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Extrusion Processing for Manufacture of Low-Density, Fine-Celled Polypropylene Foams
~ a n i E. Naguib
A thesis submittd in conformity with the requirements for the Degree of Doctor of Philosophy
Department of Mechanical and lndustrial Engineering university OF Toronto
@Copyright by B. E. Naguib 2001
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Extrusion Processing for Manufacture of Low-Density, Fine-Celled Polypropylene Foam
Hani E. Naguib Degree of DacCor of Philosophy, 2001
Department of Mechanical and Industrial Engineering University of Toronto
A continuous extrusion process for the manufacture of low-density, fine-celled
poiypropylene foams is presented. Due to its outstanding Functiond characteristics and low
materiai cost, polypropylene foarns have been considered as a substitute for other
thennoplastic foams in industriai appIications. However, only limited research has been
conducted on the production of polypropylene foarns because of the weak melt strength, and
no research has been conducted to investigate the mechanisms that govern the expandabiiity
of poIypropylene foams. This thesis presents the effective scrategies for increasing the
volume expansion ratio as welI as the mechanisms goveming the foam density of
polypropylene foams. The basic strategies taken in this study for the promotion of a large
volume expansion ratio of polypropylene foams are: (a) to use a branched materiai For
preventing ceIl coalescence; (b) to use a long-chah blowing agent with low diffusivity; (c) to
lower the melt temperature for decreasing gas Ioss during expansion; and (d) to optirnize the
processing conditions in the die for avoiding premature crysrdlization. The effects of
processing and materiais parameters on the foam morphoiogies of poIypropyIene materiais
were thoroughly studied using a single-screw tandem foam extrusion system. A careFuI
analysis of extended experimentd resuIts obtained at various processing conditions indicates
that the final votume expansion ratio of the extruded polypropylene foams blown with butane
is governed either by Ioss of blowing agent or by crystdlization of the polyrner matrix. By
tailoring the processing conditions in the die, ultra low-density, fine-celied polypropylene
foarns with very high expansion ratio up to 90-fold were successfuIly produced from the
branched poiypropyiene resins. Fundamental snidies have also been conducted to investisate
the effect of various processing and materials parameters on the thermodynamic, thermal and
meit fracture behaviors of polypropyIene melts with foaming additives that influence the ceIl
morphology of poiypropylene foams.
ln memory of my Aunt Aida, who passed away during my Ph.0. program.
1 know you still hear me and pray for me. 1 thank you for al1 the pars you took care of me, for your love, support and encouragement. 1 could not make it without you. You will aiways be in my
heart and my thoughts
In memory of my Mother who was and still is my inspiration in my life
70 my great Father to whom 1 owe al1 that 1 am and al1 that 1 am going to ber
and al1 that 1 can ever hope to be
Acknowledgments
1 would like to start by giving thanks to the beneficent and mercifuI God, my Lord
and Savior Jesus Christ, for he has covered me, supported me, preserved me, accepted me
ont0 him, had compassion on me, sustained me, and brought me to this hour.
1 would like to express my sincere gratitude to my supervisor Professor Chu1 B. Park
for providing guidance and encouragement throughout my iesearch. I will never forget the
support he gave at my most difficult times.
1 would Iike also to thank my Ph. D. thesis committee: Prof. Steve Bdke from
Department of Chemical Engineering, Prof Shaker Meguid, Prof Beno Benhabib €rom the
department of Mechanical and Industrial Engineering.
Special thanks to Prof. O. A, Aziz, Prof. M. Al Gammal, Prof. A. AbdeI Messih and
Mr. Albert Mickail for their guidance throughout my career and the time and effort spent.
My gratitude is extended to the Department of Mechanical and Industnal Engineering
at the University of Toronto for providing the University of Toronto Open Fellowships, the
Ontario Graduate Scholarship, as well as, the NSERC Postdoctoral fellowship. Also, i would
like to thank Boredis AG. Company, in Austria for their funding and support in this project.
I would also like to thank my colleagues in the Microcellular Plastics Manuhcturing
Laboratory for their help and Friendship over the p s t four years. They include Simon Park,
Dmitry Ladin, Dr. Valentina Padareva, Anthony Yeung, Amir Behravesh, Ghaus Rizvi,
Esther Richards, Patrick Lee, Linda Lin, Remon Pop-iiiev, Deepak Fernandes, Dr. Chris
Song, Dr. Sang Mae Lee, Dr. Yuejian Liu, Minhee Lee, Xiang Xu, Donglei KU, Wanlin Chen,
Chris Ozolins, Anjan, Xioyang Guan, Rehan Khan, Gumjan Guo, and Haiou Zhang.
Especially, sincere gratitude poes to Simon and Dmitry who shared common thoughts. 1
would like to thank al1 the undergraduate students that worked with me Erin Youn, Young fi,
David Allas, Suzan Oh, Won Park, Carri Li, Joyce Lam, Fames Koo, Brandon Lee, Michael
Lee, Ibrahim Abu Eisha, and others. Also, 1 wish to acknowledge the professiond technical
support from Mike Smith, Len Rooseman, Jeff Sansome and Dave Eisdaile in the Machine
TooI Laboratory, from Wendy and Amanda in the generai office, from Mary Rose, Teresa,
Dan and Sheila in the purchasing department, and from Oscar in the computer support. A
special thanks to Esther Richards for proof reading my thesis and for Dimtry Ladin,
Mohammed Attia, DongIei Xu, Linda Lin, Ghaus Rizvi, Xiang Xu, WanIin chen and Lianne
Ing for the time they spent for helping me at the final stage of the thesis.
iv
1 would like to thank my family al1 over the world; from Egypt: 1 would like to thank
my wife for her continuous efforts to provide me with her love and support, my great father,
my beloved brother Rami, his wife Mary and Daniel, to my beloved cousins: Sarni, Nagi and
their families, Shadi, and to a11 my aunts and uncles especidly Aunt Mary, Uncle Soria] and
Uncle Atef for their continuous support and encouragement. My family in Canada: my God
father Uncle Esmat, to my beloved cousins: Tarek, Kim, Mark, Magdy, Samer, Rosemary,
Chris and Sally. A special thanks to my parents in Canada Uncle Said and Aunt Mona for
their continuous care and love. My family in the States: my God mother Nadia for her
continuous love and support, my uncles and riunts: Osiris, Nela, Menes, Madelaine, Kamilia,
and my cousins: Mona. Marc, Sergio, Franco, and their families, Gehan and James.
1 would like to thank father Reuiss, father Georgeos, father Makari from St. Georges
Church in Toronto, father Amonios and father Messaeil form St. Mark church in Toronto. A
special thanks to the youth group in my church, who supported me with their prayers and
provide me with more love to rny Lord Jesus Christ.
1 also would like to thank many friends ihab and Lydia, Essam and Marianne, Adel and
Sally, Bassem and Dalia, Sameh and Gehan, Roger and Mira, Jules and Mary, Joseph
Armanios, Talaat and Mona, Nagui and Dalia, Dimitry Saad, and Alfred Mobayad.
Finally, 1 would like to praise my Lord with David the prophet and king by saying
I I O LORD. yorr have searched me and you knoiv me. YOU knoiv ivhen C sir and ivhen I rise; You perceive my rfroirgl~ts from afar. Yori discent my going otrr and my king down:
Yoir are faniriliur ivith al1 niy ivays. Before a word 13 on Iny tongrle You know ir cumplrrely. O LORD. Yori hem me in--bellinriund before: Yori have laid yoiir hand iipon me.
Such knoivledge is too wondefil fur me. tao lofty for me to airain. Where can I go fiom p u r Spirit? Wliere can Iflreji-om yuurpresence?
I f I p up to the heavens. You are there; iff make tny bed in the depths You are there. IfI rise on the wings of the daivn, iflsettle on rhe fur side of the sea, even tlrere your hand ivill giride me. your nghr hand will hold me fast.
IfI say, "Siirely the darkness will hide me and the light become nighr urotitrd me," even the darkness will nor br dark ro Yoti; the nigfrt will shine like the duy,
for darkness is as tighr to You. For You creared my inmost being: You knit me togerher in my mother's womb.
I praire You because I am fearfitlly and ivonderfulfy made; Your ivorks are rvonderfuf. I knorv rharfiill ivefi.
My frame ivas not hiddenfrorn Yori when 1 was made in the secret place. When I ivas ivoven together in the depths of the earrh. Your qves saw my rinformed body. Al1 the duys ordainecifor me iuere wiften in Your book before one of rhem came to be.
Hoiv precioiw to me are Your rhorighrs, O God! Hoiv vast fs the strm of them! Were 1 to cortnr them. they would ournumber the grains of sand
When I nwuke, I am $tilt with you "- Psalm 139 (1-19)
. . ABSTRACT ........ .. . . . .. . .. . .. .. .. .. . .. , . . . . . . . . .. . . . . . . . . .. . . . . . - . . . - . . .. - . . . . . . . . . . . . . . . . .. . . . . . . . . I I
. *. Dedication .......................................................................................................... 111
Acknowledgrnents ......................................................................................................... iv
Table of Contents ........................................................................................................... vi
List of Tables ............................~.~~......~....~...~..........~....................................................... xi * *
List of Figures ................................................................................................................... xu + * *
Nomenclature ................................................................................................................. xxlu
Chapter 1 Introduction
I , 1 Prearnble .......................................................................................................... 1
7 1.2 TherrnopIastics Foarns .............. .. .......................................... .......................... - 3 1.3 Low-Density, Fine-Celled Foam Processing ............................................. -
1 A Objective of the Thesis ................................................................................ 3
1.5 Overview of the Thesis .............................................................................. . 5
Chapter 2 Literature Review and Theoretical Background
2.1 Background on Polypropylene Foam Processing ............................................ 7
2.1.1 Cross-Linked Polypropylene Materials ...................................... 7
2.1.2 High-Melt-Strength Polypropylene Materials ................................. 8
2.1.3 Low-Density, Fine-Celled Polypropylene Foarns .............................. 9
2.2 Background on the Properties of Plastic Melts and Plastic/Gas Solutions ..... 13
2.2.1 PVT f roperties .......................................................................... 13 * * 3.2.2 Crystallization Behavior .................................................... 14
2.2.3 Melt Fracture ..................... .. ................................................... 16
2.3 Theoretical Background .............................................................................. 1 8
2.3.1 SoIution Formation .................................. . ...................................... 18
2.3.2 Nucleation ........................................................................................ 22
2.3.3 Cell Growth ..................................................................................... 26
2.3.4 Crystallization Kinetics ................................................................... 30
Chapter 3 Design & Construction of the Experimental Equipment
................................ ...... 3.1 Conceptual Design of Tandem Extrusion System .. 36
3.1.1 Introduction ...................................................................................... 36
3.1.2 Andysis of the Tandem Extrusion System for Foarn Processing .... 37
3.1.3 Detaiied Anaiysis and Furthet Decomposition oFFRs and DPs ...... 40
3.2 Detailed Design of the Tandem Extrusion Line ....................................... 4 1
3.2. i Overview of the System ................................................................... 41
.... 3.2.2 First Extruder in the Tandem Line ............................................. 42
3.2.3 Gris Injection Equipment ....................................... .. ................. 42
3.2.4 Second Extruder in the Tandem Line ............................................... 42
3.2.5 Gear Pump ........................................................................................ 43
3.2.6 Diffusion Enhancing Device ............................................................ 43
3.2.7 Heat Exchanger ................................................................................ 45 3.2.8 Filament Die ..................................................................................... 46
3.2.9 Coolin; SIeeve ................................................................................. 46
Chapter 4 Investigation of Fundamental Properties of Polypropylene
Materials with Foaming Agents
4.1 Introduction . ......................... .., ...................................................................... 52
4.3 Measurements of PVT Properties of Polypropylene Materials ....................... 52
4.2.1 Experimentai Equipment and Procedure ...................................... 55
. 4.2.2 Results and Discussion ...............,............+.................................. 58
............ 4.2.2.1 Effect of dissolved butane on the specific volume 58
...................... 4.2.2.2 Effect of branching on the specific volume 59
4.2.2.3 Effect of processing temperature and pressure on
the specific volume ............................................................ 59
4.2.2.4 Error Analysis ................................................................... 60
412.3 Conclusions ................................................................................... 61
4.3 Measurements of Thermal Be havior of Polypropyiene Materials .................. 61
.......................................... 4.3.1 Experimental Equipment and Procedure 62
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........................................................................ 4.3.2 ReguIar DSC Results 63
4.3.2.1 Effect of Branching ........................................................... 63
4.3.2.2 Effects of Foaming Additives ........................................... 64
4.3.2.3 Effect of Cooling Rate ....................................................... 65
4.3.3 High Pressure DSC Results .......................... .. .............................. 65
4.3.3.1 Effect of Hydraulic Pressure ............................................. 66
4.3.3.2 Effect of Dissolved N7_ ....................................................... 66
4.3.3.3 Effect of Dissolved CO2 ................................................... 67
.................................................................................... 4.3.4 Conclusions 68
..... 4.4 Measurements of the Onset of Melt Fracture of Polypropylene Mateciais 68
.......................................... 4 .4.1 Experimentrtl Equipment and Procedure 70
..................................................................... 4.4.2 Results and Discussion 72
4.4.2.1 Effect of branching on the criticat shear stress ................. 72
4.4.2.2 Effect of processing temperature on the critical
......................................................................... shear stress 73
4.4.2.3 Effects of foaming additives on the critical shew
.................................................................................. stress 73
........... 4.4.2.4 Effect of biowing agent on the critical shear stress 71
...................................................................................... 4.4.3 Conclusions 74
4.5 Summary ........................................................................................................ 75
Chapter 5 Production of Low-Density Fine-Cell Polypropylene Foams
5.1 introduction .................................................................................................... 95
5.2 Strategy for Promoting Low.Density, Fine-Ceiled Polypropylene
Foams ............................................................................................................. 95
5.3 Experimental Materiah .............................. .. ............................................... 98
..................................................... 5.3.1 SeIection of Polymeric Materials 98
........................................................... 5.3.2 Selection of Foaming Agents 99
................................................................................ 5.4 Experimentd Procedure 100
5 5 ResuIîs and Discussion .................................................................................. 101
5.5.1 Effect of Processing Temperature on Volume Expansion and
.................................................................................... Ce11 Density 101
5.5.2 Effect of Blowing Agents on Volume Expansion and Ce11
........................................................................................... Density 103
5.5.3 Effect of Nucleating Agents on Volume Expansion and Cell
........................................................................................... Density IO6
5.5.4 Effect of Long Chain Branching on Volume Expansion and
.................................................................................... Cell Density 108
5.5.5 Effect of Materials Blending on Volume Expansion and Ce11
........................................................................................... Density IO9
5.5.6 Effect of Re-Extruded Propylene Materials on Volume
.................................................... Expansion and CelI Density ... 1 1 I
5.5.7 Effect of Die Geometry on Volume Expansion and Cell
........................................................................................... Density 112
5.6 Statisticai Anaiysis of the results ............................................................ 113
5.7 Summary and Conclusions ............................................................................ 113
Chapter 6 Fundamental Mechanisms of Volume Expansion Behavior of
Polypropylene Foam Filaments
6.1 introduction ......................................... 140
................................. 6.2 Qualitative Modeling of Volume Expansion Behavior 141
6.2.1 Fundamental Mechanisms Governing Volume Expansion of
Polypropylene Foarns ................ ... ............................................. 14 1
6.2.1.1 Gas Loss ................................... .... ........................ 141
6.2.1.2 Crystallization ..................... .. .................................. 142
......... 6.2.2 Visualization of Expansion Behavior Using a CCD Came ra 143
................................................................... 6.2.2.1 System Setup 143
6-2-22 Expenmental Procedure .................................................. 143
6.2.2.3 Effect of Processing Temperature on the initial Expansion
.................. Rate and Final Diarneter of the Extudate 144
6.3 Theoretical Mode1 for Calculating the Expansion Ratio. Cell Size.
and CeIl Wall Thickness from the Obsewed Foarn Profile ..................... 145
................................................... 6.3.1 Development of a Theoretical Mode1
6.3.2 Determination of Expansion Ratio. Cell Size and Ce11 Wall
Thickness of Extruded Foarns frorn the Observed Profiles 149
.......................................................................... 6.4 Surnmary and Conclusions 150
......................................................................... Chapter 7 Summary and Conclusions 163
..............* ..........**..*.*..*****...**.. Chapter 8 Recommendations and Future Work ... 167
References ..................................................................................................................... 168
..................................................................................... ....................... Appendix 1 ... 181
........................................................................................ Appendix 2 ........................ .. 186
Table 4- 1
Tabie 5.1
Table 6.1
Table A . L
Table A.2
Table A.3
Table AA
Table AS
Table A.6
Table A.7
Table A.8
Table A.9
Table A . t O
Propenies of the polypmpylene resins used in determining the
onset of melt fracture ........... .. ............................................................... 76
Properties of the poiypropylene resins used in the foaming
experiments ............................................................................................ 116
Observed number of celis per cross section at various temperatures ...... 151
Independent Variable Settings (Single Constraints) ............................... 173
........................................................................ Sample Prepararion Sheet 174
Response Data ......................................................................................... 175
................... Overall Error S tatistics: Tnnsformed Volume Expansion .. 176
GeneraI Regression Statistics: Tnnsformed Volume Expansion ........... 176
.......... Regression ANOVA Statistics: Transformed Volume Expansion 176
Model Coefficients: Trmsformed VoIume Expansion ......................... 177
Model Terrn Ranking: Transfonned Volume Expansion ........................ 177
........................ Mode[ Term Ranking: Transforrned Volume Expansion 178
Residuals Data: Transformed Volume Expansion .................................. 179
Figure 2.1
Figure 2.2
Figure 2.3
Figure 3.1
Figure 3.2
Figure 3.3.
