are
much more common. Barrier flighted melting sections will typically provide higher
rates, lower discharge temperatures, and a more stable discharge pressure. Barrier
flighted melting sections are constructed by positioning a second flight (or barrier
flight) in the transition section such that the solids are maintained on the trailing side
and the molten resin on the pushing side. A schematic of a barrier flighted screw with
a Maddock-style mixer is shown by Figure 10.4, and a schematic of a cross-section of a
barrier melting section is shown in Figure 10.5. The resin that is melted near the barrel
wall is conveyed across the barrier flight and collected in the melt conveying channel. The key design parameters include the position of the barrier flight, the depths
of the channels, and the undercut clearance on the barrier flight. The undercut clearanc