Figure 3.4
Figure 3.5
Figure 3.6
Figure 4 . L
Figure 4.2
Figure 4.3
LIST OF FIGURES
............................................................... Solubility of Gas into polymers 34
.......................... Effect of the variation ofenergy on the bubble growth 34
Model of a nucleated cell inside a polymer matrix ................................. 35
.......................................... A schematic of the tandem extrusion system 48
........................................................... A schematic of the rnixing section 48
....................................................... A schematic of the gear pump 4 9
......................................................... A schematic of the heat Exchanger 50
............................................................... A schematic of the filment die 50
.............................................................. A schematic of a cooling sleeve 51
Experimental setup for the PVT measurements of
Polypropylene/butane solutions .............................................................. 77
Degassing oven for the PVT measurements ......................................... 78
Caiibration of the Degassing Oven .......................................................... 78
Figure 4.4.(a) Effect of the dissolved butane on the specific volume for
............................................................... linear polypropylene materials 79
Figure 4,4.(b)
Figure 4.4.(c)
Figure 4.5.(a)
Figure 4.5.(b)
Figure 4.5.(c)
Figure 4.6
Figure 4.7
Figure 4.8.(a)
Figure 4.8.(b)
Figure 4.9.(a)
Figure 4.9.(b)
Effect of the processing temperature on the specific volume
........................................................... for linear polypropylene materials 79
Effect of the processing pressure on the specific volume for
linear polypropylene materials ................................................................ 80
Effect of the dissolved butane on the specific volume for
........................................................... branched polypropylene materials 8 1
Effect of the processing temperature on the specific volume
for branched polypropylene materiaIs ..................................... ,.. 8 1
Effect of the processing pressure on the specific volume for
branched polypropylene mitterials ........................................................ 82
Design of cooling system for the high-pressure DSC ceIl ....................... 83
DSC thermograms for linear and branched polypropylene
materials ................................................................................................... 83
Effect of talc on the crystallization behaviors of linear and
branched polypropylene resins .................... ...... .............................. 84
Effect of talc on the degrees of crystdlinity of linear and
branched polypropylene resins ................................................................ 84
Effect of GMS on the crystalIization behaviors of iinear and
branched polypropylene resins ............................................................. 85
Effect of GMS on the degrees of crystdlinity of linear and
branched poiypropylene resins ............................................... ,.,. 85
Figure 4-10 Effect of cooling rate on the crystalIization behaviors of linear
and branched polypropylene resins ......................................................... 86
Figure 4.1 1 Effect of pressure on the crystallization behavior of linear
.......................................................................................... polypropylene 87
Figure 4.12 Effect of pressure on the crystailization behavior of branched
.......................................................................................... polypropylene 87
Figure 4.13 Effect of hydraulic pressure on the crystallization behaviors of
Iinear and branched polypropylene resins ............................................... 88
Figure 4.14 Effect of dissolved & on the crystailization behaviors of
Iinear and branched polypropylene resins ............................................... 88
Figure 4.15 Effect of dissolved CO? on the crysrallization behaviors of
linear and branched polypropylene resins ............................................ 89
Figure 4.16 Experimental setup for rneasuring the onset of melt fracture .................. 90
Figure 4.17 Effect of long-chin branching on the critical shear stress ................... ,,. 91
Figure 4.18. Effect of processing temperature on the critical shear stress ................... 92
Figure 4.19 Effect of processing temperature on the critical shear rate ...................... 92
Figure 4.20 Effect of talc on the cnticai wall shear stress ........................................ 93
Figure 4.21 Effect of GMS on the critical shear stress .......................................... 93
Figure 4.22 Effect of dissolved butane on the critical shear stress ............................. 94
Figure 5.1 The melt strength and melt extensibility of Iinear and
branched polypropylene materials ............................................................... 1 17
............... Figure 5.2.(a) The viscosity of Iinear and branched polypropylene materials 118
Figure 5.2.(b) The capiltary viscosity rneaurements of linear and branched
polypropylene materials .................................................................... 1 18
Figure 5.3.(a) The expansion ratio versus the melt temperature for branched
....................... polypropylene materials usinp butane as btowing agent 1 19
Figure 5.3.(b) The blowing agent efficiency versus the melt temperature for
branched polypropylene materials using butane as blowing
....................................................................................................... went a 1 19
Figure 5 .34~) The cell density versus the rnek temperature for branched
....................... polypropyIene materials using butane as blowing agent 120
Figure 5.3.(d) The cell density versus butane contents for branched
........................................................................ polypropylene materials 120
Figure 5.3.(e) The die pressure versus the melt temperature for branched
....................... polypropylene materials using butane as blowing agent 120
Figure 5.4.(a) The expansion ratio versus the melt temperature for linear
polypropylene materials using butane as btowing agent ................... .... 121
Figure 5.44b) The blowing agent efficiency versus the melt temperature for
linear polypropylene materials using butane as blowing agent ............. 12 1
Figure 5.4.(c) The cell density versus the melt temperature for linear
....................... polypropylene materials using butane as bIowing agent 122
Figure 5.4.(d) The ceIl density versus butane contents for linear
polypropylene materials ..................................................................... 127
Figure 5.4.(e) The die pressure versus the melt temperature for linear
....................... polypropylene materials using butane as blowing agent 122
Figure 5.5.(a) The expansion ratio versus the melt temperature for branched
polypropylene materials using CO2 as blowing agent ........................ 123
Figure 5.5.(b) The blowing agent efficiency versus the melt temperature for
branched poIypropylene materials using COl as blowing agent ........... 123
Figure 5.54~) The cell density versus the melt temperature for bnnched
........................... polypropylene materials using COr as blowing agent 124
Figure S.S.(d) The ceIl density versus COz contents for branched
polypropylene materials ............................. ,., ..................................... 124
Figure S.S.(e) The die pressure versus the rnelt temperature for bnnched
polypropylene materials using COI as bIowing agent ........................... 124
Figure 5.6.(a) The expansion ratio versus the melt temperature for branched
poIypropyIene materials using a blend of butane and COz as
blowing agent ......................................................................................... 125
xvi
Figure 5.6.(b) The blowing agent efficiency versus the melt temperature for
branched polypropylene materiais using a blend of butane and
CO2 as bIowing agent ............................................................................ 125
Figure 5.6.(c) The cell density versus the melt temperature for branched
polypropylene materials using a blend of butane and CO? as
......................................................................................... blowing agent 126
Figure 5.6.(d) The ce11 density versus relative CO? percentage for branched
polypropylene materiaIs using a blend of butane and CO? as
blowing agent .............................................................................. 126
Figure 5.6.(e) The die pressure versus the melt temperature for branched
polypropylene materiais using a blend of butane and CO1 as
............................................................ blowing agent ........................ ,., 126
Figure 5.7.(a) The expansion ratio versus the melt temperature for branched
polypropylene materials using butane as blowing agent for
various taIc contents ...................,.................................................... 117
Figure 5.7.(b) The blowing agent efficiency versus the melt temperature for
branched polypropytene materials using butane as bIowing
agent for various talc contents ............................................................... 127
Figure 5,7.(c) The ceIl density versus the melt temperature for branched
polypropylene matetials using butane as blowing agent for
various talc contents ........................................................................ 128
Figure 5.7.(d) The cell density versus talc contents for branched
polypropylene materials using butane as biowing agent ...................... 128
Figure 5.7.(e)
Figure 5.8
Figure 5.944
Figure 5.9.(b)
Figurc 5.94~)
Figure 5.9.(d)
Figure 5.9.(e)
The die pressure versus the melt temperature for branched
polypropylene materials using butane as blowing agent for
various talc contents .............................................................................. 128
The melt strength and melt extensibility of linear and
........ branched poIypropyIene materials biend ....................................... 139
The expansion ratio versus the melt temperature for linear and
branched polypropylene materials blend using butane as
blowing agent .................................., 130
The blowing agent efficiency versus the melt temperature for
linear and branched polypropylene materials blend using
......................................................................... butane as blowing agent 130
The ce11 density versus the melt temperature for Iinear and
branched polypropyiene materials blend using butane as
......................................................................................... blowing agent 13 1
The ce11 density versus linear polypropylene percentage for
Iinear and branched polypropylene materials blend using
butane as blowing agent ................................................................... 13 L
The die pressure versus the melt tempenture for linear and
branched poIypropylene materials blend using butane as
blowing agent ......................................................................................... 13 1
Figure 5. LO.(a) The expansion ratio versus the melt temperature for re-
extruded branched Plpolypropylene materials using butane as
................................................................ ................. bIowing agent ,.,, 132
xviii
Figure 5.10.(b) The blowing agent eficiency versus the melt temperature for
re-extruded branched P L polypropylene materials using butane
as blowing agent .................................................................................... 132
Figure 5.10.(c) The cell density versus the melt temperature for re-extmded
branched P 1 polypropylene materiah using butane as blowing
agent ................................................................................................... 133
Figure 5.10.(d) The ce11 density versus butane contents for re-extmded
branched P1 polypropylene materials ............................................... 133
Figure 5.10.(e) The die pressure versus the melt temperature for re-extmded
branched PI polypropylene materials using butane as blowing
......................................................... agent ............................ .. 133
Figure 5.1 I.(a) The expansion ratio versus the melt temperature for re-
extruded branched P2 poiypropylene materiais using butane as
blowing agent ..................................................................................... 134
Figure 5.1 I .(b) The blowing agent efficiency versus the melt temperature for
ce-extruded branched P3 ~olypropyIene rnateriaIs using butane
as blowing agent ............................................................................... 134
Figure 5.1 L(c) The ceIl density versus the melt temperature for re-extruded
bmched P2 polypropylene materials using butane as blowing
agent ..................~..~..~...........~..~............................................................... 135
Figure S. I 1 .(d) The cell density versus butane contents for re-extruded
branched P2 polypropyfene materids ................................................... t 35
xix
Figure 5.1 1 .(e) The die pressure versus the melt temperature for re-extruded
branched PZ polypropylene mixeriais using butane as blowing
agent ....................................................................................................... 135
Figure 5. [?,(a) The expansion ratio versus the melt temperature for branched
polypropylene materials using butane as blowing agent for die B ........ 136
Figure 5.12.(b) The blowing agent efficiency versus the melt temperature for
branched polypropylene materids using butane as blowing
agent for die B ..................................................................................... 136
Figure 5.12.(c) The ceIl density versus the melt temperature for branched
polypropylene materials using butane as blowing agent for die B ........ 137
Figure 5.12.(d) The cell density versus butane contents for branched
polypropylene materials for die B .................................................... 137
Figure 5.12.(e) The die pressure versus the melt temperature for branched
polypropylene materials using butane as blowing agent for die B ........ 137
Figure 5.13.(a) The expansion ratio versus the rnelt temperature for branched
polypropylene materials using CO2 as blowing agent for die B ............ 138
Figure 5.13.(b) The blowing agent efficiency versus the melt temperature for
branched polypropylene materiais using CO7 as blowing agent
for die B ................................................................................................. 138
Figure 5. H.(c) The ce11 density versus the melt temperature for branched
poIypropyIene materids using CO2 as blowing agent for die B ............ 139
Figure 5.13.(d) The ceIl density versus COl contents for branched
polypropylene materiais for die B .................................................... 139
Figure 5.13.(e) The die pressure versus the rnek temperature for branched
Figure 6.1
Figure 6.2
Figure 6.3
Figure 6.4
Figure 6.5
Figure 6.6
Figure 6.7
Figure 6.8
Figure 6.9 (a)
Figure 6.9 (b)
Figure 6.10
Figure 6 . I I
polypropylene materiais using CO2 as blowing agent for die B ............ 139
Effect of processing temperature on the extrudate shape ...................... 152
Effect of gris Ioss and crystallization on the volume expansion ............ 153
Fundamental volume expansion mechanism of polypropylene
............................................................... foams ........................... ,.. 154
Images of the foarn extrudate coming out of the die ............................. 155
Effect of processing temperature on the initial expansion rate .............. 156
Effect of processing temperature on the maximum and final
.......................................................................................... diameters 156
Extrudate diameter as a function of the distance frorn the die .............. 157
Description of ceIl shape mode1 ......................................................... 158
.................... The caicuiated volume expansion ratio based on Mode1 1 159
The caiculated volume expansion ratio based on Model 2 .................... 159
Caiculated average ce11 size ..... .............................................................. 160
CalcuIated ce11 waIl thickness ............................................................... 160
xxi
Figure 6.12 The calculated volume expansion ratio based on Mode1 2 .................... 161
Figure 6.13 Calculated average ce11 size From the observed profile ......................... 162
Figure 6.14 Calculated ceIl wall thickness from the observed profile ................... ... 162
matrix representing the complete relationships between FRs and DPs
surface area (cm')
monolayer thickness (ml
gas concentration (ggJgpiF3
solubility of pas in the polymer (cm31g) or (ggJgPly,)
concentration of ;as molecules in solution (#lm3)
capillary diameter (mm)
diffusiviry (cm%)
diffusivity coefficient (cm%)
design parameter
activation energy for diffusion (J)
frequency factor of gas molecuIes joining the nucleus (11s)
funcrionid requirement
linear growth rate
sheet thickness (cm)
Henry's law constant (cm3 [STP]/g- Pa)
sotubility coefficient (cm3 [STPIlg. Pa)
surface nucleation rate
Boltzman constant (.Tm
lamina factor
capillary Iength (mm)
xxiii
LCB long chah btanching (nl1000 carbon atoms)
consistency parameter in power law (Pasn)
mass flow rate (g/min)
Melt flow rate (gllOmin)
Molecular number (Kglmol)
mass uptake at time t (g)
molecular weight (Kglmol)
equilibrium mass uptake after an infinite time (g)
power law index
number of cells per cross section area
number of cells in a defined area ( ex t ) mm'
cell population density (cellslcm3)
nucleation rate (#lm3. s)
gas injection pressure (Pa)
interna1 pressure of the bubble (Pa)
saturation pressure (Pa)
pressure (Pa)
dciving pressure (Pa)
volumetric flow rate (m3/s)
volumetric Bow rate (cm3/s)
spatial coordinate in the cylindrical polar coordinate system (mm)
gas constant (JK)
Reynolds number
xxiv
Y,"
= specific gravity (g/cm3)
= time (s)
= time required for the completion of absorption (s)
= temperature (K)
= crystallization temperature (K)
= glass transition temperature (K)
= fluid velocity (mis)
= fluid velocity components ( d s )
= radial velocity component = k'l r' (mls)
= bubble voIume (m3)
= volume expansion ratio
= work required to generate a bubble of radius r in a body of a liquid (J)
= spatial coordinate in the rectanguiar Cartesian coordinate system (mm)
= ultimate crystallinity for t=-
= relative volumic crystallinity at tirne t
= absolute crystallinity at crystallization time t
= spatial coordinate in the rectmuigular Cartesian coordinate system (mm)
= spatial coordinate in the rectangular Cartesian and cylindrical polar
coordinate systems (mm)
= shear rate ( Ils)
= apparent shear rate (11s)
= surface tension (Pa)
= true shear rate (Ils)
xxv
= shear viscosity (Pas)
= blowing agent efficiency
= actuai volume expansion ratio of the foam
= viscosity of polymer rnatrix (N- slm')
= spatial coordinate in the cylindricd poIar coordinate system
= bulkdensity (g/cm3)
= amorphous region densi ty (kg/m3)
= crystalline densiiy (kg/m3)
= gas density (kglm3)
= polymer density (kp/m3)
= stress tensor components (Pa)
= wallshearstress(Pa)
= specific voiume (cm3/g)
= Gibbs free energy of crystal formation (8
= free energy (1)
= heric of fusion (1)
= molm heat of sorption (J)
= pressure drop (Pa)
xxvi
Introduction
1.1 Preamble
The use of plastic foams in today's technology continues to grow at a rapid pace
throughout the world. The reasons for this growth include the light weight, excellent
strengthlweight ratio, superior insulating abilities, energy absorbing performance, and
comfort features of plastic foams. Foams cm be prepared from any polymer, by introducinp a
gas within the polymer matrix. Selection of polymers for industrial foam applications
depends upon their properties, their ease of manufacture, and the economics of the foaming
system [Il. Polymeric foams comprise a wide variety of materials, with a wide range of densities.
Their main applications include furniture, transponation, bedding, insulation, packaging,
appliances, sports applications, shock and sound absorbers, etc. Foams are usualIy classified
as flexible, semi-flexible, or rigid, and can be fabricated to any desired degree of hardness,
They c m be manufactured by a variety of processes, depending on the application. TypicaI
processing methods include extrusion, injection molding, blow molding, rotations! molding,
themoforming, etc. [ I l . in the years following its discovery, polypropylene went through such a dynamic
industrial developrnent that it is today one of the most widely used polymeric rnaterials, and
still bas a bright future. The world market of polypropylene has grown from around 1.5
million tons in the 1970s, to over 25 million tons in year 2000 [2]. Such trernendous growtfi
is due to the outstanding combination of cost performance, excellent physical propeaies,
smng and continuous expansion of process versatility, and environmental friendly processes
and materials during manufacturing, use, and recycling stages [2].
Due to its outstanding functionai characteristics and low material cost, polypropylene
foams have been considered as a substitute for other thennoplastic foarns in industnai
applications. PolypropyIene materiai is a member of the semicrystalline polyolefin famiIy.
which is resistant to chemicais and abrasion. Polypropylene material has a number of
odvantages over polystyrene and polyethylene [3]: it h a a higher rigidity cornpared to other
polyolefins; it offers higher strength than polyethylene and better impact strength than
polystyrene; and it provides a higher service temperature range and good temperature
stability. However, only limited research has been conducted on the production of
polypropylene foarns because of their weak rnelt strength which rnakes it diffcult to be
foamed compared to the other plastics [4].
1.2 Thermoplastics Foams
Most therrnoplastics foarns are produced by the expansion process, which is based on
the expansion of a gaseous phase dispersed throughout the polyrner melt. The gaseous phase
rnay be generated by separation of a dissolved gas, vaporization of a volatiIe liquid, or
release of gas frorn a chemical reaction. Regardless of the type of blowing agent, the
expansion process comprises three major steps: nucleation, bubble growth, and stabilization.
Nucleation or formation of expandable bubbles can begin within the polymer melt that has
been supersaturated with the blowing agent. Once nucleated, a bubble continues to grow as
the blowing agent diffuses into it. This growth will continue until the bubble stabilizes or
ruptures [5] . Thennoplastic foarns possess a cellular structure. and this interna1 strucnrre
provides unique properties that enable the foarned plastics to be used effectively for various
industrial applications [l]. Foarned plastics can be classified in different ways: by nature as
flexible or rïgid, by dimension as various sections, by density as low and high density, by
structure as open or closed cell, and by ceIl density as fine celled or microcellular+
1.3 Low-Density, Fine-Celled Foam Processing
Low-density, fine-celled foams are chmcterized by a foarn density lower than 0.025
@cm3, a ceII density higher than 10' cells/cm3, and an even ce11 size distribution. The
rationale behind the production of low-density, fine-celled foams is that if bubbles that are
smaller than the existing critical flaws within the plastic can be introduced in sufficient
numbers, then the mechanical properties c m be rnaintained while reducing the material
density [q. Low-density foams with a volume expansion ratio higher than thirty-fold have
been excIusiveIy blown using environmentally hazardous long-chah blowing agents such as
CFCs, HCFCs [Il. However, CFCs and HCFCs are environmentaily hazardous, and because
of the Montreal protocol, these blowing agents will be eliminated from use. A chemicai
blowing agent that generates gales) may also be used for the low-density foam processing
that involves cross-linking [7]. The rationde of this process is to block the escape of blowing
agent through the foam skin in order to promote large expansion by cross-linking of the
polymer in the skin area. Furthemore, the cross-linking of the polymer melt in the core area
also helps increase the volume expansion because the cross-linked polyrner in the core area
retards the diffusion rate of blowing agent within the polymer. However, excessive cross-
Iinking of the polyrner melt is not desirable because the melt may be too stiff to allow ceIl
expansion. Although this method can be used to effectively produce largely expanded foams,
this process is considered reIatively expensive because of the additional cross-linking step,
which is required before foaming. AIso, cross-linked foarns are not recyclable.
Gaseous blowing agents such as CO2 and N2 have been used in plastic foam
processing to manufacture relatively high-density foarns with volume expansion ratios in the
range of 1.2 to 30 fold (mainly Iess than IO fold) [4,8-191. Many patents describe well-known
methods of inert gas-based foam processing for the manufacture of relatively high-density
foams [IO-Id]. Since the inert gas blowing agents have higher volatility, and therefore higher
diffusivity than the long-chah blowing agents, gaseous blowing agents can escape erisily
during expansion [1,20,21]. Therefore, it is very dificult to obtain a low-density foam with ü
large expansion ratio of over thirty-fold with an inert gas. On the other hand, the long-chain
blowing agents have Iow diffusivity because of their low volatility. This low diffusivity
offers a tremendous advantage in controlling cell growth to achieve a very high expansion
ratio of over thirty-fold since it is easier to control ceII growth because of the slow growth
rate and because gas escape can be prevented. Although the gaseous blowing agents could
not be used for low-density foam processing because of the technological difficulties. they
have been effectively used for producing fine-celled or microcelIuiar foams because of its
high volatility [IS-211.
1.4 Objective of the Thesis
The main objective of this research is the improvement of understanding of an
extrusion foaming process for the manufacture of towdensity, fine-celled propylene foams:
This research entails 1) the development of effective strategies for the production of low-
density, fine-celled polypropylene foams; 2) the identification of the fundamental
mechanisms governing the volume expansion behavior of polypropylene foam; 3) the
measurement of the Fundamental properties of polypropylenelgas solutions such as the PVT
data, thermal properties, and critical shear stress.
Despite the advantages of polypropylene materials over polystyrene and
polyethylene, most commercially availrtble foams have been made out of polystyrene and
polyethylene because of the technological difficulties in foam processing of polypropylene
materiais. In this context, the main purpose of this research is to develop effective strritegies
that can be used for the production of Iow-density, fine-celled polypropylene foams. [t is
aimed that the foam density be lower than 0.0125 &m3 (or equivalently the expansion ratio
bigger than fony-fold) and the cell-population density be higher than 1o6 cells/cm~n
response to a request from industry [22], Also, the hndamental mechanisms that govern the
volume expansion ratio of polypropylene foaming needs to be identified in order to
determine the critical processing and materials parameters for controlling the qualities of
polypropylene foams.
Fundamental studies will also be conducted in order to determine the mechanisms
governing the volume expansion ratio of polypropylene foams. When a gas dissolves in a
molten polymer, the polymer sweils due to gas absorption. By meüsuring the extent of
swelling of the poIymer due to a given pressure, tempenture, and concentration of dissolved
gas, the pressure-specific volume-ternpenture (PVT) data of polymerlgas solutions c m be
acquired. The knowledge of the values of the thennodynamic propenies of polymerlgas
solutions is of a crucial importance for the successful processing of plastic foams because
properties such as the surface tension of a polymerlgas solution [23], the solubility and
diffusivity of the gas in the melt [24], and the shear and extensional viscosities of
polymerlgas solutions [25-271, are strong functions on the PVT data of the respective
polymerlgas solution. Consequently, acquiring the PVT data and their reiationships for a
particular polymerlgas solution is essentiai to the understanding of its foaming behaviour and
improving current foaming technologies.
The thermal behavior of a given polyrner-additive system is an important parmeter
for plastic foam processing because the bubble growth phenomenon c m be governed by
crystailization: volume expansion of foam can be completed at the moment of crystalIization
even if the gris saturated in the polyrner matrix has not completely diffused out to the cells
yet. Therefore, understanding of the crystaiIization behavior is critical in plastic foam
processing of semi-crystalline material. The branched structure of the resin used to prevent
ceIl coaiescence will affect the crystallization temperature. Also, the additives used to
enhance the nucleability and surface quality will affect the crystallization temperature.
Finally, the presence of the blowing agent in the polymer matrix will also affect
crystallization, In this context, it is proposed that the crystallization behaviors of
polypropylene materials be investigated at various processing and materials parameters.
One of the major factors that affect the volume expansion behavior of polypropylene
foams is extmdate distortion or meit fracture. Therefore, in order to get high volume
expansion for the polypropylene foams, it is important to avoid the onset of melt fracture. In
the case of foaming extrusion, the skin of extnided foam is stretched as the expansion occurs.
As a consequence, the foam skin becornes shiny as long as the expanded foarn does not
contract due to gas loss. Even if me1t fracture occurs and a sharkskin is generated on the
extmdate surface, the foam expansion of the extrudate causes the foam skin to be stretched,
and thereby, the trace of the sharkskin cm be easily removed. Therefore, it is very difficult to
detect the onset of melt fracture by simply observing the surface qudity of the fully
expanded foam. Due to the difficulties of determining the onset of melt fracture for polyrner
foam, this research will present a system developed for merisuring the critical wrill shear
stress by monitoring the early stage of the extnidate at the die exit before it expands.
1.5 Overview of the Thesis
Chapter 2 presents a literriture survey on the foaming of polypropylene materiais. It
inchdes the low-density foam production, the fine-ce11 production and the manufacture of
new foamable-grade high-melt-strength potypropylene materials. Since one of the objectives
of this thesis is to identify the fundamental mechanisms governing the volume expansion of
polypropylene foams, the theoretid aspects for ce11 growth, cell nucteation and
crystailization kinetics are also included in this chapter,
Chapter 3 describes the design of the extrusion system in the conceptual stage. Since
an axiomatic design rnethod provides a scientific and systematic bais for andyzing existing
systerns and designing new systems, it is employed to analyze the tandem extrusion systern.
The details of this analysis and design are described in this chapter. The detailed design of
different components in the tandem extrusion system is also given. These include the
extruders, a difision enhancing device, a heat exchanger, die, and other components,
Chapter 4 describes the fundamentai measurements conducted on the polypropylene
materials with foaming agents. The pressure-volume-temperature (PVT) relationships for
polypropylene/butane solution are identified- The effects of materials branching, processing
temperature and pressure, and the gas content on the specific volume measurements of
polypropylene materials are determined. Other measurements include the crystallization and
the onset of melt fracture measurements. For crystallization measurements, the effects of
materials branching, foaming additives, and dissolved gas on the thermal behaviors of
polypropylene materials were identiiied. Finally, with respect to the onset of rnelt fracture
measurements, the effects of materials branching, foaming additives, and dissolved gas on
the rnelt fracture behaviors of polypropylene materials were investigated.
Chapter 5 describes the production of low-density, fine-celled polypropylene foams,
and experimental results are presented to verify the feasibiiity of the proposed ideas. It
presents an effective strategy for the control of ceil growth tu achieve a desired volume
expansion ratio. The effects of processing parameters such as the temperature, the materials
parameters such as the blowing agent, nucleating agent, and long chain branching of
polymer, and the die geometn'es such as the diameter and length of tube on the cell
morphology of polypropylene foams were investigated.
In Chapter 6, the fundamental mechanisms goveming the volume expansion ratio of
polypropylene foarn filaments were identified. Firstly, the cell growth phenornena were
described based on our experimental observation. Secondly, the system setup used for
capturing the image of foam expansion was described, ThirdIy, the procedure for monitoring
the foam expansion mechanisms and the image andysis were elucidated. FinalIy, the effects
of processing and materials parameters on the foam expansion mechanisms were depicted
based on the extrudate images.
Chapter 7 provides a summary and the concIusions of the research. Suggestions for
future work are presented based on the research results in Chapter 8.
Chapter 2
Literature Review and Theoretical Background
2.1 Background on Polypropylene Foam Processing
There are very few studies chat have investigated the production of polypropylene
foams. Polypropylenes possess weak rneit strength, which increases the difficulties of
foaming when compared to other plastics [4j. When the melt strength is too weak, the celI
walls separating the bubbles may not have enough strength to bear the extensional force and
may rupture very easily during foaming. As a result, the foamed polypropylene products
usually have a high open-ce11 content and thus are unsatisfactory for many applications.
2.1.1 Cross-Linked Polypropylene Materials
In order to improve the foamability and therrnoformability, material modifications or
new resin developments have been conduceed for polypropylene materials. Earlier efforts
inciuded cross-linking of polypropylene resins, which significantly improved volume
expandability, cell uniformity, and thermoformability of the foarns [28-3 11. Nojiri et aI. 1281
produced polypropylene foams by cross-linking of polypropylene with trïacylate or
trïmethacrylate. The cross-linking agent and the polypropylene polymer were mixed in an
extruder. Then, the extruded sheet was irradiated with y-radiation to introduce cross-linking
in the polypropylene. The radiated sheer was finally foamed on a wire net housed in a
constant temperature bath. They claimed that the foam produced by their technique had a
good therrnoformability, a high eiongation, a uniform ceil structure and a dcnsity of about
0.035 &m3. In another scudy by Nojiri et al. (291, they claimed that they could pmduce
polypropylene foams with a volume expansion ratio higher than 20-fold by chernicd cross-
Iinking. They developed a chemicd cross-linking process using silane compounds and hot
wacer treatment. SimiIarly, in a study by Lee and Wang [30], polypropylene structurai foams
were prepared by cross-linking the polypropylene with vinyi trimethoxy silane. However,
they claimed that a suitable amount of the cross-linking agent has to be used in order to
produce good foams. When the polypropylene was not cross-linked enough, the pressure
produced inside the nucleated bubbIe was much higher than the ce11 waIl could withstand,
and consequently, more open cells were formed. The smallest average cell size found in their
study was 520 pm. Kitagawa et al. [3 11 descnied a method of producing a crosslinked,
uniformly foarned article of a polypropylene resin. They claimed that the resulting foam
article had better heat resistance than the existing ones. The polypropylene resin was
prepared by mixing it with the blowing agent at low temperature. An ethylene-propylene
random copolymer resin which contained 2 to 10 wt% of ethylene was used as the
polypropylene resin. Feichtinger [32] disclosed a new cross-linked foamable composition of
linear polyolefins blended with polypropylene. He claimed that the new composition consists
of new cross-linked polyolefin copolymer that shows improvements in strength, toughness,
flexibility, heat resistance and sealing temperature as compared to conventional low-density
polyethylene compositions. He used both homo and copolymer polypropylene resins and the
obtained densit), was 19 to 160 kg/m3.
Although cross-Iinking promoted high expansion and better ce11 morphology, it is
difficult to process and is hazardous in nature. Therefore its use for achieving low-density,
fine-celled foam is Iirnited.
2.1.2 High-Melt-Strength Polypropylene Materials
A special grade of polypropylene called high-melt-strength (HMS) polypropylene has
been developed and has been found to be satisfactocy to manufacture foamed prridiicts.
Efforts have dso been made CO promote long-chain branching [33-371 to improve the melt
strength of polypropylene materials. Bradely and Philips [33] ernployed CFC I l4 injected
into conventional polypropylene and HMS polypropylene melt streams using a commercial-
scde single screw tandem system with a filament die. Calster [34] presented a newly
developed grade of polypropylene specifically for foam, sheet, pipe and profile extmsion,
and a broad range of densities was achieved, from 600 to 35 kg/m3. or equivdentiy 1.5 to 25
fold expansion ratios- Bavaro [35] has presented the production of low-density HMS
polypropylene foarn sheets and planks down to 32 kg/m3, or equivdently 28 fold expansion
ratio. He used HMS polypropyIene resins blended with an endothermic blowing agent and
gyceroi monostearate aging modifier, and this pellet blend was plasticated while injectinp
liquid isobutane blowing agent directly into the melt. The resultant residblowing agent
solution was conditioned and extruded while maintaining a solution temperature of 138 OC.
Panzer et al. [36] claimed that adding HMS polypropylene to Iinear standard one
improved the processability of polypropylene materials in extmsion as weIl as in other
technologies such as blown film with air cooling, extrusion coating, thermoforming, blow
motding and fiber spinning. Ratzesh et al. [37] developed a new grade of H M S propylene
polpers, They ciaimed that the new material is a combination of high melt strength with
high meIt extensibility, which is due to the introduction of long-chain branches into the
structure of polypropylene polymer. In this way, the polypropylene processing can be
improved so that the low-density polypropylene foams can be manufactured by extrusion
processes known from polyethylene and polystyrene.
Vanous other grades of HMS propylene polymers [38-391 have also been presented.
Thoen et al. [38] presented new HMS polypropylene resins for extrusion foams. They
claimed that with the materials better control of cell growth could be achieved due to high
melt strength and drawability, resulting in a closed cell structure. The resulting foarn had a
density in the range of 25 to 700 kg/m3. Park [39] introduced ii reactor made HMS
polypropylene employing a special catalyst system with unique polymerization process
technology. He ciaimed that this material is designed for foarning applications with improved
melt strength and melt elasticity while retaining ail physical, mechanical, optical and
chemical resistance properties of the conventional polypropylene.
In general, compared to the conventional linear polypropylene, this branched grade
has a high enough melt strength and melt elasticity that solves the bubble stability. Therefore,
this HMS -de enables wide usage of polypropylene materials as it overcornes the weak
melt strength for linear materials.
2.1.3 Low-Density, Fine-Celled Polypropylene Foams
Various methods for the production of polypropylene foams have been reported. in
addition to these rnethods, several authors have investigated the foaming behaviors, or more
specifically the ce11 nucleation and volume expansion behaviors, and the final ceII density
and foarn density of polypropylene foams using various nucleating and blowing agents. A
brief review of these techniques and results are presented below.
Holland and FeIlers [40] described a method to manufacture poIyethylene or
polypropylene foam, a quantity of powdered po1yrner consisting of polypropylene or
polyethylene powders was mixed with a blowing agent consisting of minerai oil, to a
substantiaily uniform biend and piaced in a closed mold. Then, the mixture was heated to a
temperature at which the polyrner was in a substantial Iiquid state, such that the intimately
mixed polymer and mineral oil formed a solution. Upon further heating, the mineral oil was
caused to undergo a phase change from a liquid to a vapor at its boiling temperature, then the
minera1 oil acted to expand the poIymer into a foam.
Fukushima et al. [4l] described a process for producing a highly expanded
polypropylene foam, which had a uniform viscosity during extrusion foming, improved melt
strength and free from surface roughness by using a polypropylene resin having a specified
meIt tension. They used in their process a mixture of trichlorotifluoroethane and
dichIorotetrafluoroethane as a volatile foaming agent.
The use of HMS polypropylene seems to be very effective in producing high qudity
foams that do not suffer from the bubble stability problem. Much more improved foamability
and themoforrnability were observed from long-chain branched HMS polypropylene resins.
Park et al [42] produced a multilayered foam sheet usefut for forming rigid or sernirigid
articles for packaging applications, comprising at least one layer of a polypropylene foam
sheet. They chimed that this thermoformable rigid or semi rigid polypropylene foüm sheet
having a srnooth surface and uniform ce11 structure and a density of at Ieast 40 K ~ / I I I ~ is
prepared by extruding a mixture of nuckating agent, a physical btowing agent and a
propylene resin having a high melt and high melt ehsticity. In another study by Park et al.
[43], they cIaimed a continuous process for producing a foam sheet of polypropylene with
the same characteristics by extruding a foaming mixture, of nucleating agent, a physical
b1owing agent and a propyIene resin having a high melt and high melt elasticity, inco a
foarned extrudate and forming the foam extrudate into a continuous foam sheet.
Park 1441 described the production of a low-density extnrded propylene pdyrner
forun comprising primarily of an expanded propylene polyrner materid wherein the foam had
a blowing agent comprising at Ieast 15 percent by weight of one or more inorganic biowing
agents. The propylene polyrner material comprised at lest about 15 weight percent of
propylene monomeric units. The foam had a density ranging from 10 to about 150 k@m3, and
an average cell wail thickness of less than about 35 microrneters. In another work by Park
[45], he described the developrnent of a closed-ce11 polypropylene foam, with at least eighty
percent closed cells, a foamability characteristic of less than about 1.8 and a foarn density of
less than about 5 lblf?. Also, he disclosed several rnethods for making such closed-ce11
polypropylene foarns. Also, Park [46] disclosed a process for rnaking a low-density, open-
celled, extruded propylene foam prirnarily comprising of an expanded propylene potymer
material wherein the foarn had a blowing agent comprising greater than 85 wt% of one or
more organic blowing agents based upon the total weight of the blowing agent. Wilkes et al.
[47] disclosed a process for producing extruded low-density propylene polymer foarns. They
used a propylene resin that had greater than 50 wt% of branched rnolecular structure, aIso
they used propylene resin that had less than 50 wt% of a propylene-ethylene copolymer.
They added a polymeric modifier, nucleating agent, and they injected an organic, non-
organic, or a blend of blowing agents. They clairned that they obtained closed cell propylene
polymer foam with a density between 10 kg/rn3 and 150 kg/rn3, and with a thickness in
excess of 1.3 cm.
Andreassen et al. [48] studied the production of polypropylene foarns using CO-
rotating twin-screw extrusion. with high melt strength and conventional linear polypropylene.
They concluded that both branching and high rnolecular weight fractions contributed to high
rnelt suength, and that only branching resulted in a pronounced strain hardening effect. They
chimed that the conventional linear polypropylene could be foarned to produce small and
rnainly closed cells. They aiso observed extrudate distonion with linear polypropylene grade
with a broad moIecular weight distribution; in addition they obtained very thin ceIl walls
using this grade of linear polypropylene.
Alteepping and Nebe [49] were able to extrude a polypropylene foam product hriving
a density Iess than 0.2 g/crn3 by foarning a composition cornprising two different viscosity
polypropylene components: a major proportion of a low viscosity polypropylene cornponent
hwing a melt viscosity of less than 2x10~ poise and a rninor pmponion of a high viscosity
polypropylene cornponent having a melt viscosity of greater than 2 . 5 ~ 1 0 ~ poise (at 190" C
and a shear rate of LOO0 sec-'). They clairned that this composition of polypropylene resins
would enhance the f o m stability significantly. They aiso believed that the use of a
nucleating agent such as hydrocerol would produce a finer ce11 structure. With a suitablc
selection of the blowing agent and the nucleating agent, they obtained a ce11 structure having
a cell size of 300 p.
Ahmadi and Hornsby [50,51] investigated the effect of the processing parameters on
the structures of the injection molded polypropylene foams in great detail. It was observed
that the polypropylene foam processed by a high injection rate of polymer melt into the mold
had a relatively uniform and fine cell structure. By contrast, the foam produced by a slow
injection rate exhibited large cells and non-uniforrnity in the cellular structure. Also, they
experienced a high melt pressure and a low melt temperature favored in the reduction of ceIl
size. The smallest average ce11 diameter found in their published data was approirimately 120
p. They concluded that the uniformity and fineness of the cellular core was favored by the
use of high injection rates combined with a low melt temperature and a high melt pressure,
On the other hand, Colton [52] was able to produce a very fine cellular structure in the
copolymeric and nucleated polypropylene, with a ceIl size of 5 pn, utilizing a microcelIuIar
batch processing technique. He claimed that the addition of sodium benzoate would also
enhance the bubble nucleation. However, the batch technique is not applicable to the
continuous extrusion process. Dey et al. [53] have succeeded in foaming polypropylene
homopolymer with carbon dioxide to foam densities between 500 and 300 kg/m3.
equivalently 1.8 to 3 fold expansion ratios, using a single screw extruder and a standard
ribbon die.
Park and Cheung [4] studied the ceil nucleation of linear and branched polypropylene
resins using CO1 and isopentane as blowing agents in extrusion. They found out that the cell
density increased with an increased amount of blowing agent for both CO2 and isopentane.
They dso observed that the phenornenon of ceIl coaiescence in Iinear polypropylene was
more than in branched polypropylene. In another study of Park et al. [54] they studied the
effect of tatc on cell nucleation in extrusion foam processing of polypropylene with CO1 and
isopentane. They concluded chat when isopentane was used, the ceIl density was dominated
by the talc concentration, and the isopentane content did not greatly affect the cell density.
By contrast, when COz was used the maximum cell density was not improved by the talc
concentration.
The above studies and patents show successful attempts for achieving either Iow-
density or fine-ceIled polypropyIene foarns. However, the strategies for achieving low-
density and fine-celled polypropylene foms were not identified. Also, there was no clear
understanding of the fundamental mechanisms governing the expansion behavior of
polypropylene foams.
2.2 Background on the Properties of Plastic Melts and PlastidGas Solutions
When gas dissolves in the polymer, the thermodynarnic, thermal, and rheological
properties are changed. Since the PVT data, cryscallization temperature, and critical shear
stress are important properties in extrusion foarning, the effect of gas on these properties tire
studied.
2.2.1 PVT Properties
To determine the effects of the gas content in polymer processing, one must
accurately estimate the composition and specific volume of the polymerlgas solution over a
wide range of pressures and temperatures, Thus far, the availability of such valuable
information has been limited to only a few specific materials in a specific range of
temperature and pressure. However, various techniques for measuring the PVT properties of
a pure polymer have been reported by severai authors [55-581. In addition, some attempts
have been made to measure the PVT relationship of a polyrner/gas sample but only in a solid
state [59], while other studies [60,61] presented a method for measuring the polymerlgas
solution in a molten state.
Measurements of the PVT properties of pure polymer systems have been well
established by several authors. Foster et ai. [ 5q used a cylinder-piston type dilatometer in a
compressible chamber to acquire the PVT data. A bellows-type dilatometet was used by
Zoller [56, 571 and Sato et al. [58]. The advmtage of the bellows type dilatometer is that it
can measure both solid and moiten states of the polymer in a wide range of temperature and
pressure, however, it cannot be used CO measure the P W properties of polymerlgas solutions.
An attempt to measure the PVT relationships of a polyrnerlgas sample was made by
Hirose et al. [59]- They determined gas absorption by measuring the change in the length of
a thin polymer sample in its solid state. The pressure range for this apparatus was up to 50
atmospheres anb the temperature was in the range of 25OC to 55'C. The major problem
associated with this type of dilatometer was the distortion of the sample (Le., warping),
which affected the accuracy of the measurements. In addition, the PVT data may be
influenced by the anisotropy of the molecuiar chains in the sample, because the polyrner
chains may be oriented due to the small thickness of the extruded samples. Moreover, the
operating pressure and temperature ranges were too low to detennine the PVT behaviors in
the molten state of polymer/C02 solutions.
Park et al, [60,6 11 presented a novel technique to directly measure the PVT properties
of polymer/C02 solutions at high temperatures. This rnethod employs on-line mesurement
of the volume flow rate with the aid of a positive-displacement gear pump. In order to veriw
the function, caiibration experiments were conducted on pure polymers such as
polypropylene and polystyrene without implementing a gas injection system and the results
were compared to the known data [60]. The system was modified to measure the PVT
properties of polymer/C02 solutions by attaching a CO? injection pump [6 11.
In general, most studies have investigated the effect of processing conditions and
blowing agents on the PVT measurements of polymer and polymerlgas solutions. There was
no reported data in the literature for the polypropylene/butane system, and no theory
described the effect of polyrner swell on the foarning behavior.
2.2.2 Crystallization Behavior
A number of studies have been attempted to investigate the crystallization behavior of
polypropylene. The effects of talc and heterogeneous nucleating agents including calcium
carbonate on the crystallization of polypropylene have been studied 162-671. Dey et ai.
studied the effect of physical blowing agents on the crystallization temperature of polymer
melts by rneasuring the thermal conductivity in a specially designed high-pressure reactor
[68]. A variety of techniques have been attempted to investigate the thermal behavior of
polyrner and compressed gas system 1691. A DSC was empIoyed to scan a polymer sample,
which was pre-satumted with gas dthough the loss of gas during the sample handling and
scanning was unavoidable [70]. A sealed high-pressure DSC pan was also used to scan a
polymer-gris system [71]. Handa et al. developed a high-pressure calorimetric technique to
investigate the glass transition depression characteristics of PMMA with compressed gas
1721. They aiso studied the effect of supercritical CO7 on the polymorphism in syndiotactic
polystyrene (73,741 and the ptasticization effects of poIymers with high-pressure gases
[69,72].
He and ZoIIer studied the crystallization kinetics of polypropylene, polyamide, and
poly(ethy1ene terephthalate) using a pressure dilatometer to follow the volume changes
associated with the crystallization process [75]+ Zhang and Handa reported the CO2-assisted
melting of semicrystalhe polyrners such as polystyrene and poly(ethy1ene terephthalate)
using a high-pressure DSC 1761. Crystallization of nmorphous polyrners induced by
supercritical COz ha been investigated and compared with thermal crystallization
[70,77,75], Amorphous poIy(ethylene terephthdate), poly(viny1idene fluoride)/poly(methyI
methacrylate) blends, and bisphenol A polycarbonate have been used for those studies. The
plasticization of CO2 c m facilitate crystalIization in certain polyrners to an extent comparable
to that achieved using an organic liquid or vapor.
Several authors have investigated the crystaflization kinetics of various materials.
PateI and SpruielI [793 presented, in the context of process modeling, a review of some of
these rnodels. Some of the earIiest pubIished results on the isothennaI kinetics of
polypropylene crystallization were reported by Keith et al. [80], Griffith and Ranby [81],
Marker et al. (821 and others [83-851. In light of this previous work, Goldfrirb [86] undertook
a study to establish the primary nucIeation rates for isotactic poiypropylene using an optical
polarizing microscope in combination with a heating stage. Lin [87] investigated the rate of
crystallization of Poly(ethy1ene terephthalate, PET) using DSC and reported a nonconstant
Avrami exponent, n, for PET indicating the existence of secondary crystaIlization. With the
aid of DSC, Chiu et al [88] analyzed the isothermal crystallizarion kinetics of isotactic
polypropylene using the Avrarni equation, the Turnbull-Fisher nucleation theory and the
HofFman-Lauritzen growth theory for modeling- They examined the temperature dependence
of the nucleation rate, linear growth rate, and Avnmi constant, and Cound maximum values
in their relations. They aiso reported the Avrami exponent of isotactic polypropylene to be 3
indicating instantaneous nucleation followed by sphemlitic growth. Lim and Lloyd [89]
employing both the Avrarni and Lauritzen and Hoffman analyses investigated the isothermal
crystallization behaviour of isotactic polypropylene. A mater curve approach was eqioyed
by Chan et al, [90j in an attempt to mode1 isothermai crystallization behaviour of PET.
Supaphoi, Phillips and Spruiell [91] detemined the bulk crystallization kinetics of high
density polyethylene under very high cooling rates. Verhoyen et al. 1921 developed a macro-
kinetic mode1 for the crystallization behaviour of semicrystalline polyrneric materials, which
takes into account the induction time, final degree of crystallinity and secondary
crystallization. Finally Supaphol et al. [93] investigated the isothemal crystallization of
syndiûîactic polypropylene.
The above studies have investigated the crystallization kinetics for various materials.
The effect of foarning additives such as nucleating and blowing agents on the crystallization
behavior were elucidated. However, there was no mention of the effect of crystallization
kinetics on the foaming behavior.
2.2.3 Melt fracture
A number of studies have been attempted to investigate the melt fricture behavior of
polypropylene. Sammler et al. [94] investigated the extnidate distortions of polypropylene
using two types of isotactic polypropylene resins of identical melt flow rate. They observed
"spurt" melt fracture behaviors in the case of high melt strength polypropylene. They claimed
that the long-chah branching is the most probable molecular origin of "spurt" melt fracture
of these polypropylene resins. Baik et al. [95] studied the melt fracture characteristics of
controlled-rheology polypropylene, using a capillary rheometer. They investigated the effects
of shex rate and die geometry on the melt fracture behavior. They observed that there was no
surface mek fracture and extnidate distonions decreased when the shear rate was increased
for a given material and processing conditions. Also, they found that the seventy of melt
fracture reduced with increasing the UD ratio.
Kazatchokov et al. [96] investigated melt fracture behavior of molten polypropylene,
using a capillary rheometer. They observed a sudden transition from a smooth extrudate to a
highly distorted one in the case of polypropylene unlike the case of polyethylene, They
clairned that the onset of gross melt fracture was detected to occur at critical shear stresses in
the range of 0.14-0.15 MPa. Also, they noted that both the critical apparent shear rate and
critical shear stress decreased with increasing the UD ratio.
Melt fracture has been studied with other materiais as weli. Hatzikirakos et ai. [97J
studied the effect of layering the polyrner wall interface with ~luorocacbon elastomer on wall
siip and melt fracture of high density polyethylene. They clru'med that the slip occurred when
the critical shear stress was above 0.09 MPa, and chat the presence of fluoro-elastomer at the
interface reduced the melt fracture. They aIso found that the onset of melt fracture was a
function of the temperature, the diameter of die, and the LID ratio of the die. In another
study, Hatzikirakos [98] investigated the onset of wdl slip and sharkskin melt fracture in a
capillary flow. He observed that the onset of sharkskin melt fracture occurred at a critical
value of the shear stress, which was higher than that of the onset of wall slip. He clairned that
the severity of melt fracture for short capillaries was due to higher extensional rate, wall
shear stress and slip velocity at the exit of short die. He also observed that an increase in
temperature delayed the onset of surface melt fracture at higher apparent shear rates, Baird et
al. [99] investigated the surface melt fracture behavior of linear low-density polyethyiene.
They conducted capillary die experiments with different UD ratios and slit die experiments
with various regions coated with fiuoro-elastomers. They obsewed a decrease in the severity
of melt fracture with incrcasing UD ratio in capilIary experiments. They also clairned that
coating the re-entrant corner and exit region eliminaced sharkskin over the achievable range
of rates in the slit die.
Jeyaseelan et al. [100] investigated the polymer melt instability in a large amplitude
oscillatory shear flow between paraIlel plates. They observed that for polyethylene, the
critical strain amplitude for instability in large amplitude oscillatory shear increased with the
critical shear rate for the onset of melt fracture in capillary flow. They also studied the effect
of using three different coatings on the slip envelope for large amplitude oscillatory shear,
and they clairned that one additive significanciy decreased the instability envelope, while the
other two increased it.
Ramamurthy [101j studied the possible criterion for the occurrence of melt fracture
for different polymers such as polyethylene and polypropylene, using a capillary rheorneter.
He observed that, in the case of Iinear low-density polyethylene, the extrudates were srnooth
and glossy below a certain critical apparent shear stress (0.14 MPa), and that the surface
roughness, which is the onset of melt fracture, occurred around the critical stress value. He
also observed that at a stress value of 0.43 MPa, the extnidates were grossly distorted. He
clairned that for lineac low-density polyethylene the critical shear stress (0.1 1 MPa) is
independent of molecular weight and melt temperature, but is slightly influenced by the
material of construction of the die.
Rosenbaum et al. [IO21 investigated the influence of using fine panicles of boron
nitride on the processability of polyolefins and fluoropolymers, using a capillary rheometer
and an extrusion system. They claimed that the boron nitride was found to eliminate
sharkskin melt fracture and postpone gross melt fracture, and that the degree of the effect was
dependent on the resin type, additive concentration, and temperature. In another paper [103],
Rosenbaum et al. studied the melt fracture behavior of two teflon fluoropolymer resins, by
canying out capillary experirnents- They observed that the onset of surface melt fracture
appeared at a criticai shear stress above O. 18 MPa. They also claimed that at higher apparent
shear rates, oscillating melt fracture occurred due to the wall slip and compressibility of the
melt.
Oyanagi [104] investigated the melt fracture and sharkskin behavior for high-density
polyethylene. He observed that a progressive increase in shear rate led to a transition from
steady state flow to spiraling and then to melt fracture. He claimed that this behavior for
high-density polyethylene was not the same as in the case of polystyrene and low-density
polyethylene. He also claimed that the pitch of spiraling and irregularities increased with an
increase in sample molecuhr weight. He concluded that there are different values of critical
shear stresses corresponding to spiraling and melt fracture.
The effect of processing conditions and foaming additives on the melt fracture
behavior of polypropylene and other materials were thoroughly investigated. There was no
Iiterature available on the effect of bIowing agents on the melt fracture behavior. Also, there
was no clear understanding of the effect of meIt fracture on the volume expansion behavior.
2.3 Theoretical Background
The foaming process in generai, comprises of three major steps: solution formation of
polymer and gas, nucleation, and ceil growth. Accordingly, a theoretical background for each
of the major processing step is prestnted in this section foliowed by the elucidation of the
crystallization kinetics and the ce11 growth mechanism.
2.3.1 Solution Formation
The key issue in continuous foarn processing is the formation of homogeneous
polyrnedgas solution since its qudity significantly determines the number of bubbles
nucleated. The correct arnount of blowing agent should be injected into the barre1 so that it is
completely mixed and dissolved into the polymer. Large voids will form if an excess amount
of blowing agent is used, which cannot be dissolved into the polyrner. Therefore, it is crucial
to determine the solubility, Le., the amount of blowing agent that can be absorbed and
dissolved into the polyrner for different processing conditions. An estimation of solubility is
critical for production of low-density, fine-celled foarn since the presence of large voids
cannot be tolerated.
Solubiliîy
The solubility is conventionally determined in a batch process over a limited
temperature range. In the batch process, a sheet of polymer sample is saturated with the
blowing agent by placing it in a high-pressure chamber connected to the blowing agent
reservoir. At this high pressure, the blowing agent will continuously diffuse into the polyrner
matrix until the solubility limit is reached (Figure 3.1). Although the solubility limit
theoretically occurs at time infinity, the instantaneous concentration of the blowing agent in
the polymer still can be obtained using the following equation [los]:
where D = diffusivity (cm%),
Mt = mass uptake at tirne t (g),
h =sheet thickness (cm),
M, = equilibrium mass uptake afcer an infinite time (g),
t = elapsed time (s).
The amount of mass uptake eventuaily tends to Ievel off at M, in the absorption process, and
the solubility Iimit can be calculated by dividing the mass uptake (M,) by the mass of the
polyrner sarnple.
On the other hand, the sotubility Iirnit of gas dissolved into the polyrner depends on
the system pressure and temperature and cm be estimated by Henry's law [106]:
where c, = solubility af gas in the polymer (cm31g) or (g&gpolymerji
H = Henry's Iaw constant (cm3 [STP]/g-Pa),
p, = saturation pressure (Pa).
The constant H is a function of temperature desciibed by:
where R = gas constant (JtK),
T = temperature (K),
4, = solubility coefficient constant [cm3 [STP]/g-Pa),
4iix = molar heat of sorption (I).
The molar heat of sorption, AH,, can be a negarive or a positive d u e depending on the
pol ymer-gas system.
Equations (2.1) and (2.2) allow us to determine the solubility of a blowing agent in
the polymer at the pracessin; pressure and temperature, and estimation of the soiubility of
CO2 in some polymers has been given in the literature [17,18]. However the solubility of
butane in polyrners was not available in the literature. In the acrual ~xtrusion foarn
processing, the ratio of gas [O polymer weight is maintained below the solubility Iimit, by
controlling the flow rate of the polymer and gas with injection amount.
The initia1 dope of the curve in Figure 2.1 corresponds to the diffusivity of the
blowing agent into the polymer matrix, and can be used to calculate the diffusivity using the
following equation 11051:
where is the value of (tfl?) at MJM- =1/2. By rearranging Equation (7.4). the time
required for the compietion of absorption cm be approximated from Equation (2.5) as given
below [107]:
Equation (2.5) indicates that the time of absorption is inversely proponional to the diffusivity
(D) and proportional to the square of the diffusion distance (ha).
The diffusivity D is mainly a function of temperature, and its influence cm be
explained by the following equation [105,106]:
where Do= diffusivity coefficient constant (cm%),
EF activation energy for diffusion (J).
Dissolution
During the batch process, the gas absorption process into the polyrner rnatrix will
automaticatly terminate when the concentration of the bIowing agent dissolved into the
polyrner matrix and on its surface has reached an equilibrium value. Thus, large voids are not
generated in the batch proccss since it is impossible for any extra gas to be absorbed by the
polymer above its solubility limit. However, in a continuous process, there is the possitiility
of retaining an undissolved amount of jas in the polyrner matrix, if excess gas is injected,
and therefore it is critical to ensure that the amount of gas injected must be below the
solubility limit at the processing conditions. The main advantage of utilizing the extrusion
foarning process, apart from it being the most cost-effective method among nurnerous
foaming processes, is the reduction of dissolution time due to higher gas diffusivity D at the
high processing temperature.
Although the proper amount of gas injection is emphasized, it does not necessririIy
guarantee formation of a uniform solution because the time required for completion of gas
diffusion into the polymer is equally important. If the required gas difision time is ronger
than the residence time of the melt, Le., the time between gas injection and nucleation, it is
obvious that a uniform solution would not be achieved.
Park et al. [18,108] studied the diffusion phenomenon in an extrusion process
containing a mixing screw. It was observed chat shear mixing promotes convective diffusion,
In convective diffusion, a high gas concentration region (gas bubble) is brought into contact
with a 1ow gas concentration region (polyrner melt). Moreover, due to stretching of gas
bubbles, in the shear field generated by the screw motion, the diffusion process is enhanced
as the interfacial area is increased. The tirne for completion of dissolution was estimated.
They also proposed use of a dissolution enhancing device containing static mixers in the
extrusion system and claimed that it would promote the dissolution process by generating
shear fields as the mixing elements repeatedly reorient the melt dong the flow direction, thus
promote solution formation.
2.3.2 Nucleation
Nucleation is a critical step in the fine-celled foaming process because a large number
of cells must be generated in order to achieve smdl-size celIs. Nucleation of bubbles in the
polyrner cm be modeled by the cIassica1 nucleation theory [109,110], which was originally
developed for a single-component system where the second phase is created by evaporation
of the iiquid when superheated. Blander and Katz [IO91 extended the classical theory to a
diffusion system in which one component is volatile and foms bubbles.
A diffusion process of the dissolved component into the nucleacion site _ooverns
bubble nucleation, and fine cells can be nucleated either randomly throughout the polymer
matrix that is known as homogenous nuckation, or prefenbly at certain sites that is known as
heterogeneous nucleation. Heterogeneous nucleation, which can be promoted by using
additives in the liquid, is more prefernble, since less energy is required for nucleation
compared with homogenous nucleation and bubbles are more likely to nucleate at the
preferred sites provided by the additive particles.
Homogenous Nucleation
Colton and Suh [ i l 1,1121 modeled the nucleation behavior in microcellular foaming
using the classical nucleation theory. According to this theory, the work required to genente
a bubble of radius r in a body of a liquid can be given by (1 111:
where the first term (y,, A,) is the work required to create a bubble with a surface area Ab
and a surface tension ~pb, and the second term (MV,) is the work done by che expansion of
gas inside a bubble of volume Vb. The difference between the gas pressures inside the bubble
and the surrounding matrix. AP, is approximated to be the saturation pressure [ I l l]. By
replacing Ab and Vb with the area and volume of a sphere nuclei, Equation (2.7) becornes:
Figure 2.2 depicts the variation of energy W with radius r. it is secn that there exists a
maximum energy, which must be overcome to forrn a nucleus that continues to grow
spontaneously. if the energy induced is less than this maximum energy (or free energy
barrier), the bubbles (with radius rce) will collapse. This activation energy for
homogeneous cell nucleation AG,:,,, is calculated by differentiating W with respect to r
which yieIds [ I 10,1111:
The nudeation rate is then represented by [110,112]
where C,, = concentration of gas molecules in solution (#lm3),
f;, = frequency factor of gas molecules joining the nucleus (Ils),
k = BoItzman constant (JK).
The classical nucleation theory for a homogenous system predicts that the higher the
saturation pressure AP, the greater the number of ceIIs nucleated because of the lower
activation energy for ceIl nucleation Il?]. This effect has been experimentally verified in the
batch process [19,111- 1 151. The saturation pressure in a batch process corresponds to the gas
concentration in the polymer as given by Henry's law (Equation (2.3)). Thus, the effect of
pressure can be intetpreted as its influence on the amount of gas absorbed into the polyrner.
When the amount of gas increases, the chance of nucleation is higher and ri larger number of
nucleated cells is achieved.
Park et al. studied the effect of the pressure drop rate in an extrusion process and
indicated that the pressure drop rate is influencing the number of cells nucleated [17]. They
induced various pressure drop rates by using viirious nucleation dies and examined the final
foam structures. Their study shows that the higher the pressure drop rate, the greater the
number of cells nucleated, which may be explained by the mechanism of cell
growthhucleation cornpetition [17]. When the pressure drop rate is high, the poiymer-gas
system experiences a certain pressure drop in a shorter tirne period. During this shorter time
period, the already nucleated cells grow Iess because less time is available for gas diffusion
to the cells. Therefore, less gas is consumed for ce11 growth, and more gas is rivailable for
further nucleation in the polymer matrix (Equation 2.t0, Co is higher). As a result, the final
foam structure contains a Iarger number of cells, and the cells are smaller since Iess gas is
consumed for each ceIl. Hence, it is essentid that a higher pressure drop rate be induced to
achieve a fine-celled structure in the foaming process.
Stewart's work on elastomers [Il31 shows that the number of nucleated cells
increases with increasing the temperature. Similarly, the results obtained by Kumar et ai.
[116] on PVC also show an increasing effect of temperature on ceIl nucleation, and Ramesh
et al. [117] demonstntd that in a PS-CO: system, increasing the temperature increases the
cell density. Goel and Beckman [ L 141 studied the nucleation behavior of a PMMA-COI
system and demonstrated that increasing the foaming temperature tends to decrease the
number of nucleated cells. Baldwin et ai. (1 181 show that iincrerising the foaming temperanice
increases the ceIl density in amorphous PET and CPET below 100°C. and produces no
further change above 100°C. On the other hand, they found that temperature does not
significantly affect the cell density in semicrystailine PET and CPET. They suggested that
nucleation occurs rnainly during the pressure release andor in the early stage of heating.
In surnmary, generating a large number of bubbles by dissolving a large amount of
gas in the polymer is a critical step to obtain a fine ce11 structure for polypropylene foarns-
However, ensuring the amount of gas dissotved in the polyrner below the solubility limit
according the processing conditions is equaily important. Nevertheless, a sufficiently high
pressure is required, as stated in Henry's Iaw, to maintain a large amount of gas dissolved in
the polymer.
Heterogeneous Nucleation
Bubble nucleation is heterogeneous when it initiated at some preferred sites by
mixing the polymer with an additive. in generaI, nucleation tends to occur at the boundary of
the matrix and additive rather than inside the polyrner matrix as with hornogenous nucleation.
At the boundaries of the matrix and additive, the fee energy barrier for nucleation is lower
than that in hornogenous nucleation (Figure 2.3); therefore, nucleation is more likely to occur
heterogeneously rather than hornogeneously. By controlling the amount of additive, one cm
generate the desired number of bubbles [lI9-1211. However, in general, it is difficuit to
generate a large number of cells of micron size using additives due to poor dispersion [Il91
and agglorneration [121] of the additive particIes. Using an appropriate nucleating acgent that
consists of very srnall particles (less than a micron) well dispersed in the polyrner matrix
without agglomeration, one could produce a large number of cells for fine-celled foam
application [119].
The shear force is aiso
system such as extmsion [122,
bubbles increases. Lee [122,123
affecting the heterogeneous nucleation rüte in a dynarnic
1231; as the shear force increases, the number of nucleated
1 developed a lump cavity nucleation model, which indicated
that the cavities on the rough surfaces of the tiny nucleating particles form potentiai sites for
bubble nucleation. M e n the gas phase in the cavity grows and rnatured by difision of the
dissolved blowing agent into the cavity or by a pressure drop, the applied shear force
enhances the chance of detaching it from the cavity to generate a bubble.
2.3.3 CeIl Gmwth
Ce11 growth is due to the continuous difision of gas from the polymer matrix into the
nucleated cells since the soiubility of grts in the polymer is decreasing with the pressure drop
(Henry's law). Concurrently, because the pressures inside the cell is greater than the pressure
in the surrounding matrix, the cell tends to grow to rninimize this difference [i l . The
viscosity, the diffusion coefficient, the gas concentration, and the number of nucleated
bubbles are goveming the growth dynamics. Foc instance, when the polymer viscosity
decreiises via a temperature increase, the rate of growth increases due co the decrease in
resistance against cell growth. Conversely, ceIl growth terminates when al1 the gas ciissolved
in the polymer rnatrix is depleted or the matrix is too stilf to allow further growth. Figure 2.3
rnodels a nucleated ceIl inside a polymer matrix charged with a biowing agent (such as gas).
When the ceil is nucleated, the gas concentration around the cell decreases. This generares a
gradient of $as concentration around the cell, which is responsible for funher cell growth.
The dynamics of cell growth have been extensiveiy studied [t24-132,137-1391,
Arefmanesh et al. El301 modeled the cell growth in a finite sea of liquid flowing in a
mold by considering the pressure drop along it. They described the effect of various
parameters on the dynamics of the growth and simuiated the effects of the diffusivity and
viscosity on cet1 growth. It was shown that the higher the diffusivity, the more rapid the cell
growth. In contrast, increasing the viscosity retarded the growth process. The results showed
that in the initial stage of growth, the viscosity hm a greater effect. The effect becomes
weaker as growth proceeds so that eventuaily the rate becomes the same at different
viscosities.
Han and Villarnizar [13 11 investigated the growth of bubbles in a foam sheet
extrusion process. They studied the pressure profile of foam along the die and found that
when the fiow rate increases, the point at which foaming becomes visible moves closer to the
die exit. On the other hand, a higher gas concentration and higher temperature tend to move
the foarning point away from the die exit. This observation is important since early ce11
growth causes deformation of ceIIs to a non-sphericd shape due to the presence of a shear
field dong the die, which mrty promote ce11 codescence- At a higher flow rate the residence
time is shorter, and hus the already nucleated cells grow less. As a resuIt, it is Iess tikeiy that
adjacent celts codesce, which would cause foam degradation.
Villamizar and Han [132] studied ceII growth in an injection molding process. They
studied the effect of the rnelt temperature, injection pressure and blowing agent concentration
on the final bubble size. It was observed that at Iower rnelt temperatures, bubble growth is
slowed down, and also fewer bubbles were detected by the naked eye. This was due to the
increase in viscosity which opposed the rate of ce11 growth and the lower diffusivity
(Equation (2.7)), which means that a less blowing agent was supplied ta the bubble. Their
results also showed that as the arnount of blowing agent increased, the final ce11 size became
greater. it is interesting to note that when the blowing agent concentration exceeded the
solubility limit, a small number of large bubbles appeared. Therefore, the arnount of injected
blowing agent must be maintained below the solubility limit. Their research also showed that
an increase in the injection pressure (thus, a reduction in the filling tirne) resulted in a
decrease in the ceIl size and a uniforrn cell distribution. This effect is similar to that of the
flow rate, which was explained earlier* As the filling time decreases, the amount of ce11
growth diminishes and the ce11 size rernains smaller. Although the final number of cells is not
mentioned, it is believed that the cell number increased as the injection pressure increased.
This corresponds to the effect of pressure drop rate as explained in Section 2.3.2, A shorter
filling time at a higher injection pressure results in a higher rate of pressure drop, which
favors the nucleation of a larger number of cells.
Ce11 growth has uusally been modeled wiihout the assumption of possible Ioss of
blowing agent. The loss of blowing agent or gas escape is significant in a batch process
although the gris loss is not easily noticed. [n a batch process, the sample skin has the highest
temperature and this prornotes the gas escape to the environment [133]. This is why the
microcellular foarns produced in a batch process have a typically low volume expansion ratio
in the range of 1.5 to 10 times although the maximum possible volume expansion ratio is
about 35 times for 5 wt% CO2 concentration. However, the loss of blowing agent in a batch
process can be minimized by Iowering the foming temperature and thereby Iowering the
diffusion of gas. if the foaming temperature increases too much, the amount of gas escape
becomes significant so that the final foarn expansion will be dtamaticaily reduced 11341.
On the other hand, for foaming in a molding process [131,132], the poIymer in the
mold is mechanicdIy constrained by the mold wails. Therefore, the volume expansion ratio
is determined by the mold cavity to shot size ratio regardIess of the gas escape through the
foam skin.
Lee and Ramesh [l20,135] scudied the effect of gas loss theoretically and
experimentally and showed that the amount of the Iost gas is influenced by the number of
cells nucleated. The lost gas amount decreases with an increase in the number of cells up to
500 cells/cm3, and then becomes almost insensitive above 500 to 2000 ce11slcm3. They
attributed this effect to the fact that less time is available for cell growth and gas escape. As
the number of cells increases, the wali thickness between the cells decreases, and thus the
distance decreases for gas diffusing from the macrix to the cell. As a result, the rate of ceII
growth is faster when the number of cells is greater. Moreover, because a large number of
cells are nucleated, the final ceII size should be smaller. Thus, a cell must grow faster to a
size which is smaller, and so the growth time is much shorter. A shoner growth time allows
less time for gas escape from the sheet surface and chus the expansion increases. Therefore,
the final foam density decreases as the number of ceI1s increases.
Lee et al. [136] studied the effects of foam sheet thickness and nucleation on
thennoplastic foam sheet extrusion. The results show that a decrease in sheet thickness
reduces the final cell size and increases che foam density (Le., decreases the amount of
expansion). They attributed this effect to the influence of heat transfer and also to the
possible gas escape from the surface of the sheet to the atmosphere rit the time of growth. Lee
et al. [136] claimed that at a Iarge thickness, the sheet behaves like an insulator thus
enhancing the bubble growth. As the thickness decreases below 0.5 mm, the amount of final
sheet thickness (or volume expansion) decreases significantly. This effect can be further
described by a surface-to-volume ratio concept- Decreasing the sheet thickness increased the
surface-to-volume ratio of the foam sheet. Since the gas loss occurs at the foam surface, this
Ioss is promoted by increasing the surface-to-volume ratio. Thus, the amount of the lost gas
is greater resultïng in a lower expansion (or a higher foam density) in a thin sheet.
Shafi et al, [137] developed a mode1 that combines nucleation with bubble growth,
They studied the effects of processing variables such as the pressure and the dissolved gas
concentration on bubble growth dynamics, nucleacion and final bubble size distribution
during free expansion polymer processing. The mode1 was based on an expansion influence
volume where each bubble is assigned a region depending on the nucleation threshold, and
this specified volume expands as the bubble grows. They found out that the growth rate
increased with an increase in the solubiIity of giis in polymer. They also claimed that the final
bubble size distribution depends on both nucleation rate and bubble growth dynamics, and
that lowering surface tension significantly increased nucleation rate and resulted in a much
narrower bubble size, a higher bubbIe density, and smaller bubble sizes.
Shimoda et al. [138] conducted a numerical simulation for polymeric foaming
extrusion processes. They combined the classicai nucleation rate and bubble growth rnodels
with a non-Newtonian fluid model of a flow in order to simulate bubble growth and
nucleation in a flow field. They examined the effects of influence volume region and the
initial equilibrium vapor pressure of bubbles on bubble size and number density calculation.
They claimed that the processing parameters such as diffusion viscosity is strongly affecting
the simulation of the ceIl nucleation and growth,
Ramesh et al. [139] developed a new bubble growth model that includes the effect of
blowing agent concentration, temperature effects on physical properties during foam
formation. Experimental data confirmed well the mode1 predictions. Gas loss, blowing agent,
and transient cooling effects are shown to be the most important factors. As found by
Rarnesh et al.. predictions of old foam models differ significantly from experimentally
observed values, since they ignored inff uence of blowing agent concentration effects in the
binary system [139].
in conclusion, ce11 growth is governed by many parameters such as the final foarn
density, ce11 size, and distribution, and can be controlled by the temperature which influences
the diffisivity and melt viscosity. Tempenture control should be performed before ce11
nucleation since the rate of cell growth is much higher in the initial stage where the cells are
small [Il. in the foaming process, it is important to consider the gas loss from the extrudate
dthough the effect of gas Ioss is generally not dnmatic [120,133,136]. However, at a high
melt temperature, where the diffisivity is high and the viscosity is low, gas escape couid be
vigorous. This could result in a dramatic drop in foarn expansion, resulting in an undesirably
high foam density. Moreover, when smailer cells are present, the wall thickness separating
the two cells becomes weaker, and thereby the rate of growth is faster, which may cause
rupture in the ce11 wall and ce11 coalescence [135J.
2.3.4 Crystallization kinetics
Semicrystailine polymers are polymers in which crystallites are dispersed into an
arnorphous matrix. The fraction of the polymer that is fully crystalline is known as the
crystallinity (or the degree of crystallinity). When a polymer is super-cooled by lowering the
temperature below the equiIibrium melting temperature, crystallization takes place and
normally proceeds in two stages: nucleation, and growth. When a long chain molecule srarts
to fold back on itself repetitively, it forms chain-folded lamellae radiating from a nucleus and
nucleation is said to have occurred. The folding length of the crystal is temperature
dependent and increases with the crystallization temperature [140,141]. The broad range of
melting temperature observed in semicrystalline polymers depends on the crystallinity
distribution, and is attributed to the distribution of chain folding length.
During nucleation, the crystallites are nucleated from the melt at a definite range of
temperature and subsequently grow during the growth phase to forrn three-dimensional
aggregates of crystallites known as sphenilites. In generai, the spherulite growth rate is faster
than the nucleation rate because the free energy requirement of the former is Iower [L42].
Nucleation is therefore the rate-determining step of polymer crystallization. When
crystallization occurs under stress, as in the extrusion process, the overall crystallization rate
will increase because the stress could cause orientation of the molecular chains thus making
them more packable [142].
Nucleation can occur in one of two ways; when nuclei appear instantaneously at the
beginning of the process, then athermal or instantaneous nucleation has occurred and it is
assumed to depend only on temperature and to be independent of time and cooling rate. The
grown crystais will therefore be of approximately equai sizes. if nuclei appear in the liquid
phase during the process, thermal or sporadic nucleation has occurred, and the activated
nuclei appear at a constant rate per unit volume.
Nucleation and Growth Rates
In gnerai, two types of nucleation are defined for polymer crystallization: primary
and secondary nucleation. During primary nucleation, when a potentiai nucleus reaches a
critical size, cqstai growth occurs quickly and spontaneously and a three dimensional crystal
is genented, the rate of which depends on temperature. Pnmary nucleation is caiied
homogeneous nucleation if no preformed nuclei or foreign surfaces Xe present. Secondary
nucleation then follows, where chah segments are dded to the existing crystal surface. The
essential difference between primary and secondary nucleation is the energy of formation of
a nucleus of critical size or the Gibbs free energy,
The classical nucleation theory developed by Gibbs [143-1451 is based on an
assumption that energy fluctuations in the supercooled phase can overcome the nucleation
barrier caused by the surface of the crystal. Based on this assumption, Turnbull and Fisher
[146] derived an expression to determine the primary nucleation rate as a function of the
crystallization temperature, using the Williams-Lmdel-Fe;erry (WLF) Il471 equation which
universally describes the temperature dependence of pdymer melt viscosity:
Based on the surface or secondary nucleation theory, Lauritzen and Hoffman [148-
1501 formulated a linear growth rate equation, which incorporates fold surface energy, lateral
surface energy, heat of fusion, and lameIlar thickness ternis into the Gibbs free energy to
describe the linear growth rate of spherulites.
Crystullizution Regimes
The relative rates of nucleation and deposition of chain segments on the crystd
surface will affect both the crystal growth rate and the sphemlite size. Hoffman [151.1521
ctassified several crystailization regimes, which describe the relacive nucleation and growth
rates. A regirne transition occurs when the ~lationship becween growth rate, G , and the
surface nucleation nte, i, undergoes a change. In regirne I, the highest temperature regime.
the chain segment deposition rate on one surface nucleus is so fast, that the crystal unit
growing on an existing crystd face is complete before the next layer is nucleated, in other
words, G varies as i. In regime il, the nucleation rate is fast compared to the growth rate,
: M ir. G = i . consequently ihe nucleation of new crystal layers occurs before deposirion on
existing Iayers is complete. This results in a downward break in the growth rate curve as one
passes through the regime 1 to regime II transition. Finaily, in the Iowesc temperature regirne,
Oregime ID, the mean separation of the nuclei approaches the width of rhe molecular stems,
and G = i again, such that at the regime II to regime III transition, an upward break in the
growth rate curve occurs. Other kinetic theories of crystailization were fomiulated by Frank
and Tosi [ 1531, Sanchez and DiMarzio [154], Sadler [155,156], and Lauritzen, DiMarzio and
Passaglia 1157-1591.
Avrami's ïïleory
Knowledge of the primary nucleation rate and the linear growth rate is usually
sufficient to calculate the overall crystallization rate. Many rneasurernents of crystallization
invoIve the macroscopic determination of crystallinity as a hnction of time. The first efforts
to quantitacivefy describe the rnacroscopic development of crystallinity in tems of nucieation
and linear crystal growth were made by Kolmogoroff 11601, Johnson and Mehl [161], and
Evans [162]. However, the classical theory of Avrami [163-1651 for phase transformation
kinetics is the most widely quoted mode1 For the analysis of isothermal nuclerition and
crystallization in polymer processing.
Despite 'its wide use, the Avrami mode1 suffers from some limitations caused by
simplified assurnptions. These assumptions are as follows: (i) there is no volume change
during crystallization; (ii) the sampie is completely trrinsformed; (iii) there is constant linear
growth rate; (iv) the nuclei have constant shape during growth; and (v) there is no secondary
crystallization occurs,
The implications and complications caused by these assumptions are dealt with in
detail in the literature [166]. It is sufficient at this point to state that macroscopic observation
of the increase in crystallinity cannot adequately descnbe the rnicroscopic mechanisms of
crystal gowth. The Avrami analysis sirnply provides a convenient representation of the
macroscopic data.
Diflerential ~canning Calorimeter
DSC is a powerful tool in the determination of the Avrami parameters as desccïbed in
his theory. Determination of polyrner crystallinities using the heat of fusion is often based on
rneasuring the area of DSC rnelting peaks above a chosen baseline. During the
crystdlization of a sernicrystdline polymeric material, heat is rejected as the rnotecules
become ordered and form crystallites. The heat evoIved (Mldt) is rneasured and recorded as
a function of time by the DSC, and the weight fraction XJt) of materiai crystallized aiter
rime t cm be cdculated from the relation:
where X,,(t) is the absolute crystallinity at crystallization time t, and X, is the ulcimate
crystallinity for t==. For isothermal crystallization experiments, the heat evolved is
evaluated while the polymer is in the isothermal condition, consequently, XJt) c m be
physicaliy obtained as the area under the crystailization peak in a plot of heat tlow versus
cime. Since the Avrami equation is expressed in terms of the volume fraction, it is necessary
to transform the weight fraction measured by DSC into a volume fraction. This can be done
using the following relation:
where p, is the morphous region density, and p , is the crystdline density.
in conclusion, the Avrami ana1ysis is a useful tool for the representation of the
macroscopic data of polymer crystallizrition. Based on this analysis, the main parameters
governing the crystdlization kinetics under isothermal conditions are the primary nucleation
rate and the spherulitic growth rate. Physical investigation of the crystallization phenomena
can also be conducted using DSC method by obiaining the heat evolved as a function of time
under various isothemd conditions. The resulting crystaltization kinetics cm be used as a
bais for establishing strategies for the production of low-density, fine-celied potypropylene
foarns.
Figure 2.1: Sorption isotherm for general gadpolymer system
A
Free Energy,
AG
+ Bubble Radius, r
Figure 2.2: Effect of the variation of energy on the bubble growth for homogenous nucleation
polymerlgas solution
Figure 2.3: Model of a nucleated ceIl inside a polymer matrix
Chapter 3
Design & Construction of the Experimental Equipment
3.1 Conceptual Design of Tandem Extrusion System
3.1.1 Introduction
An axiomatic design method developed by Suh at the Massachusetts Institute of
Technology [167] was employed to create a conceptual design of the tandem extmsion
system. in brief, this design method starts with the definition of Functional Requirements
(FRs) from the perceived needs; this defines the design problem. With the FRs, one can
generate a number of possible physical entities corresponding to each FR. These physical
entities are called Design Parameters (DPs). The retationship between the FRs and the DPs
can be expressed by a matrix equation:
{ FRs } = [A] { DPs} , (3-1)
The elernents in the A matrix cm be either "X", which denotes a suong relationship between
the corresponding FR and DP, or "O", which denotes a weak or absent relationship. The A
matrix presents the cornplete retationships between the FRs and the DPs; it can help the
designer better understand and improve the design.
3.1.2 Analysis of the Tandem Extrusion Systern for Foam Processing
The four functionai requirements (FRs) identified for producing low-density, fine-
celled polyrner foam are as follows:
RI = plasticarion of polymer;
FR2 = formation of a polymer/gas solution;
FR3 - - nucleation; and
Flt - - expansion.
To satisfy the above FRs, the following design parameters (DPs) are proposed:
DPI = heat provided by the plasticating screw motion in an extruder and
externally mounted band heaters;
DP2 = gas injection, diffusion system and gear pump;
DP3 = thermodynamic instability created by a nuclerition die: and
DP4 = a cooling system that can cool the polyrner melt without decreasing
its pressure.
in response to FRi, a single-screw extruder is chosen to plasticate the polymer. There are
two heat sources in the system. The primary source is the frictional heat generrtced by the
motion of plasticating screw. The secondary source is the externaily mounted band heaters.
The band heaters are important in the start-up period but they are not suitable as the only heat
source for this process due to the low thermal conductivity of the polymer.
FR2 is satisfied by using a gas injection system with a metering device, and a
diffusion enhancing device to produce a homogenous polymerlgas solution. The amount of
gas is metered by a gas injection pump, which supplies the gas under high pressure into the
polymer melt in the extruder. The resuIt is a two-phase polymerigas mixture, The shear field
generated by the motion of plasticating screw stretches the gas bubbles and increases the
interfacial surface area when the polymer is conveyed in the extruder barrer. As a result, the
gas can diffuse into the polymer matnx more quickly. To further assist the diffusion process
and homogenize the polymer/gas solution, a diffusion enhancing device consisting of static
mixers is employed.
in response to FR,, a rapid pressure drop is chosen to generate thermodynamic
instability and thereby promote high nucleation density. The thermodynamic instability is
created by a sudden drop in gas solubility in the polymer/ gas solution. As mentioned in
chapter 2, the gas solubility in polyrner is proportional to the pressure of the polymer melt. A
rapid drop in the pressure results in a rapid decrease in the solubility of gas in polymer melt.
Ttierefore, a high nuclei density can be created by dropping the pressure of the polymer/gas
solution rapidly.
A coaling system is employed to satisfy FR+ The most dominant parameter that
affects the growth of the nucleated cells is the melt temperature. As the gas from the
polyrner melt diffuses into the nucleated cells, the concentration of the dissolved gas
decreases in the region near the cells. This results in a concentration gradient in the melt,
which drives the dissolved gas towards the cell. The diffusion of gas into the cells causes the
cells to grow. The gas diffusion rate increases with temperature, and therefore, if the melt
temperature is too high, the gas cm easily escape from the polymer to the environment
instead of contributing to cell growth. A high temperature can also promote ceIl coalescence
since the melt strength decreases when the temperature increases. As the cells grow, the
walls between them will be stretched and could easily be broken due to the weak melt
strength. Thus, adjacent cells will join together, and the ceIl structure and the cell-population
density wilI deteriorate. Since gas diffusion and ce11 coalescence can be controlled by
lowering the temperature, a cooling system is chosen for satisfying m. A cooling system
consisting of a second extruder is introduced.
The FRs-DPs relationships cm be described in the following matrix:
The diagonal elements Aii of the design matrix are al1 "X", simpIy because each DPi
is chosen directly to accomplish the corresponding FRi- An examination of al1 the non-
diagonal elements of the matrix is required in order to determine the effects of each DPi on
the ather FRs.
Since gas is injected in the polymer after the polymer is cornpletely plasticated, it has
no effect on the plastication process. As a result, element AL2 should be zero. EIernents AL3
and A14 should aiso be zero because nucleation, expansion and shaping take place further
down strearn and therefore, will have no effect on the plastication stage. The shear field
generated by the plasticating screw motion affects the polyrnerlgas mixing. Therefore,
elernent should be non-zero. Element A3 should be zero because nucleation takes place
after formation of the polymerlgas solution. The amount of gas in the polymer melt is
dependent on De4. The addition of a gear pump decouples FR3 frorn these DPs. The gear
pump ensures a constant polymer flow rate as long as the gear speed and the extruder speed
are not varied. This feature implies that the amount of gas injected into the polyrner meIt,
once set, will becorne independent of the temperature variations and the die exit settings.
Therefore, elernent AZ4 is zero.
The spindle speed of the plasticating screw affects the systern pressure and therefore
affects the nucleation rate. As a result, element Afl should be non-zero. Element A3? should
be non-zero too, because the arnount of gas in the polymerlgas solution affects the nucleation
density. Elernent A3J should be zero because the cooling system of a second extruder can
cool the polymer rnelt without sacrificing its pressure.
As desciibed previously, the temperature of the polymerlgas solution affects the
volume expansion ratio. Thus, hI should be non-zero because the shear heat generated from
the piasticating stage affects the temperature of the polymer flow, Element should be
non-zero since the arnount of gas injected affects the expansion ratio. Elernent should be
non-zero because gas loss to the environment is localized if the cell-population density is
high.
From the above considerations, Equation (3.3) can be written as:
X X X O DP, X X X X - DP,
Equation (3.4) is a lower triangular rnatrix, which irnpties that the functional requirernents
can be achieved if the design parameters are implemented in the proper sequence.
Having arrived at a satisfactory decoupled design at the parent levet, it is necessary to
decompose these FRs and DPs to lower level hierarchies and zigzag between the functional
domain and the physical domain to arrive at an appropriate design for satisfying the initial
FRS.
3.1.3 Detailed Analysis and Further Decomposition of the FRs & DPs
Decomposition of FR4 & DP, (Second Level)
The cooling systern is required to perform two main functions. FirstIy, the melt
temperature has to be uniformly lowered to the optimum level, without losing the processing
pressure, in order to prornote ceIl growth without coalescence. Secondly, the surface
temperature of the extrudate has to be further reduced to prevent the gas from escaping
prernaturely to the atrnosphere through a weak surface layer. Therefore, the two FRs for this
syseem are:
FR,, = Cool the polyrner melt homogeneously to the optimum level
without losing the processing pressure;
FR,? = Cool the surface of the extrudate to prevent escape of gas to the
atmosphere.
In response to FR4,, a cooling extruder and a cooling section are incorporateci. The
shear mixing in the extruder can hornogenize the rnelt temperature whiIe cooling it, and the
thrust force provided by the screw motion can maintain the processing pressure. The cooling
section has a sleeve through which a cooling fluid flows to cool the polymer funher. The
static mixers in the cooling section ensure the homogeneity of temperature in the melt.
Although there is sorne pressure loss due to the cooIing in the cooling section, the high
pressure provided by the cooling extruder can compensate for this loss. To cool the surface
of the extrudate, an appropriate system is chosen. The DPs are then given as:
DP,, = Cooling extruder and cooling section;
DP,? = Polyrner surface cooling system.
The design equation can be represented as:
A,? is zero because the surface cooling cannot affect the inside temperature of the melt due CO
the low thermal conductivity. However, if the melt temperature is highec than the desired
temperature, the cooling of the surface cooling system will have to be increased ro achieve
the required surface temperature; this leads to the conclusion that A,, is non-zero. Equation
(3.6) wiIl become:
Therefore, the cooling system design is decoupled.
The polymer surface cooling system can be further decomposed to the third Ievel in
order to arrive at a proper design.
3.2 Detailed Design of the Tandem Extrusion Line
3.2.1 Ovewiew of the System
Based on the conceptual design of the overall system for low-density, fine-celled
foam extrusion in the previous section, the design of the detailed components is carried out in
this section. The tandem extrusion system consists of two single-screw extruders, gas
injection equipment, a gear pump, a diffusion enhancing device, a heat exchanger, and a
fitament die. The first extruder is used for plasticating the poIymer min, the gas injecrion
equipment is used for injecting gas into the polymer melt, while the second extruder provides
mixing and initial cooling for the polymer mek The gear pump controls the polymer melt
fIow rate, independent of temperature and pressure changes, the diffusion enhancing device
ensures the homogeneity of the polymer/blowing agent soIution and the heat exchanger
provides further cooling for the polymer melt to suppress cell coalescence. Shaping and ceIl
nucleation are accomplished in the die. Figure 3.1 shows a schematic of the tandem extrusion
system. The detailed design of each component is presented in the following sections.
3.2.2 The First Extruder in the Tandem Line
The extruder used in the single screw extrusion system consists of a W1 laboratory
extruder (Brabender: 05-25-000) with a 5 hp variable speed drive unit (Bnbender: Prep
Centre, Model D52T). The screw is a single stage mixing screw (Brabender: 05-00-144)
with a 30:l UD ratio. The purpose of the mixing stage is to enhance the mixing of the
blowing agent and the polymer melt. A schematic of the mi~ ing section is shown in Figure
3.2.
3.2.3 Gas Injection Equipment
The two main components of the gas injection equipment are a positive displacement
syringe pump and an in-house designed gas injection port. The pump is capable of injecting
the blowing agent into the polymer melt at a maximum pressure of 51.7 MPa (7500 psi) with
a wide range of flow rates ( fom O.OI mllmin to IO7 mI/min, depending on the pressure
condition).
At the hem of the gas injection port is a flow restrictor. Because of the compressibility of
gas, the pumping of gas into the barre1 is readiIy affected by the pressure fluctuation in the
barrel. Variations in the injection rate could affect the consistency of foaming significantly.
One solution is to maintain a high pressure difference between the gas injection pump and
the barre[. The choice of restrictor depends on the required pressure difference and the
desired gûs flow rate.
3.2.4 The second Extruder in the tandem Line
The second extruder in the tandem extrusion system consists of a 1%" extruder
(KiIlion: KN-150) with ri buïit-in 15 hp variable speed drive unit. The screw motion c m
generate shear heat in the polymer melt, however. since the second extruder is intended for
cooling, it is necessary to keep the generated shear heat to a minimum. In this context, a
smdl iength/diarneter (Ml) ratio of the screw (18:l) was chosen in the screw design.
Moreover, a compression ratio of 1: 1 was chosen since the second extruder is intended for
rnaintaining the pressure of the polymer melt.
Polymer rnelt is fed into the barre1 of this extruder from the firsc extruder through an
injection port. Typically, the injection pressure of the melt is around 27.58 MPa (4000 psi).
Because of the high pressure, it is necessary to have a dynamic sealing in the second extruder
to prevent the melt Leaking backward into the motor assernbiy through the clearance between
the screw and the barrel. The proposed solution was to miike extra flights on the screw
Iocated before the injection port. As the screw rotates, the thrust force generrited by the
screw motion of these extra fiights would push the polymer melt forward preventing the
leakage to the back. A detailed design of the second extruder was proposed by Young [I68].
3.2.5 Gear Pump
A gear pump (Zenith: PEP-I l), with a '/2 hp motor drive (Pacific Scientific: Model
SR), a speed control unit (Zenith: ZeDrive) and a temperature controller (Eurothem Controls:
Mode1 94), is used in the extrusion system to control the polymer solution flow me. The
gear purnp consiscs of two closely intermeshing gem that rotate in a counter-rotating manner
to convey the polymer melt from one end to the other. A schematic of the gear purnp is
shown in Figure 3.3.
3.2.6 Diffusion Enhancing Device
A diffusion enhancing device is used to ensure the polymer melt and the blowing
agent are mixed hornogeneously. It was originally designed by Park [169j and modified by
Behravesh [170], It consists of an in-house design, rniId steei body, a static mixer (Omega:
FMX844 lS) enctosed in a mild steel case, and two band heaters (Omega: MB 1GlJ I A t &
MBIG2A1 Al). The rationale behind this design is to use the static mixer to promote shex
mixing, and to maintain a high melt temperature to promote a high difision rate of the
blowing agent into the polymer rnelt.
in order to determine the required number of static mixer elements, a calculation was
done based on an equation provided by the manufacturer's technical brochure [L? LI:
where Re = Reynolds number,
Q = flow rate, gallmin
S = specific ,mvity
p = viscosity, cP
d = inside pipe diameter, in
A typical flow rate (Q) of the system is in the range of 5-12 cm3/min (1.3-3.2xl0-' gallmin).
The specific gravity (S) is approximately 0.9 for polypropylene, and the inside pipe diameter
(6) is 0.0127m (0.5 in).
The viscosity is influenced by the shtar m e and temperature. The apparent shear rate
( j,,, ) in a circular channel can be calculated by the following equation [17 11:
With a channel diarneter (4 of 0.07 m (0.28 in) and a flow rate (Q) of 10 cm3/rnin (2 .64~10~~
gdrnin), the apparent shear rate is approximately equal to 4.95 11s. From the manufacturer's
data sheet (Borealis), the apparent viscosity of bmched polypropylene is 916.108 Pa's
(916108 cP). With these values, the Reynolds nurnùer can be cakulated using Equation
(3.7):
The pressure arop across the static mixer c m be detennined using the following equation
[171]:
where I = laminar factor (from the manufacturer), and
A P = pressure drop, psi.
With a flow rate (Q) oF0.303 m3/min ( ~ . ~ L c x I o - ~ gal/min), a viscosity (p) of 916108 cP. and a
larninür factor (0 of 0.05, the vaIue of head loss in the static mixer is:
AF = 120.92 psi (0.8 17 MPa) (3. L 1)
This value is reasonable when compared to a typical system opemting pressure of 27.58 MPa
( 4 0 psi). Based on the calculated results, it can be concluded that a static mixer with six
elements cm fulfill the rnixing requirernents (1701, Additional mixing of the polymer melt is
provided by the mixing stage of the extruder's screw.
3.2.7 Heat Exchanger
As mentioned in Chapter 2, cooling the polymer meIt will increiise the melt strength
and suppress ce11 coalescence. It is important to cool the polymer melt homogeneous1y
because non-uniforrn temperature distribution could induce in-homogeneous ce11 growch,
resuIting in a irregular cell structure. The heat exchanger utiiized in this research was
designed by Behravesh [170]. It consists of a static mixer (Labcore: H-04669-12) encased in
a mild steel body with embedded cooling channels as shown in Figure 3.4. This static mixer
is stntcmrally different €rom the one used in the diffusion enhancing device. This static mixer
used in the herit exchanger promoted polymer transport in the radiai direction such that the
core material is constmtly exchanging with the boundary materiai, whereas the mixer of the
diffusion enhancing device does not.
Since the temperature cm directly affect the foaming behaviour of the polymer, it is
important to control the temperature of the heat exchanger precisely. Using band heaters
alone with temperature controllers do not provide adequaie control because of lack of a
cooting source, consequently, a high pressure gas is introduced CO provide cooIing for the
heat exchanger. The temperature controller controls the air flow through the cooling channel
in the heat exchanger using a solenoid valve based on the polymer melt temperature. When
the temperature is above the set point, the solenoid value opens to let the high pressure air
expmd and flow through the chmnel in the heat exchanger. Because of the cooling induced
during isentropic expansion [172], the temperature of air is reduced significantly and the cold
air removes the heat from the heat exchanger. When the temperature is below the set point,
the band heaters are switched on to provide the necessary heating. Using this band
heatedhigh pressure air arrangement, temperature control of &loC was easily maintained,
3.2.8 Filament Die
The principal function of the filament die is to induce a thenodynamic instability,
which promotes nucleation of a large number of cells. The section of the die, which induces a
rapid pressure drop, is characterized by a small diamcter and is referred to as the nucleation
section or capillciry section. In this case, two nucleation dies were designed to induce a rapid
pressure drop. For polypropylene/butane solutions, it was found that a die of 0.46 mm (0.018
in.) diameter and 7.62 mm (0.3 in.) length, and a larger die of 0.76 mm (0.030 in.) diarneter
and 12.6 mm (0.5 in.) length, resulted in a desirable processing pressure in the processing
temperatures that are of interest. These filament dies were therefore manufactured since the
foaming behavior of the polyrner rneIt solution at different temperatures was not known.
The effects of die diameter and length on the pressure drop and the pressure drop rate
have been explriined in details in the literature [169,170]. It is clear that a change in diarneter
affects the pressure drop rate, while a change in diameter does not influence the pressure
drop rate. However, both the die diameter and length affect the pressure drop in the die. For
instance, by decreasing the die length without changing the die diarneter, the pressure drop
decreases while the pressure drop rate remains the same. Thus, a shorter die may be used
when a high-pressure drop is experienced during processinp. Figure 3.5 show a schematic of
the filament dies used.
3.2.9 Cooling Sleeve
As discussed earlier, the strategy for preventing the gas escape is to cool the extrudate
surface whiIe it flows into the nucleation section of the filament die, This cm bbe achieved
through cooling of the die using high pressure air, to cool the entire nucleation section of the
die. Thus, there is a need to provide a seded channel around the die to circulate the
pressurized air. It would be inefficient to make a channel for each die. especiaIly when a
large number of dies are to be exiunined. An alternative design is to make a sleeve wirh a
large circumferentiai grwve, as shown in Figure 3.6. When the sleeve is mounted on the die
and sealed at each end using two O-rings, the cooling air can be circuhted througti the
channel.
Figure 3.1: Schematic of the tandem extrusion system
Figure 3.2: Schematic of the extruder mixing section
Figure 3.3: Schernatic of the gear pump (counesy of Zenith).
hot polymer melt -
cooling oil cooling
static mixer channel \ oil pressure transducer
seat
cotd polyrner melt
+
cooiing oil
high ternphanire thermocoup1e O-ring seat
Figure 3.4: Schematic of the heat exchanger (courtesy of Behravesh) 11701
Figure 3.5: A schematic of the filament die
Heat exchanging air grooves for
nucleation ,' nozzle \ - -
- -
circulation
Figure 3.6: A schematic of the cooling sleeve to be mounted on the nucleation nozzle
Investigation of Fundamental Properties of Polypropylene Materials with ~oaming Agents
4.1 Introduction
This chapter presents the investigation of various processing and materials parameters
on the themodynamic, thermal and melt fracture behaviors of polypropylene melts with
foaming additives under different processing conditions. This will aIlow us to develop the
strategies for achieving low-density, fine-celled polypropylene foams and to identify the
fundamentai mechanisms governing the volume expansion ratio of polypropylene foarns in
the next two chapters.
4.2 Measurements of PVT Properties of Polypropylene Materials
In this section, the effect of dissolved butane, processing temperature and pressures,
and materials bnnching on the PVT relationships of polypropylene matenals in a moIten
state are investigaced. The basic principle involved in the measurement of PVT properties of
polyrner/gas solutions is to measure the specific voIume of the solution by deterrnining the
mass and volume flow rates of the polymerlgas solution at ditierent temperatures and
pressures. In this context a dilatometer based on a foaming extruder with a new degassing
oven to facilitate the measurernent of the mas flow rate of polymedgas solutions from the
foaming extruder is presented [173]-
The system utiIized for the measurement of the PVT properties of propylenehutane
solutions is based on the previous system developed by Park et al. [60,61]. The previous
system was found to be inadequate for experiments involving butane. Since, the diffusivity
of COz in the polyrner is much higher than butane, the degassing of COz can be more easily
accomplished, For the case of butane, the gas does not emily go out of the polymer becriuse
of the low diffusivity of butane. An effective degassing system was cherefore designed and
implemented to achieve more accurate degassing of the butme from the propylene foams.
Xthough detailed information about the system cm be found in the previous pubrications
[60,6 11, the overall function of the foaming extruder-based dilatometer, and the procedure to
determine the specific density of propylenehutane solutions are briefly described below.
An important element in the measurement of the PVT relationships of the
propylenehutane solutions is the formation of a homogeneous and uniform single-phase
polymerlgas solution in a themodynamically well-defined pressure and temperature
condition. This is achieved through the use of a tandem foarning extruder shown in Figure
4.1. Melting of the propylene is achieved in the first extruder through the frictional shear
generated by the screw motion and through the use of extemally mounted band heaters.
While in the molten condition, gas is injected into the polymer in metered amounts using a
syringe pump. As the propylenehutane mixture is conveyed in the extruder barrel. the shear
fields generriteà by the plasticating screw and irregular mixing blades stretch the gas bubbles
under high pressure and result in dissolution of the gas into the polymer matrix [17,18].
After formation of the single-phase polymerlgas solution. the temperature of the solution is
controlled in the second extruder and further homogenized using a heat exchanger consisting
of a static mixer. A positive-displacement gear pump is then used to control the volume tlow
rate of the polymerlgas solution. A variable resistance valve is installed after the gear pump
to control the pressure of the solution flowing through the gear pump. The solution then
exits throush a small filament die, which facilitates collection of samples for analysis.
The specific volume of the propylene/butane solution foned at each temperature and
pressure was then determined using Equrition 4.1.
where subscripts p and g refer to the polyrner and gas respectively, and Q ~ ~ ~ ~ , , ~ ~ ~ ~ , , ~ ~ ~ ~ , and
mpçx(,,ighZ highP, refers to the volume and mass fiow rates of the propylenehutane soIution,
mesured in cm31g and glmin, respectively. The volume flow rate of the solution is
determined by:
The RPM of the gear pump was set constant for each series of experiments, whiIe the
throughput per unit revolution of the gear pump was determined from calibration
experiments performed by Park et al. [60]. The m u s flow rate of the propylendburane
solution was dererrnined at high experimentnl temperatures and pressures using Equation 4.3.
- mp+,v (IiighT. hjgh PI - mp lhighT. high PI
' ',g(higfr T. high P)
where rhp(highT.highP) represenrs the mass flow rate of the polymer in the polymerlgas
solution, determined with the aid of the degrissing system descnbed. The m a s flow rate of
the gas in the polyrner/gas solution, mX(,lighT,highP), was computed as:
- high T. high P) - mg.k;ispump(mom T.injcction P)
- - Qg.gu pvmp (momï'.injeciion PI
Pi(rwm~.injcction P)
The volume flow rate of gas in the :as pump.~,,pu,p ,,,,,,, lionPl, was collected for each
experiment by reriding the gris injection pump, whiie the gas density at room temperature and
gas injection pressure, pi (momZinjcnianP), was obtained from the thermodynamic tables for
butane [174]. The validity of equiition 4 holds only when the equilibrium point is reached.
The system was considered to be at equilibrium when there was no fluctuations in the gas
flow rate readings of the gas pump.
Since gas loss from the extruded foarn is unavoidable [20,2I], simpIy weighing the
collected extnrdde foam btown with butane woutd not provide an accurate measmement of
the mass Bow rate of the polyrner mett or the poIymerhutane solution. The rnass flow rate of
the solution c m then be determined by Equation 3, using the sepmteIy measured gas and
polymer m a s flow rates. Therefore, it is proposed that the mass flow rate of polymer melt
only be measured by substantially degassing the extruded polymerlbutane solution insrmd of
rneasuring the mass flow rate of the solution.
The principles of Axiomatic Design [167] were employed to design a special
degassing oven to remove al1 traces of the blowing agent from the polymerigas solution (see
Figure 4.2). The main design considerations were: (i) to devoid the polymerlgas solution
sample of gas and (ii) to determine the rnass of the sample. To achieve these objectives, a
heating mechanism and a weighing device were chosen, respectively. The main requiremenr
for the heating mechanism was to achieve and maintain high temperatures above 200°C. To
satisfy the requirement, a physical enclosure with appropriate insulation was designed ta
provide a finite space for heating, a heat source was chosen with the appropriate range, a
thermocouple and temperature controller were used to maintain the set temperature, and air
circulation was provided to maintain temperature uniformity. It was known that any
weighing device would be adversely affected by very high temperatures, and would result in
incorrect mus readings, Consequently, physical separation of these rwo functions was
desirable. A balance with the desircd resolution and capacity was chosen. The balance couId
take measurements in one of two modes: above or below the balance. It was determinrd that
physical separation would best be achieved by positioning the balance above the heating
mechanism. The overail design was then constructed to specifications, and tested for
functionality.
4.2.1 Experimentai Equipment and Procedure
4.2.1.1 Experimental Setup
The PVT measurement system for the propylenehutane solution consisted of a
tandem foaming extrusion line described in chapter 3 equipped with a variable resistance
valve, a filament die (see Figure 4.1), and a degassing system (see Figure 4.2)- An explicit
description of this setup can be found in the ceferences [60,6 11.
An integral part of this study involved using a degassing system to achieve maximum
gas loss from the foam sampIe. As shown in Figure 4.2, the degassing oven was divided into
two separate sections. The lower section was used for heating and degassing the propylene
foarns and consisted of a cenrnic heater (Omega: CRFP-126/120A). tnsulation was provided
for the walls in the form of glas wool fibers. The wall directly opposite the ceramic heater
was hinged to the frame to allow access to the heated chamber, and a circulation h n was
inserted through one wall. The upper portion of the oven was left exposed to the
environment with a simple platform on the top. A high-resolution balance (Scientech:
1 1 144- 12) equipped with beIow balance weighing function was mounted on this platform. A
weighing assembly was constructed of a small diarneter aluminum rod and thin duminum
plate. The rod section of this assembly was attached to the bottom of the scale to enable
below-balance weighing. The plate section was positioned in the oven to support the sample
during the degassing process. Finally, a Type J thennocouple was positioned in the lower
charnber and the temperature controlled using a temperature controller (Fuji: P m - T A Y 1-
4w-
4.2.1.2 Experimenîai Materiais
A linear and a branched polypropylene material with MFRs (ISO 1133, 230 OC/2.16
kg) of I l dg/min and 2.3 dg/min, respectively, were chosen as the poIymer materials for the
experimentation in the presect study. Both polypropylene materials were supplied by
Borealis AG Austria. The blowing agent used in this study was n-butane supplied by
Matheson Gas Company.
4.2.1.3 Experimental Procedure
The band heaters located dong the tandem foaming extmsion line equipped with a
gas injection system were turned on for approximately half an hour before running the
system to ensure proper plristication of the propylene materials that were charged in the
extruder through a hopper. Initially, the speed of the calibrated gear pump was set and the
system was started. While the system ran, butane was injected from the gas pump, and the
ga was mixed with and dissolved in the polymer melt in the first extruder. Then, the molten
polymerlgas mixture advanced through the second extruder, the heat exchanger, the g a r
pump, and exited the system through the filament die. Before the measurements, the system
was kept running for 20 minutes or more to reach a steady-state condition. The weight of
butane with respect to the weight of the polymer was maintained constant throughout each
experiment to ensure consistency. This required repeated measurements of the mass flow
rate of the polyrner as well as the mass flow rate of butane being injected by the gas pump.
The temperature of the filament die was maintained to be 220 O C throughout the experiments.
Since measurements had to be taken at a thermodynamically defined pressure and
remperature, control of these two parameters was essential to the measurement procedure.
The pressure upsueam of the gear pump was controlled by varying the rotationd speed of the
second extruder, while the downstream pressure was controlled using a variable resistance
valve 1601. The temperature of the rnelt was controlled in the second extruder using band
heaters placed around the barrel. The temperature was then Further homogenized unifonnly
as it passed through the static mixer, while sirnultaneously being cooled.
At the instant when the inlet and the outlet pressures of the gear pump were equalized
and no fluctuations in the pressure readings were observed, extruded polymerlgas solution
samples were collected for 1 minute at the die exit. Three separate sets of samples were
collected. After collection, the samples were placed in the degassing oven CO Facilitate
degassing of the sample. After 8 minutes (which represents the degassing time described in
the next section) the mas of each sample was recorded and the volumetric and mass flow
rates of the polymerigas solution determined. Finally, the specific volume of the
propylenelbutane solution was calculated based on Equation 1.
4.2.1.4 Calibration of the Degassing Oven
in order to reduce the error involved in merisuring the polymer tlow rate. the time
required to degas the polymerlgas solution sample was detennined frorn calibration
experiments. It should be noted that if the degassing tirne was too short, then the residud gas
would cause an error. The degassing oven was therefore caiibrated to detennine the
maximum time required for cornplete degassing of the samples. Extruded polymerigas
samples were collected at each experimentai condition and placed in the oven. The mass of
each sample was recorded at one-minute intervals. Of ail the cdibration experiments
perforrned, it was obsewed that the Iongest degassing times required occurred ar
expenmentd conditions of 170 O C , and 15% butane content. The result is depicted
graphically in Figure 4.3. It was observed that a npid decrease in the m m occurred beyond
a certain time. This is beiieved to be due to the degradation of the plastic materiai at the hi$
temperature. The threshold value was 8 minutes and this was selected as the maximum time
required for degassing during the experiments at ail conditions. It should be mentioned that a
significant amount of degassing of the propylene foam occurs immediately after exit from the
die, because of the small cross-section of the sample and the high temperature of the filament
die, which facilitates desorption of gas, however a foam structure is still evident. After
degassing in the oven, the foam structures disappeared, and only few bubbles of very small
sizes could be seen. Degassing was assumed complete, when the mass remained unchanged
for three consecutive readings.
42.2 Results and Discussion
4.2.2.1 Effect of dissolved butane on the specific volume
Figures 4.4 (a) and 4.5 (a) represent the effect of varying the amount of dissolved
butane in the solution on the specific volume Cor linear and bnnched propylene, respectively.
For both linear and branched propylene, the specific volume linearly increased as the amount
of injected butiuie increased, When the butane content was increased from O wt% to 15 wt%.
the specific volume increased by spproximately 02798 cm31g for linear propylene. This
increase was approximately constant over the temperature range 170°C to 210°C and the
pressure range 13.8 MPa to 37.6 MPa.
It is believed that, due to the high solubility of butane in propylene, al1 of the injected
butane was dissolved in the polymer, This reduces the solution density and consequently
increases the specific volume. For O wt% butane (pure propylene melt), the specific volume
increased by 0.0222 cm31g over the range of temperature from 170°C to 210°C. This was
found to be consistent with published resuIts 1601. With 15 wt% butane, the specific volume
increase was only 0.0045 cm3ig. This change seems to be reduced for temperature, however,
with respect to pressure, the sensitivity became higher. Over the temperature range 180°C to
220°C. Park et al. 1601 obtained a specific density increase of 0.050 cm31g when the CO2 was
increased from O wt% to 4 wt% for the PSICOl system. For the propylene/butane solution,
the specific volume increased approximately by 0.100 cm3tg when the butane content was
increased fom O wt% co 4 wt%. For the same wt% of gas, the swelling was therefore
observed to be higher for butane, This would indicate that butane has a higher plasticizing
effect than CO?.
4.2.2.2 Effect of branching on the specific volume
A significant difference between the specific volume behaviors o f linear and
branched propylene materials was observed (Figures 4.4 and 4.5). This cornparison reveaIs
that at rnost experimental conditions seiected, the specific volume was found to be higher for
the linex propylene material. This is of interest, because for polyethyiene, branched
materials are known EO have Iower density by occupying more volume. For the branched and
linear propylene materials of interest, the densities ac room temperature were almost the same
(0.91 @cm3). However, when the temperature was increased to the high range of 170°C to
190°C, the specific volume of the branched propylene was lower. When butane was
dissolved into the propylene matrix, the polymer swelled; it is believed that in branched
propylene, the long chain branching restricted the mobility of the molecular chahs. As the
butane dissohed in the propyIene matrix, the difference between the specific volumes of
linear and branched propylene resins increased, especially at high pressures. It was dso
noted that the branched propylene materiai showed a higher sensitivity with respect to the
change in pressure compared to the Iinear resin. The higher sensitivity of the branched resin
with respect to the pressure became more severe as the gas concentration increased. It is not
dear why the specific volume of branched resin shows a very high sensitivity with respect to
the pressure at hi& butane concentrations. A further study is required to clarZy this issue.
4.2.23 Effect of processing temperature and pressure on the specific volume
The graphs presented in Figures 4.4 (b) and 4.4 (c) itlustrate the effect of the
processing temperature and pressure respectively, on the specific volume of
propyIenehutane solutions for five different butane contents (Le., O%, 2%, 5%, IO%, and
15%) for linea. propylene. Sirnilar results for branched propylene are shown in Figures 4.5
(b) and 4.5 (c). These resuits reveal that at O wt% butane, the sensitivities of the specific
volume with respect to the temperature and pressure were similar to the results obtained by
others [57,58]. However, when butane was dissolved in the propylene melt, the specific
volume became more sensitive to a change in the pressure for both materials, whereas the
specific volume with respect to temperature was not changed much.
4.2.2.4 Error Analysis
Due to the dynamic nature of PVT measurements, a number of sources of error were
possible. These include the errors associated with the residual gas, and the controi and
measurement of temperature, pressure, and gas concentration. It is believed that the error
associated with the residual gas in the polymer matrix was negligible because of the newly
designed degassing system. Evidence in support of this assumption could be observed in the
degassed polymerlgas solution, as only very few gas bubbles of very small size could be seen
in the degassed propylene~butane solution. On the issue of temperature measurement and
control, examination of the results reveals that the average sensitivity of the specific volume
with respect to temperature was approxirnately 0.005 crn3/g0~. The error range associated
with the temperature control of our systern was detemined to be t 2°C [60,61]. This
corresponds to an error of t 0.0 10 cm31g in the specific volume, With regards to the pressure
control and measurement, our system error was determined to be approximately k 0.69 MPa.
From our results, the sensitivity of the specific volume with respect to pressure was
approximately 3.600~10-~ c r n 3 / g - ~ ~ a . The error range in the specific volume associated
with the sensitivity due to pressure is Cherelore 2.500~10-~ cm3/g. The average sensitivity of
the specific volume with respect to the butane content was found to be relatively high at
0.020 cm3/g-butane wt%. However, ctiere is no possible way of measuring the final butane
content in the extmded propylene foam CO determine our systern error with regards to the
butane content. Further research and development are therefore needed to quantify and
minirnize the error in the specific volume associated with the butane content. Since the
butane content of the foam is also inextricably Iinked to the butane flow rate. improvements
to the system must also be made in terms of accurate gas flow rate control.
4.2.3 Conclusions
In this study, the effect of dissolved butane on the PVT relationships of linear and
branched propylene materials in a molten state were investigated. Based on the experimental
results, the following conclusions can be drawn,
1. The specific volume of the propylenehutane solution increased significantly with an
increase in the percentage of butane injected in both branched and linear propylene. Both
the linear and branched propylene resins sweHed significantly as the butane permeated
into the resins.
2. At al1 experimental conditions selected, the specific volume was found to be higher for
the linear propylene than for the branched propylene.
3. When butane is dissolved in the propylene matrix, the sensitivity of the specific voIurne
with respect to pressure increased with the butane content for both the linear and
branched resins, whereas there were no significant changes in the sensitivity with respect
to the temperature. EspeciaIly, the branched resin exhibited a very high sensitivity with
respect to the pressure at a high concentration of butane.
4.3 Measurements of Thermal Behavior of Polypropylene Materials
The thermal behaviors of linear and branched poiypropylene with foaming additives
were investigated using a high-pressure differential scanning calorimeter (DSC).
Specifically, the effects of material branching, dispersed additives, and dissolved blowing
agent on the crystallization temperature of polypropylene resins were elucidated. The effect
of dissolved blowing agents was determined using the high pressure DSC cell with carbon
dioxide and nitrogen. The effect of hydraulic pressure was identified by performing DSC
study employing an inert gas such as helium, which has a very low solubiiity in the polymer
matrix. Foaming additives such as tdc and GMS as welI as processing parameters such as the
cooling rate dso played major roies during the crystallization process [175],
4.3.1 Experimental Equipment and Procedure
4.3.1.1 Experimental Setup
The crystallization experiments were perfonned w ith a DSC (TA Instruments, DSC
29 IO), A regular DSC ce11 and a high-pressure DSC cell were used in the experiments.
4.3.1.2 Experimental Materials
A linear and a branched propylene material were chosen as the polymer materials.
They are the same materials used in section 4.2. Talc (A7 with top-cut of 7 microns,
Naintsch) and glycerol monostearate (GMS, Pationic 909, PATCO Polymer Additives) were
used as foaming additives in our investigation. HeIiurn (BOC Gas 99.9 % purity), nitrogen
(BOC Gas 99.9 % purity) and carbon dioxide grises (BOC Gas, 99.5 % purity) were used in
high-pressure experiments.
4.3.1.3 Experimental Procedure
The regular DSC ce11 was used for investigating the effects of branching, additives
and cooling rates on the crystaliization behavior of polypropylene materials. The high-
pressure DSC cell was used for investigating the effects of dissolved gases on the
crystallization behavior of polypropylene materials. The calibration of both DSC cells was
done using indium. Samples for DSC experiments (typicai weight 3-4 mg) were taken from
the extrudate in a form of a very thin disk (typicai thickness 150-200 Pm) [176]. For the
nonisothermal experiments, the smples were heated up to 220 OC and kept at this condition
for 30 minutes to ense the thermal history [89], and then the samples were cooled down at
10 'Clmin (if not specified) to 50 O C . Next, the samples were heated at 10 'Clmin up to 200
OC. During the cooling and heating processes, the crystallization and melting patterns were
recorded- In the case of high pressure DSC ceII, the sarnples were pressurized to 1.37, 2.75,
4-13 and 5.5 1 MPa using helium or nitrogen or carbon dioxide.
4.3.1.4 Design of cooling system for the high pressure DSC ceIl
In order to investigate the effects of dissolved gas on the crystallization behavior of a
poiyrner, a high-pressure DSC ceIl (TA Instruments, DSC 2910) was used for this
experiment. Due to the poor cooling capability in the pressure DSC cell. only the heating
mode has been mainiy used in the high-pressure experiments.
A cooling system was developed as shown in Figure 4.6 to provide the existing high-
pressure ce11 with a cooling capability. The cooling systern enabled the measurernents of the
crystallization behavior of a polymer melt under high pressure by mounting a cooling coi1
connected to a liquid nitrogen dispensing unit which supplies liquid-nitrogen at a controlled
rate. Pnor to the cooling coi1 modification, the high pressure DSC could only attain non-
uniform cooling rates o l 10 "Clmin. Critical experiments were conducted to verify the
pressurizing and cooling functions of the modified pressure cell. The pressure was raised
inside the modified high-pressure ce11 from atmospheric pressure to 3.51 MPa, and at this
pressure a uniform cooling rate up to 30 "Clmin was successfully achieved. Also, the
crystallization behavior of polypropylene materials measured with the high-pressure DSC
ce11 was compared to that with a regular ce11 operating at atmospheric pressure. The
crystallization thermograrns of polypropylene materiais were determined at atmospheric
pressure using both the DSC cells and were found to be nearly identical at the same cooling
rates.
Finally, using this design, the measurements of crystallyzation kinetics under high
pressure was successfully carried out using the modified high-pressure DSC ceIl.
4.3.2 Regular DSC Celf Results
4.3.2.1 Effect of Branching
Figure 4.7 shows the cooling sections of the DSC thermograrns obtained for linear
and branched propyle - a materials without any additives at a cooling rate of IO 'Chin. It was
observed that branching of propylene chahs significantIy promoted the crystallization
kinetics of polypropyiene resins by increasing the crystaliization temperature about 20 OC.
This result supports the findings of previous studies [177,178], Figure 4.7 also shows that the
peak of linear polypropyiene materiais has a shouIder in the peak (or double peak) caused by
two-stage crystallization [66,179,180]. It can be understood that the interactive motion of the
polypropylene matrix layer at the particle surface changes the crystallization speed of
polypropylene matrix [66].
4.3.2.2 Effects of Foaming Additives
The thermal behaviors of linear and branched propylene resins with foaming
additives such as talc and GMS were also investigated as a part of this study. The
experiments were conducted at a cooling rate of 10 "Clmin. The concentrations of talc and
GMS were changed from O to 1.6 wt% and O to 1 .O wt%, respectively.
Figure 4.8 (a) shows the effect of talc amount on the crystallization temperature of
polypropylene materials. The effect of talc was more dominant in Iinetu polypropylene
material than in the branched one. After showing a sharp increase in the crystallization
tempenture as the talc concentration increased from O to 0.2 wt%, the crystallization
tempenture did not change much above 0.2 wt%. Figure 4.8 (a) shows that the crystallizacion
tempenture of branched polypropylene material is only 5-10 OC higher than that of the linear
material in the actual foam processing with the talc content ranging from 0.8-1.6 wt9.
However, without the addition of talc particles, the crystallization ternperiture of brrinched
polypropylene material is 10-20 "C higher than that of linear one. The degrees of crystallinity
of linear and branched propylene resins were also measured and the results are shown in Fig.
4.8 (b). It was observed that the degrees of crystallinity of branched and linear propylene
materiais increased moderately as the talc concentration increased.
Very sirnilar results were obtained for the case of GMS (Figs- 4.9 (a) and 4-9 (b)),
The GMS, used as an aging modifier, also increased the crystalIization temperatures and the
degrees of crystallinity of Iinear and branched propylene resins. However, because of its
strong tendency of migrate to the surface of the extrudate 118 11, the uniformity of GMS in
the polypropylene matrices could not be ensured. It is believed that the variations shown in
Fig. 4.9 (a) are due to the non-uniform distribution of GMS particles. The DSC samples for
this study were cut perpendicularly to the flow direction from a thin filament extrudate (about
4 mm in diameter) to minimize the effect of non-uniformity of GMS particles in the radiai
direction of filament extrudate.
4.3.2.3 Effect of Cooling Rate
The effect of cooling rate on the crystallization kinetics of linear and branched
propylene resins was also investigated. The cooling rate was varied from 0.1 Wmin to 50
"Clmin. Figure 4-10 shows the dependence of crystallization temperature on the cooling rate
for Linear and branched propylene materials. It was observed that the crystallization
temperatures of polypropylene resins were very sensitive to the change of cooling rate: the
crystallization temperatures decreased by 25-30 OC as the cooling rate increased from 0.1
"C/min to 50 'Clmin.
4.3.3 High-Pressure DSC Cell Results
Figures 4.1 1 and 4.12 show the effects of pressure on the crystallization behriviors of
linear polypropylene materials (Fig. 4.1 I ) and branched polypropylene materials (Fig. 4.12)
with various grises. The dependence of crystallization temperature on the pressure was quite
different for different gases. It was obvious that the gas at high pressure significantly affected
the crystallization kinetics of polypropylene materials. As a gas permeates into the pulymer
rnatrix under high pressure, the dissolved gas may change the rate of polymeric segmental
motions such as rearrangement into crystals. Furthemore, since the amount of gas dissolved
in the polymer increases with an increase in pressure [106], the magnitude of the change in
the plasticization and crystallization will be more pronounced at a higher pressure.
On the other hand, the crystallization kinetics of the polymer melt will aIso be
affected by the hydnulic pressure externally applied by the gas. When the solubility of gas in
the polymer is very low or negligible, the kinetics of crystallization will be governed by the
hydraulic pressure. But when the solubility of the gas in the polyrner is considerably high,
both the hydraulic pressure and dissolved gas will play individual roles during the
crystailization process in high-pressure experiments. The effect of dissolved gas on the
crystailization kinetics cm then be extracted by subtracting the effect of the hydraulic
pressure from the overall crystallization behavior of the material under high pressure.
4.3.3.1 Effect of Hydraulic Pressure
The effect of hydraulic pressure on the crystallization behaviors of polypropylene
materiais was determined from the high-pressure expeciments with helium (He). Since the
solubility of He in a polymer is very low, typically one order of magnitude lower than that of
NI and two orders of magnitude lower than that of CO? [106], the effect of the dissolved He
in a polymer melt would be negligible. Therefore, any change of the crystallization kinetics
under an elevated pressure of He can be considered as the effect of hydraulic pressure.
Figure 4.13 shows the onset and peak crystallization temperatures as a function of the
He pressure. It was noted that by increasing the pressure in the DSC cell, thc onset and peak
crystallization temperature increased for both linex and branched propylene resins: the onset
and crystallization peak temperatures for branched polypropylene material increased by
about 6 O C when the pressure was increased from the atmospheric pressure to 5.5 MPa, but
for linear polypropylene material, the temperatures increased by only about 3 OC for the same
amount of pressure change. As the hydraulic pressure in the DSC cell increases, the rnobility
of the polymer matrix molecules decreases, and hence accelerates the crystallization process,
resulting in a higher crystallization temperature. Figure 4.13 also shows that the increase of
the crystallization temperature for the branched polypropylene material was even more
pronounced than that of the linear material.
The crystaIlization temperature shown in Fig. 4.13 was not changed much above 2.75
iWa and no further effect of hydraulic pressure was observed at higher pressures. it was not
clear if the plateau region observed was due to the dissolved He in the polypropylene
materiais at elevated pressures, or if it reflected the actual effect of hydraulic pressure on the
crystallization behavior. Once the solubility of He in polypropylene resins is measured, this
issue will be clarified.
4.3.3.2 Effect of Dissolved N2
The effects of dissolved N2 on the crystallization behaviors of Iinear and branched
propylene materials were extncted by subtracting the hydraulic pressure effect (Fig. 4-13]
from the overail thermograrns (Figs. 4.1 1 and 4.12). Figure 4.14 shows the resuIting onset
and peak crystallization temperatures as a function of the Nz pressure. N2 has a relatively
higher solubility than He 11061, and therefore, the effect of dissolved gas with N? on the
crystallization kinetics of polypropylene materials will be more pronounced when compared
to the case of He.
It was observed that the onset and peak crystallization temperatures did not change
much at pressures below 1.4 MPa. Since the solubility of & in the polymer is very low at a
low pressure, the overall crystallization behavior in the low pressure range was mainly
governed by the hydraulic pressure effect as shown in Figs. 4.1 1 and 4.12. The effect of
dissolved N2 on the crystallization was more pronounced at a higher pressure (above 1.4
MPa), the onset and peak crystallization temperatures decreased moderately up to 5.5 1 MPa,
It should be noted that these compensated results shown in Fig. 4.14 may have some
errors due to the dissolved He in the polypropylene rnatrix at high pressures, causing an
incorrect hydraulic pressure effect. However, it is believed that the error range would be
srnail.
4.3.3.3 Effect of Dissolved COt
As in the case of N:, the effects of dissolved CO2 on the crystallization behaviors of
linear and branched propylene materials were extracted by subtracting the hydraulic pressure
effect (Fig. 4.13) from the overall thermograrns (Figs. 4.1 1 and 4-12). Figure 4. t 5 shows the
onset and peak crystallization ternperatures for the linear and branched polypropylene
materials as a function of the CO2 pressure. Tt was observed that the dissolved CO2
suppressed crystallization of the polypropylene materials significrtntly. The decreased
crystallization temperatures due to the dissolved CO2 was more pronounced in the
thennogram of branched polypropylene material cornpared co the linew one.
At high pressure (5.51 MPa), the increased crystallization temperature due to the
branched structure was compensated for by the dissolved CO? and the crystallization
temperature of branched polypropylene material becarne almost the same as that of the Iinear
one. This indicates that the effect of branching on the crystallization kinetics may be
neglected at elevated CO2 pressures above 5.5 1 MPa. However, because of the limitations of
our high-pressure equipment, experiments at pressure higher than 5.51 MPa coutd nor be
performed. On the other hand, the crystailization temperature of Linear Pt decreased Iineariy
as the pressure increased.
Tt is believed that the Iarger magnitude of change in che crystailization of
polypropyIene resins under a high pressure with CO: compared ta the case of NI_ wris due to
the solubility difference of both gases. Because of the higher solubility of CO? compared to
N2 [182], the effect of dissolved CO, was more pronounced even at a low pressure. A large
amount of dissoIved CO2 increases the f ee volume of the polymer [183], and enhances the
mobiIity of polyrner chah The resulting crystallization temperature of polypropylene
materials is therefore lowcnd.
4.3.4 Conclusions
A series of experiments were conducted to investigate the effects of materia1
branching, foarning additives, cooling rate, hydraul ic pressure, and dissolved gas on the
crystdlization behaviors of poiypropylene resins. Helium was empioyed to estimate the
hydraulic pressure effect on the crystallization behavior for the high-pressure experiments
with N2 and CO2 blowing agents. The f o a i n g additives considered in this study included
taIc ruid GMS. The experiments conducted in this study tead to the following conc1usions:
1. Branching in the poiypropylene rnatrix caused a significant increase in the
crystallization temperature.
2. The foaming additives such as talc and GMS increased the crystalIization temperature
of polypropylene materiais.
3. The crystallization temperature was a sensitive function of the cooling rate and it
decreased as the cooling rate increased.
4. Crystailization of polypropylene materials was enhanced as the hydraulic pressure
increased-
5. But the dissolved I$ and COz lowered the crystailization temperatures of
polypropylene resins. In particular, high-pressure COz decreased the crystailization
temperature significantiy because of the high sohbility.
4.4 Measurements of the Onset of Melt Fracture of Polypropytene Materials
The melt fracture behaviors of Iinear and branched polypropylene resins with
foaming additives were investigated. The effects of branching, processing temperature,
additives, and blowing agent on the surface melt fracture of polypropylene materiais were
thoroughly studied. A CCD carnera was installed at the die exit to precisely observe the onset
of surface melt fracture of extruded foams. The critical wall shear stress was determined for
various linear and branched polypropylene resins using a capillary die [184].
An experimental setup is designed to investigate the surface melt fracture behavior of
polypropylene materials melts and polypropylene/butanc solutions under various conditions,
Figure 4.16. This system is based on the tandem extrusion systern described in Chapter 3.
Due to the difficulties involved in determining the onset of melt fracture for polymer
foam, a visual approach is employed. A CCD camera is instalted at the die exit and the foam
extrudate is carefully monitored. In the case of foaming extrusion, the skin of extruded foam
is stretched as the expansion occurs. As a consequence, the foarn skin becornes shiny as
alortg as the expanded foarn does not contract due to gas loss [20]. Even if melt fracture
occurs and sharkskin developed on the extrudate surface, the foam expansion of extrudate
causes the foam skin to be stretched, and thereby, the trace of the sharkskin gets easily
removed, Therefore, it is vrry difficult to detect the onset of melt fracture by sirnply checking
the surface of the fully expanded foam. However, it is believed chat the onset of surface melt
fracture cm be observed by monitoring the early stage of extrudate using a CCD camera
before it gets expanded.
A capiIlary die of length L and radius R is used to calculate the critical shear stress at
which meIt fracture occurs [185]. A pressure transducer is mounted in the die More the
capiIIary section. in order to minimize the entrance effect, the UD ratio was chosen to be 35.
and the wall shear stress is calculated in terms of the die pressure, Pd, and the die geometry
[ISS]:
The corresponding wall shear rate, yUPP, is determined by 11851:
The polyrner flow rate is controlled by the rotational speed of the gear pump regardless of
the temperature and pressure fluctuations in the barrels. Appendix 1 shows a detailed analysis
of the Flow in capillary die.
4.4.1 Experimental Equipment and Procedure
4.4.1.1 Experimental Setup
Based on the above design, an experimental setup is constructed to study the effects
of branching, processing ternpenture, additives, and blowing agent on the surface rnelt
fracture behaviors of polypropylene materials melts and polypropylenehutane soIutions. The
setup consists of the tandem extmsion system described in Chapter 3, a ccipillary die of O. I O
cm diameter and 3.55 cm in length, and a cooling sleeve for the precise control of dit:
temperature. A CCD camera (Pulnix) is mounced at the exit of the die and connected to a
computer processor in order to accurately visudize and monitor the onset of rnelt fracture.
4.4.1.2 Experimental Materials
The materials used in this scudy were two linear standard polypropylene resins and
three high-rnelt-strength (HMS) branched polypropylene resins supplied by Borealis AG.
They are denoted in this study as Linear PI, Linear P2, Branched Pl, Branched P2, and
Branched P3, respectively. The materials properties, inciuding MFRs (ISO 1133,230 'Cf?. 16
kg), molecular weights and numbers ( M W , Mn), and the degrees of long chah branching per
1000 carbon atoms (LCB), are summarized in Table 4.1. The foaming additives used in this
study were talc and glycerol mono stearate (GMS) as the cell-nucleating agent and the aging
modifier. respectively [18 11. The blowing agent used in the experiments was n-butane, C.P.
(Matheson, 99.0%).
4.4.1.3 Experimental Procedure
The polypropylene resins were processed in the extrusion setup and the onset of me[t
fracture was investigated by obsewing the die pressure, the surface qua1it.y of the extrudate,
and the image of the extrudate €rom the CCD canera.
Firstly, experiments were conducted without gas injection for investigating the effects
of branching, melt temperature, and foarning additives on the melt fracture. Brrinched
materials with various degree of long chah branching were used to invescigate the effect of
brmching on the melt fracture. For investigating the effect of processing temperature, the
melt temperature was precisely controlled using the tandem extrusion system. Vanous
amounts of talc and GMS were added to the polypropylene materials to investigate the
effects of foaming additives on the surface melt fracture of the extrudate, The onset of
surface melt fracture was determined by directly obsewing the surface quality of the
extrudate and also by analyzing the CCD image of the extmdate. As the die pressure
increased by increasing the gear pump speed, the occurrence of melt fracture was detected.
When melt fracture occurred, the critical shear stress was calculated by reading the
corresponding die pressure.
Secondly, experiments were conducted with gas injection to study the effect of
dissolved gas on the melt fracture. A metered amount of blowing agent was injected and
dissolved into the polypropylene melt to form a single-phase polypropylene/butane solution.
The fomed single-phase polypropylene/gas solution entered the die and foaming was
allowed. The onset of melt fracture on the foamed surface was carefully observed using the
CCD camera as the die pressure was increased with a higher speed of gear pump.
4.4.1.4 Calibration of the CCD camera image for detecting the onset of surface melt
fracture
In order to precisely deterrnine the effect of dissolved gas on the onset of s~irface melt
fracture of polypropylene materials, the image of the extrudate captured with a CCD camera
For the early stage of foaming was carefully analyzed. It was obsewed that for a pure
polymer melt without any dissolved gas. and without melt fracture the extrudate was steadily
exiting the die. In this case, the extrudate shape was uniform and smooth. However, at the
onset of surface melt fracture, the extrudate corning out of the die started to oscillate
verticaiiy in a direction perpendicular to the flow. It was also confirmed that surface
irregularities started to appear on the surface of the extrudate at this moment. Since the
amplitude of oscillation was very smdl at the onset of surface melt fracture, the oscillation of
extrudate could be detected only by the magnified image of the extrudate from the CCD
camen.
On the other hand, when gas was dissolved in the polyrner and the expansion of
extrudate occurred due to foaming, the appearance of surface melt fracture on the foam
surface was not synchronized with the moment when the foam extrudate started to oscillate.
in other words, even when the minute-scale oscillation of the extrudate was detected from the
captured CCD image, there was no visible change on the extrudate surface and the degree of
straightness of foam extrudate. It was believed that surface melt fracture actually occurred at
the moment of oscillatory motion of the extrudate. However, the trace of melt fracture on the
foam surface must have been removed because of the stretching of foam skin during
expansion. The foam skin at this moment was typically very shiny. By contrat, when the
degree of oscillation of the extrudate became Iarger so as to be visible even to the naked eye,
the foam extrudate became wavy. However, the foam surface was still shiny and smooth.
In conclusion, the onset of surface melt fracture could be detected effectively using a
CCD camera by observing the image of the extrudate shape as it starts to oscillate for both
pure polyrner and Foam.
4.4.2 Results and Discussion
4.4.2.1 Effect of branching on the critical shear stress
The effect of branching on the surface melt fracture behavior of polypropylene resins
is shown in Figure 4.17, without the use of additives. It was observed that the critical shear
stress linearly decreased with an increase in the degree of long chah branching, in other
words, the surface melt fracture of polypropylene resin was promoted by the degree of
branching, According to the experimental results, the critical shear stress decreased
approximately by 0.0175 MPa with an increase of 0.1 nJ1000c in long chain branching The
decrease of criticaI shear stress may be attributed to the increase of melt elasticity as the
degree of long chah branching increased. An increase in melt elasticity can increase the slip
at the die wall, and hence promote surface melt fracture [6 11.
4.4.2.2 Effect of processing temperature on the critical shear stress
The effect of processing temperature on the surface melt fracture behavior of
polypropylene resins is shown in Figure 4.18 for al1 the linear and branched materiais, In this
experiment, the processing temperature was varied from 180 OC to 210 OC. It was observed
that the critical shear stresses of linear and branched polypropylene resins were aImost
insensitive to the processing temperature in the range of 180 OC to 210 OC.
The critical shear rate at the onset of surface melt fracture was also calculated as a
function of the processing temperature using Equation 4.6. Figure 4.19 shows that the
apparent shear rate at the onset of surface melt fracture linearly increased as the temperature
increased from 180 OC to 210 OC for al1 the polypropylene resins. The shear rates behavior
observed in this study conform to the behavior obsewed in other studies [95,96].
4.4.2.3 Effects of foaming additives on the critical shear stress
The effects of the dispersed foaming additives, Le., talc and GMS, on the critical
shear stress of polypropylene resins were also investigated. All the experiments were
conducted at 190 OC. Figures 4.20 and 4.2 1 depict the critical shear stress for Linear P 1 and
Brrinched P 1 materials as a function of the talc and GMS contents. respectively. For both talc
and GMS cases, it was observed that the critical shear stress of Branched PI increased
sharply as the amount of dispersed additive increased from O to 0.4 wt%. A further increse
in the arnaunt of dispersed additives only increased the critical shear stress slightly. By
contrast, the critical shear stress of Linear PI did not change much with the foaming
additives. For both talc and GMS, the critical shear stress for Linear Pl increased sli~htly
(less than 0.005 MPa) as the concentration of additives increased from O to 2.4 wt%.
It is interesting to note that the well-known lubricating effect of GMS was not
distinguished in this study [186]. On the other hand, the lubricating effect of talc, as a ceII
nucieating agent, for the branched polypropylene materials was also surprizing whereris the
Iubricating effect of boron nitride, which is another well-known ce11 nucleating agent in f om
processing, has been reported before [174]. It is speculated that there may be some
rdationship between the cell nucleating ability of the additive and its lubricating effect.
Further study needs to be conducted to c1arifL this issue.
4.4.2.4 Effect of blowing agent on the critical shear stress
The effect of the dissolved blowing agent on the criticai shear stress of polypropylene
resins is shown in Figure 4.22. It was observed that the critical shear stress increased
significantly as the amount of dissolved butane increased: for both linear and branched
polypropylene resins, the criticai shear stress increased by 0.025 MPa as the butane content
increased from O to 20 wt%. These results rnean that the surface rnelt fracture is significantly
suppressed by the dissolved butane. The decreased criticai shear stress of Branched PL by
branching was alrnost recovered by the dissolved 20 wt% butane. On the other hrind, the
critical shear stress of Linear Pl was also significantly increased by the dissolved butane.
4.4.3 Conclusions
Experirnental studies were camed out to investigate the effects of branching,
processing temperature, foaming additives, and blowing agent on the critical shear stresses of
linear and branched polypropylene resins. An experimental setup was designed to elucidate
the surface rnelt fracture behaviors of polypropylene melts and polypropylendbutnne
solutions under various conditions. Efforts were made to accurately control the processing
temperature, to well disperse the foaming additives in the polypropylene melt, and to fom a
single-phase polypropyleneibutane solution. A CCD camera wris installed at the die exit to
precisely monitor and analyze the onset of surface melt fracture of the extrudate at the early
stage of foam processing. From the expenrnents conducted in this study the following
conclusions cm be drawn:
1. An on-line technique for detecting the onset of surface rnelt fracture for extmded foam
has been developed by visualization of extrudate using a CCD camera.
2. The long-chain branching of polypropylene rnaterials significantly decreased the critical
shear stress of the resins.
3. The critical shear stress was aimost insensitive to the die temperature; however, the die
temperature significantly affected the criticai shear rate at the onset of surface mett
fracture.
4. The foaming additives of talc and GMS increased the critical shear stress of branched
polypropylene materiats. However, they did not affect the critical shear stress of Iinear
polypropylene rnaterials much.
5. The dissolved butane significantly increased the critical shear stresses of linear and
branched polypropylene resins.
4.5 Summary
Based on the results shown in this chapter, the effects of processing conditions and
foarning additives on the amount of swelling of polypropylene melts, the c~stallization
point, and the onset of melt fracture were identified. The specific volume was highly
dependent on the processing temperatures and pressures, the degree of branching, and the
dissolved butane. The crystallization temperature of polypropylene materials was a sensitive
function of branching, foaming additives, dissolved b1owing agent, and the cooling rates. The
onset of melt fracture was affected by the degree of bnnching, the addition of foaming
additives, and the dissolved butane. These parameters are the bais for development of the
strategies for the production of low-density, fine-celled polypropylene foams, and for
identifying the fundamental mechanisms that governs the volume expansion ratio of
polypropylene foams.
Table 4.1. Materials properties of Iinear and bcanched polypropylene materials
- Material
Branched P 1
Branched PZ
Branched P3
Linear P 1
Linear P2 I
k[FR
(gllornin)
2.3
4.8
3.1
I I
2.8 I
LCB degree
(n/lûûûc)
0.2 1
0.17
O. 13
O
O I
MW
(Wmol)
418
416
327
360
500 I
Mn
(KgImol)
3 1.7
48
46.5
70
97 I
Figure 4-1: Schematic of the tandem extrusion Iine that provided the extruded samples for measuring the PVT propenies of polypropylene/butane solutions
Figure 4.2: Schematic of the degassing oven 1. High Resolution Balance 2. Heater 3. Below Balance Weighing 4. Sample Holder 5. Circulating Fan
8 5 I
1 2 3 4 5 6 7 8 9 t O
the (mh)
Figure 4.3: Calibntion of the degassing oven
Butane Content (%)
(a) Specific volume vs butane content for Iinear polypropylene
12 14 16 18 20 22 24 26 28 Pressure (MPa)
(b) Specific volume vs pressure for linear polypropylene
Figure 4.4: Changes of the PVT data of Iinear propylene
Temperature (C)
(c) Specific volume vs temperature for linear polypropylene
Figure 4.4: Changes of the PVT data of linear propylene (cont'd)
1.20 C 1
O 2 4 6 8 10 12 14 16 Butane Content (%)
(a) Specific volume vs butane content for branched polypropylene
J L.
12 14 16 18 20 22 24 26 28 Pressure (MPa)
(6) Specific volume vs pressure for branched poIypropylene
Figure 4.5: Changes OF the PVT data of branched polypropylene
1.20 . 170 190 210
Temperature (C)
(c) Specific voIume vs temperature for branched poiypropylene
Figure 4.5: Changes of the PVT data of branched polypropylene (cont'd)
I 1
4.6 Design of Cooling Capability for High-Pressure DSC Cell
Crystallization hmperature (C)
Figure 4.7. DSC thennograms of Iinear and branched polypropylene resins
+Eranched P l (onst) +ûranched Pl (peak) +Linear P l (onal) +Linear P l (peak)
100 ! 1
Concentration of Talc (%)
Figure 4.8 (a) Effect of talc on the crystailization behaviors of linear and branched polypropylene resins
0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0
Concentration of Talc (%)
Figure 4.8 (b) Effect of talc on the degrees of crystdlinity of linear and branched polypropylene resins
-6-Branched P l (onsht) . -a-Branched P l (ptak)
+ Linear P l (peak) , 100 .
0.0 0.2 0.4 0.6 0.8 1 .O 1.2
Concentration of GMS (%)
Figure 4.9 (a) Effect of GMS on the crystallization behaviors of linear and branched polypropylene resins
Concentration of GMS (%)
Figure 4.9 (b). Effect of GMS on the degrees of crystdlinity of Iinear and branched poiypropylene resins
Cooling Rate (Clmin)
Figure 4.10. Effect olcooling rate on the crystdlization betiaviors of linear and branched polypropylene resins
i oo J I
O 1 2 3 4 5 6
Pressure (M Pa)
Figure 4.1 1. Effect of pressure on the crystallization behavior of Linear polypropylene
Pressure (MPa)
Figure 4-12. Effect of pressure on the crystallization behavior of Branched poiypropylene
--
n n w O w
1 - -e-Branched P l (onat) +Branched Pl (peak) :
I i -e- Linear P l (onmt) 1 +Linear P l (peak)
100 4 O 1 2 3 4 5 6
Pressure (MPa)
Figure 4.13. Effect of hydraulic pressure on the crystallization behaviors of linear and branched potypropylene resins
r - .. -- -. - . - - - . - --- - - - - - ---- / -€+Branched Pl (onat) 1
135 i +Oranchad P l (peak) - --I +Linear PI tonatl
Pressure (MPa)
Figure 4. L4. Effect of dissoived N2 on the crystalIization behaviors of linear and branched polypropylene resins
O 1 2 3 4 5 6
Pressure (MPa)
Figure 4-15. Effect of dissolved CO2 on the crystallization behaviors of linear and branched polypropyiene resins
Figure 4.16. Experimental setup for measuring the onset of melt fracture
Long Chain Branching (rdl000c)
Figure 4-17. The effect of long-chain branching on the critical shear stress
0.05 1 175 180 185 190 195 200 205 210 215
Processing Temperature, C
Figure 4.18. The effect of processing temperature on the critical shear stress
Pracessing Temperature, C
Figure 4.19. The effect of processing temperature on the ctiticd shear rate
0.0 0.5 1 .O 1.5 2.0 25 3.0
Talc Content, %
Figure 4.20. The effect of talc on the critical wall shear stress
GMS Content, %
Figure 4.21. The effect of GMS on the cnticd stiear stress
Butane Content, %
Figure 4.22. The effect of dissolved butane on the critical shear stress
Chapter 5
Production of Low-Density, Fine-Celled Polypropylene Foams
5.1 Introduction
As discussed earlier, during the extrusion of polypropylene foams, the biowing
agent that promotes foam expansion may escape through the exterior skin of the foam
resulting in contraction of the foam [20,21,120]. Since the expansion of foam relies on
the presence of blowing agent within the foam, it is desirable to devise a means for
preventing gas loss in order to achieve a large volume expansion ratio.
There are few studies that have investigated the foaming behavior of
polypropylene, as it possesses a weak melt strength, which increases the difficulties of
foaming when compared to other plastics. When the melt strength is too weak, the ce11
walls separating the bubbles may not have enough strength to bear the extensional force
of the expanding bubbles and may rupture very easily during foaming. As a resuIt,
foamed polypropylene products usually have a high open-ce11 content and thus rire
unsatisfactory for many applications.
This chapter studies the production of low-density, the-celled polypropylene
foams on the tandem extrusion system and experimental results are presented that verify
the feasibility of the proposed ideas. It presents an effective strategy for promoting low-
density, fine-celled polypropylene foams. The effects of processing parameters such as
the temperature, the materials parameters such as the blowing, nucleating agents, and
long chah branching, and the die geometry on the production of low-density, fine-ceiled
polypropylene foams are investigated.
5.2 S trategy for Promoting Low-Density, Fine-Celled Polypropylene Foams
The basic approach for the promotion of a large volume expansion ratio with
poiypropylene materiais is fourfold [187]: 1) to use a branched material for preventing
cell coalescence; 2) to use a long-chah blowing agent with low diffusivity; 3) to Iower
the melt temperature for decreasing gas loss during expansion; and 4) to optimize the
processing conditions in the die for avoiding tao quick crystallization.
One of the most important strategies for promoting large expansion of
polypropylene foams is use of branched polypropylene rnaterials. Because of the weak
melt strength, cell codescence will be severe in the extnided foams with linear
polypropylene rnaterials [4]. When cells are coalesced. the cell population density and
cell size uniformity are not only deteriorated, but the volume expansion ratio is aIso
greatly sacrificed because of accelerated gas loss through opened ceII wdls. Therefore,
using branched polypropylene rnaterials is essenrial for Iow density as welI as fine-ce11
foam processing.
Another strategy that i s required to obtain low-density polypropylene forims is to
use a long chain blowing agent [l06]. Since the inert gris blowing agents have highet
volatility, and herefore higher diffusivity than the long-chah blowing risents. gaseous
bIowing agents can escape easily during expansion [20,21,i88]. Therefore, it is very
difficult to obtain a low-density f o m with a large expansion ratio of over forty-lold with
an inert gas. On the other hand, the long-chain blowing agents have Iow diffusivity
becriuse of their low volatility. This low diffusivity offers a tremendous advantage in
controlling cell'growth to achieve a very high expansion ratio of over forty-fold. With a
lower diffusivity, the cell growth rate is slower, and thereby gas escape can be more
easily blocked. As a consequence, a high volume expansion ratio can be easiIy achieved.
As the thickness of the ceII waIIs decreases in low-density foam production, the
cell-to-cell diffusion rate is increased, and therefore, the diffusion rate of gas escaping out
of the foam to the environment increases. Tt should be noted that the blowing agent which
has diffused into the nucleated celIs eventually tends to diffuse out to the atmosphere
because the compiete separation of two phases is therrnodynamicaily more favourable
[L72]. Gas escape through the thin walls will decrease the arnount of gas available for the
growth of cells. As a result, if the celis do not ireeze, they tend to collapse causing foam
contraction. in our foam extrusion process of potypropylene materiais with butane or
pentane, this volume contraction due to gas loss dter an initial voIume expansion has
been observed as with our previous foam expansion studies of other materials
[20,21,188]. As a consequence, the final product had a high foam density. In order to
produce low-density foams. gas diffusion through the skin of the foam extrudate must
therefore be prevented.
One way of preventing gas escape from the foam is to freeze the skin of the
extrudate [20,2 1,1881. Since the diffusivity drops dramatically as the temperature
decreases [106,189], the rate of gas escape can be substantially reduced by freezing the
skin of the foams. The surface of the extrudate can be quenched by lowering the die
temperature. Therefore, the basic strategy for prornoting large volume expansion in the
extruded foam is to freeze the foarn skin by controlling the die temperature [20].
The melt temperature also needs to be lowered to obtain a high volume expansion
ratio. The temperature of the polymer melt Rowing into the die significantly affects the
amount of blowing agent that escapes to the environment because the diffusion rate of the
blowing agent in the ceIl walls can be retarded by Iowering the temperature of polymer
melt [20]. The polypropylene melt with dissolved blowing agent should be cooled
uniformly down to a very low temperature using a cooling device before the melt gets to
the die. The cooling device can be a heat exchanger consisting of a homogenizing static
mixer on the inside and a cooling medium such as high-pressure gas or cold water
running outside the heat exchanger.
On the other hand, if the melt temperature is too low, the expansion of foam may
not be fostered. The high stiffness of mett at a low melt temperature will retard the cell
growth rate in the beginning and too quick crystaltization of melt will stop cell growth
dunng foam processing. As a result, the final expansion ratio rnay not be high enough.
Therefore, in order to maximize the volume expansion ratio of polypropylene foams, the
stiffness of the polymer meit and the timing of crystallization should be properly
controlled by choosing an appropriate temperature of the polymer melt flowing in the die.
In other words, proper selection of the meIt temperature is extremely important in
controlling the volume expansion of extruded foams.
5.3 Experimental Materials
5.3.1 Selection of Polymeric Foarns
The materiais used in this study were two commercial polypropylene resins
supplied by Borealis GmbH: standard linear pdypropylene m i n and high meIt strensth
(HMS) branched polypropylene m i n . The MFRs (ISO 1133, 230 "C/2.l6 kg) of the
linesir and bnnched propylene materials are 2.8 dg/min and 3.3 dgmin, respectively. The
melt strength and meIt extensibility of these materials cm be seen in Figure 5.1, and both
were measured by using the so-caIled Rheotens method with die UD=30/2, 130 mm/s2
draw-down and melt temperature of 200 OC, fed by Haake lab extruder [187]: a melt
strand is extruded through a capillary die and pulled down with an increasing velocity (at
constant acceleration) by using a pair of wheels, and the force is measured tiIl rupture
occurs. The maximum force is caIIed the melt strength, and the draw-down velocity at
break is a mesure of me1t extensibility. Obviousty, the linear polypropylene has a low
Ievel of melt strength and melt drawability. In contrary, the long-chain branched
propylene polymer shows a very high melt strength in combination with an almost
doubled melt extensibility [l87]. The non-linear increase in force indicates stnin
hardening which is well-known from long-chain bmched materials, such as low-density
pol yethylene.
Figure 5.2 (a) gives the results of the dynamic testing of the polymer rnelts at
230°C by using a Rheornetrics DSR2000 in a frequency sweep test mode. The viscosity
relates to the molecular weight: because the bnnched material has much broader
molecular weight distribution chan the linear one, the viscosity is more dependent on the
shear rate (i. e. shear tfiinning). Whereas at higher frequencies (shear rates) the viscosity
of branched materiai is well below ehat of the Iinear one, the higher viscosities at the very
low frequencies (shear rates) indicate a rnuch higher zero shear viscosity and a significant
contribution of high molecular weight parts. This is due to the post-reactor branching step
in the manufacture of branched materid. The viscosity data obtained from capillary
viscosity rneasurements at 230 "C (Rheograph 2001 from Goertfen) are &en in Figure
5.2 (6) Again, at higher shear rates that correspond to poIymer processing, the long-chah
branched material has a slightly lower viscosity than the linear material. The materials
properties are summarized in Table 5.1.
5.3.2 Selection of Foaming Agents
The type of blowing agent used in plastic foam processing is critically important
in detemining the cell morphology of produced foams. Typically, a physical blowing
agent is used for the low-density foam processing when cross-linking is not involved,
whereas, a chemicai blowing agent is used for low-density foam processing without
cross-linking [190]. Low-density foams with a volume expansion ratio higher than 40
times used to be produced using environmentdly hazardous blowing agents such as
CFCs, HCFCs, etc. in extrusion foam processing. However, because of Montreal
Protocol, these blowing agents have been or will be eliminated from use [19I]. Although.
some efforts have been made to utilize inert gases for low-density foam processing
[20,2 1,1881, long chain blowing agents such as pentane and butane are commonly used in
the low-density foam processing. In order to determine the type of blowing agent CO be
used in the low-density polypropylene foam processing, the maximum achievable
theoretical volume expansion ratios of pentane and butane were calculated based on 5 W C
% of blowing agent. The maximum theoretical volume expansion ratio V, of extruded
foam can be calculated as [?O]:
polymer volume + gas volume v, = polymer volume
where v, is the specific volume of blowing agent at the crystallization temperature of
polypropylene materials (130 OC) and vp is the specific volume of poIypropylent
materiais at room temperature, The specific volumes of pentane and butane are 430 cm3/g
and 600 cm3/g respectively [174,192], and the maximum theoreticai expansion ratios are
21 and 30, respectively, at 5 wt % of blowing agent. With this information, butane is
selected as a blowing agent based on the assumption that the blowing agent efficiencies
are almost the sarne for these blowing agents. The blowing agent efficiency is defined as:
where: cp is the actual volume expansion ratio of the foam.
Talc A7 with a particle size l a s chan 7 pm was used as the nucieating agent. The
talc was supplied by the in-house facilities of Bo~aIis AG.
5.4 Experimental Procedure
The polypropylene pellets, mixed with a fixed amount of taic, were First fed into
the barre1 through the hopper and were completely melred by the screw motion. A
metered arnount of biowing agent was then injected into the extrusion barrel by a positive
displacernent pump at a given weight percentage and mixed with the polyrner melt
Stream. When the gas was injected into the extrusion barre[, the remaining section of the
first screw and the second screw generated a shear field to cornpletely dissolve the gas in
the polymer melt via convective diffusion [la]. The single-phase poIymer/gris solution
went through the gear pump and was fed into the heat exchanger where it was cooled to a
designated melt temperature. The cooled polymedgas solution entered the die and
foaming occurred at the die exit. En general a filmentary die of 0.45 mm diameter and
7.6 mm length was used unless specified. The melt and die temperatures were
synchronised for simplicity in this study. While fixing ail the other materials and
processing parameters such ris the screw speed, gear pump speed, biowing agent content
and barre1 temperatures, the meIt and die temperatures were lowered incrementally and
samples were randomly collected at each designated temperature oniy after there was no
change observed in the melt pressure.
The foam samples were characterized using an optical microscope (Wild
Heerburgg) ancüor a scanning eiectron microscope (SEM, Hitachi 5 10). As a part of the
preparation for characterization of foam structure, the foam samples were dipped in
Iiquid nitrogen and then fractured to expose the cellular morphology,
Characterization of Foam Samples
The volume expansion ratio and the cet1 population density were the structurd
foam paramete* measured. The expansion ratio of foam was deterkined by measuring
the weight and volume of the sample. The volume expansion ratio (9) of each sample
was calcuIated as the ratio of the bulk density of pure polypropylene materials (p,) to ~ h e
bulk density of the foam sarnpk ( p, ) as follows:
The cell population density (N) was calculated as the number of cells per unit
volume with respect to the unfoamed poiymer. First, the number of ceils (nb) in a defined
area ( e x t ) mm' was detemined and then extnpolated to find the total number of cells
per cubic centirnetre:
5.5 Results and Discussion
5.5.1 Effect of Processing Conditions on Volume Expansion and Cell Density
Figure 5.3 (a) depicts the expansion ratio versus the melt temperature for
branched polypropylene materials. Equivalently, the blowing agent efficiency is plotted
against the melt tempenture in Figure 5.3 (b). The arnount of blowing agent used was 5,
10, 15 and 20 wt % and the amount of talc was fixed to be 0.8 wt%. It was observed that
the volume expansion ratio was a sensitive function of the rneit tempenture over the
range investigated. When the melt temperature was as high as L80 O C , the achieved
volume expansion ratio was only about 2 foId. This means chat when the meIt
temperature was too high, most of the gas escaped through the hot skin Iayer of foam
dunng expansion. Kowever, when the temperature of the poiyrner melt was lowered, a
high volume expansion ratio was achieved. Tn other words, the gris Ioss was reduced and
more gits remained in the foam, resuIting in greater votume expansion at a Iower melt
temperature.
It was observed chat there was an optimum temperature that produced the
maximum expansion ratio. When the pmcessing temperature was higher than this
optimum value, the voiume expansion ratio increased as the temperature decreased. IL is
believed that the gas loss was reduced by lowering the temperature because OF the
decreased gas diffusivity. This implies that the voIume expansion behavior was Iirnited
by the gas loss phenomena in this temperature range. The PVT properties also affect the
gas Ioss phenomena due to their direct influence on che viscosity of polyrner materialand
thereby the diffusivity of gas, which eventually has a direct bearing on the expansion
behavior of foams.
When the processing temperature was below the tempenture at which the
maximum volume expansion occurred, the volume expansion ratio decreased as the
temperature decreased. It is believed that the earIy crystailization of polypropylene
materiais because of very low processing temperature was the reason for the lower
expansion ratio. In other words, the poiymer me1t was solidified too quickly before dl the
dissoIved gas could be fully used for the foam expansion.
On the other hand, the melt fracture may dso affect the resuitant expansion ratio
of foams because the non-regularity of surface would not be favounble to expansion in
the radial direction of the filament, Since the melt fracture occurs when the wall shear
stress exceeds the critical die stress [185], the likelihood of occurrence of the melt
fracture increases at a low die tempenture. As the wall shear stress increases further,
ultimateiy gross melt fracture occurs and extrudate becomes extremeiy non-uniform. As
can be anticipated, the expansion ratio is reduced with the onset of surface irregularities
and continues to decrease with increasing melt fracture.
Figure 5.3 (c) depicts the ceII density versus the meIt temperature for branched
polypropylene materiais. Equivalently, Figure 5.3 (d) depicts the ce11 density versus the
amount of butane injected in the polymer melt. The figures show that the ceIl density was
unaffected when the tempenture of the rneit was varied. Since the sarne die was used in
ail the experiments, varying the met tempenture resulted in a variation of the pressure
built due to the resistance of the die as shown in Figure 5.3 (e). This impiied varying the
melt pressure had no major effect on the celi density. This rnay be attributed to the high
solubiIity of butane in the polymer. AIthough, there was no specific litenture data
availabIe about the solubility of butane in poIypropyIene at high pressures and high
tempentures, the solubility of haiogenized hydrocarbon (FC I i) in polypropyiene was
reported to be as high as 90 WC% at 13.8 MPa and 200 OC [I93]. It was therefore expected
that the solubility of butane in polypropylene would be also high. The amount of injected
butane in the range of 5-20 wt% must have been far below the solubility Iimit under the
processing conditions and the entire amount of injected butane was therefore dissolved in
the polymer melt regardless of the processing pressure. In other words, the
thermodynamic instability induced from the dissolved butane and polypropylene matrix
was not affected by the processing pressure. As a result, the driving force to nucleate
bubbles would not Vary as the processing pressure was increased. It appeared thac the cell
density was mainly determined by the talc content [51] and the amount of butane.
5.5.2 Effect of Blowing Agents on the Volume Expansion Ratio and Cell Density
5.5.2.1 Effect of the Blowing Agent Amount
One of the most critical factors affecting the foaming behavior of polypropylene
is the amount of blowing agent injected. Figure 5.3 (a) shows that the largest expansion
ratio achieved was a strong function of the amount of butane injected. The largest volume
expansion ratios achieved for 5, 10, 15 and 20 wt % butane contents were 20, 50, 68 and
90 fold, respectively. On the other hand, the efficiency of blowing agent for 5, 10, 15 and
20 wt 95 butane contents was 69%, 84%. 76% and 7596, respectively (Figure 5.3 (b)).
Figure 5.3 (e) shows chat the working pressure was significantly lowered as the amounc
of injected butane increased from 5 wt % to 20 wt 95. The lowered pressure must have
been due to the plasticizing effect of dissolved blowing agent in the polymer matrix (251.
It was also notable that the maximum volume expansion temperature was lowered with
an increase in the butane content because of the plasticizing effect.
The bloking capability of butane to obtain a low-density foam is well known
together with other long chain blowing agents such as pentane and propane in the foam
processing of po[ystyrene and polyethylene (11. However, the effectiveness of butane for
the foaming of polypropylene materials has not been verified in the literature. It seems
that butane is a very effective blowing agent for polypropyiene in achieving a Low-
density foam as Iong as the resin has high melt strength. However, because of its high
flarnmability, the amount of butane used for achieving a desired expansion ratio should
be minimized.
Figure 5.4 shows the foaming behavior of linaar polypropylene materials with 0.8
wt % talc. The maximum volume expansion ratios achieved for 5, 10, 15 and 20 wt %
butane contents were 3, 32. 38 and 46 times, respectively. The volume expansion ratio
was not significantly increased as the amount of injected butane increased above 10 wt %
unlike in the case of the branched materials. This seemed to be due to the severe ceIl
coalescence sustained by the foam extrudate, and ce11 coalescence seemed to be
accelerated by the plasticizing effect of linear poiypropylene materials as the amount of
dissolved gas in the polymer melt was increased. Correspondingly, the efficiency of the
blowing agent was IO%, 539'0, 41% and 38% for 5. 10, 15 and 20 wt % of butrine
contents, respectively. Figure 5.4 (e) shows a significant drop of the working pressure
due to the piasticizing effect of blowing agent as the amount of injected butane increued
Crom 5 to 20 wt %.
Figures 5.3 (d) and 5-4 (d) show thrit increasing the concentration of bu tane in the
polyrner rnelt resulted in a higher ceII density. This result was expected because o f a
greater driving force to nucleate bubbles at a higher blowing agent concentrrition. In other
words, when the potymer melt had more dissolved blowing agent, a greater
therrnodynamic instability was induced as the polymer melt exited from the foaming die
because of the solubility drop. As a consequence, a higher cell density was devetoped in
the extmded foarns. Simihr trends had been observed for isopentane-polypropylene and
CO2-polypropylene combinations [4]. However, as in the case of the isopentane-
polypropyIene system. the sensitivity of the ceIl-density increase with respect to the
butane content was not as high as that with respect to the COl content in polypropylene.
Furthemore, because of a high flammability, the amount of butane used in the foam
processing should be rninirnized and it wouId not be a good strategy to increase the
butane content to increase the cell density.
5.5.2.2 Effect of Blowing Agent Type
Figure 5.5 iilustrates the foarning behavior of braoched polypropylene with COz
as blowing agent. The injected amount of COz was 2, 5, 10 and 15 w&%. Figure 5.5 (a)
shows the effect of the processing temperature on the expansion behavior of branched
polypropylene using CO2 as blowing agent. Equivalently, the blowing agent effectiveness
is plotted against the die temperature in Figure 5.5 (b). In the case of CO2, low volume
expansion was obtained at higher temperature, however, by decreasing the temperature a
sIight increase in the volume expansion was observed, Also, the volume expansion
increased slightIy when the amuunt of CO: increased from 2 to 15 wt%. The maximum
volume expansion achieved was 18 using 15 wt% CO?. This result cari be attributed to
the Iow solubility of CO: in the polymer rnelt. Also, the high diffusivity of CO2 from the
extrudate promoted gas loss and resulted in a low volume expansion.
Figures 5.5 (c) and 5.5 (d) depict the efkct of using CO: as a blowing agent on the
ce11 density of bnnched polypropylene materials. The figures show that the ce11 density
increased (up to 10' cells/cm3) when the COz amount increased from 2 to 10 wt%,
however, by further increasing the arnount of CO?, the ceil density decreased. In general
a high ceIl popdation density was achieved when using COzcompared to butane, this is
due to the fact that CO2 cm act as a nucleating agent as well as a blowing agent El],
Figure 5.5 (e) shows the effect of incieiuing the amount of C01on the processing
pressure üt the die. The figure show that by increasing the arnount of CD2 from 2 ro 15
wt% , the processing pressure at the die decreases. This is due to the phsticating effect of
gas in the polyrner, However, the processing pressure for CO2 was much higher than that
of butane, as the solubility of butane in the polymer melt and hence the plasticizing effect
is much higher than that of CO-.
5.5.2.3 Eflect of the Blowing Agent blending
The effect of using bIends of butane and CO1 as blowing agents on the volume
expansion was investigated in this study. The bIends used were 75/25, 50/50, and 25/75
wt%, butane/C02, while the total blowing agent amount was fixed to be 10 wt%, and the
nucleating agent used was talc with a fixed amount of 0.8 wt%. Figure 5.6 (a) shows the
effect of blending on the volume expansion for various die tempemures. Tt was observed
that the volume expansion increased when the mount of butane in the blend increased.
The volume expansion for pure CO2 was around 4-fold and by increasing the amount of
injected butane to 25 wt%, the volume expansion did nat change rnmuch. However, when
the amount of butane increased to 50, and 75 wt%, the volume expansion increased to I O
fold and 32 fold respectively. When using pure butane, the volume expansion reached a
maximum value of 50 fold. As discussed in an eartier section, the reason for the increase
in the volume expansion is attributed to the high solubility and Iow diffusivity of butane
in the propylene melt compared to CO2. It seems that the blending of blowing agents is
not favorable in terms of cell nucleation and expansion. The existence of CO2 in the
blend adversely affected the expansion of propylene. For example, the expansion ratio of
foams blown with 5 wt% of the biowing agent mixture (5 wt% butane and 5 wt% CO2) is
much smaller than those bIown with S wt% butane only. Equivalently, Figure 5.6 (b)
describes the effect of using blends on the blowing agent efficiency.
Figure 5.6 (c) shows the effect of blending on the cell density for various die
temperatures. Equivalently, Figure 5.6 Id) depicts the cell density versus the relative
amount of CO2 injected in the polymer meit. In general, it was observed that the ce11
density increased when the amount of CO7 in the blend increased. The ceIl density for
pure butane was around 5x10~ cells/cm3 and by increasing 25, 50 wt% of COz, the cell
density did not change much. However, when the rimount of CO2 increased to 75 wt%,
the cell density increased to 8x 106 cells/cm3. For the pure CO2, the cell density reached a
maximum value of 8x 10' cells/cm3. As discussed in an earlier section, the reason for the
increase in the cell density is attributed to the fact that CO2 acts as both a nucleating agent
and a blowing agent. It seerns that the existence of butane does not affect the nucleatinp
behavior of COz, Figure 5.6 (e) describes the effect of using blends on the die pressure. It
was observed that the die pressure increcised when the amount of CO2 increased in the
blend. This is due to the low solubiiity of COz compared to butane in the propylene
materiais. By having more CO2 in the blend, the overall solubiIity of the mixture
decreased, thereby decreased the plasticising effect.
5.5.3 Effect of Nucleating Agents on the Volume Expansion Ratio and Cell Density
The nucleation and volume expansion behaviors of HMS branched polypropylene
were investigated using various talc contents as nucieating agents for IO wt Ojo injected
butane. The effects of changïng the tdc content on the foamability of polypropylene
materials were studied.
Figure 5.7 (a) shows the effect of increasing the amount of talc on the volume
expansion behavior of bnnched poiypropylene resins. It was observed that at high
temperature, the volume expansion ratio decreased as the tdc amount increased. It is
believed that as the talc content increased, more nucIeated cells are produced in the
polymer melt and the ceIl growth rate increased. The blowing agent, which has difised
into the nucleated cells eventually diffused out to the atmosphere, because the complete
separation of two phases is thermodynamically most favorable [188]. Gas escape through
the thin walls and eventually through the foam skin decreased the amount of gas available
for the growth of cetls. Also, the high diffusivity of the blowing agent at the elevated
temperatures increased its diffusion rate during expansion. As a result, the polymer me1t
that formed the ceIl walls did not freeze at an appropriate time after the ce11 structure was
forrned, and the cells collapsed causing foam contraction because of gas escape from the
cells. The elastic recovery of the polymer me1t also caused the stretched ceIl walls to
shrink, and thereby caused the foam to contract.
it was also observed that by adding 0.8 wt% talc to the poiypropylene rnelt, the
maximum volume expansion increased up to 50-fold. However, when the rimount of talc
increased above 0.8 WC%, the maximum volume expansion ntio decreased again. This
means that there was an optimum talc content to achieve a maximum volume expansion
ratio in the foam processing of polypropylene materials. This c m be attributed to the
unnecessq high pressure buih in the die when the arnount of talc increased. This causes
early crystailization of the foam extrudate. which led to freezing of the foam and an
increase in its stiffness, which the caused ce11 growth to stop.
Figure 5.7 (b) shows the blowing agent efficiency (q), defîned in Equation (5.2).
as a function of the die tempenture. It was observed that the blowing agent efficiency
decreased with an increase of the amount of talc. At high temperature, this is due to the
increased ce11 nucleation, which accererated the gas loss, and at low temperature this is
caused by the increased stiffness of the f o m extmdate.
Figures 5.7 Cc) and 5.7 (d) show the effect of using talc as ri nucleating a, aent on
tfie ce11 density of branched polypropylene resins using 10 wt% injected butane, The
figures show that for a higher talc content, a higher ceIl density of polypropylene foams
was achieved. Even a small amount of tdc particles significantfy increased the ce11
density. The cell density increased up to 6x10' cells/cm3 when the amount of talc
increiised frorn O to 2A wt%. By increasing the amount of talc in the polymer meIt it
creates more nucleation sites by reducing the Gibbs free energy required to form bubbles
[143]. This promotes heterogeneous nucleation, and a high ceIl density is achieved.
Figure 5.7 (e) shows the effect of increasing the amount of talc on the processing
pressure at the die for 10 wt% injected butane. The figure shows that by increasing the
amount of talc, the processing pressure at the die increases. It is believed that the taIc
particles increased the viscosity of the melt resulting in a higher-pressure profile at the
die.
5.5.4 Effect of Long Chain Branching on the Volume Expansion Ratio and Cell
Density
In this section, the foaming behaviors of branched and linear polypropylene
materials are compared, A cornparison of Figures 5.3 (a) and 5.4 (a) shows that a higher
expansion ratio was achieved from the branched polypropylene materials (up to 90 times
for 20 wt. Q butane) whereas from the linear polypropylene materials, the maximum
volume expansion was only 46 times for the sarne content of butane. This can be
explained by the different extensional behavior of the melts (see Figure 3.1): in the case
of the long-chain branched polypropylene with high melt strength and high melt
extensibility, ceIl coalescence is much more effectively prevented compared to the linear
polypropylene. Consequently, the linear polypropyIene material could not reach large
volume expansion like the branched polypropylene. CeIl coalescence seemed to be more
dominant at a high concentration of butane because of the plasticizing effect as
mentioned earlier.
Figures 5.3 (c), 5.4 (c) and Figures 5.3 (d), 5.4 (d) show that the average ceIl
density for linear polypropylene materials was slightly higher than that of branched
materials using 5 wt % injected butane. On the other hand, at higher amounts of blowing
agents, the average cell density for branched polypropylene materiais was slightly higher
than that of linear materials. It was not ciear whether the trend of the increasing ce11
density in branched resin was due to the inherently better ce11 nucleability of the branched
resin or due to a lower degree of ce11 coaiescence. Also, there is a high probability that
cell coalescence in the ce11 growth stage of the linear resin :vas severe at higher amounts
of blowing agent, and consequently, the resultant ceII density was deteriorated despite a
high ce11 nuclei density obtained in the ce11 nucleation stage.
5.5.5 Effect of Materials Blending on the Volume Expansion Ratio and Ce11 Density
The effect of using blends of branched and Iinear polypropylene resins on the
volume expansion behavior was a central issue because low-density polypropylene foam
is the primary target of this study. The foaming additive used in this study was talc as the
cell-nucleating agent with a fixed amount of 0.8 WC%. The blowing agent used in the
experirnents was n-butane, with a fixed amount of 10 WC%. The applied blending ratios of
linear to branched components were 80/20,50/50, and 20180 wt8, nspectively [194].
The melt strength and melt extensibility of these materials can be seen from
Figure 5.8, both measured by using the Rheotens method described in section 5.3.1. The
pure linear PP (sample 1 ) has low levei of melt strength and melt drawability cornpared
to the pure branched HMS propylene polymer (sample 2) as discussed before. When
looking at the blends of pure linear and pure branched propylene polymers (samples 3.4,
5) it can be clearly seen chat the melt strength, meIt extensibility and strain hardening
behavior increase with the amount of added branched HMS material.
Figure 5.9 (a) shows the effect of blending on volume expansion for various die
temperatures. in general, it was observed that volume expansion increased as the arnount
of branched material in the blend increased. Aithough the maximum volume expansion
for pure linear material was about 25 fold, by increasing the branched material content to
20,50,80, and 100 wt%, the volume expansion ratio increased to 28,30,40, and 50 fold,
respectively. Again. the reason for the improvement of the volume expansion with an
increase in the branched materid content can be attributed to the increase in the melt
strength and melt extensibility (see Figure 5.8), which prevented ce11 coalescence. Figure
5.9 (b) shows the blowing agent efficiency (q).
As observed in Figure 5.9 (a), the temperature required for obtaining the
maximum votume expansion ratio decreased as the content of branched resin increased.
This is due to the lower viscosity and the increased melt strength and melt extensibility
together with strain hardening behavior of the branched resin, broadening the foaming
window towards lower tempentures. The blowing agent efficiency increases with an
increase of the amount of branched min, It should be noted that a high blowing agent
efficiency would be desirable in terms of ffarnrnabiïity and blowing agent cost because
Iess amount of blowing agent will be required for achieving the same volume expansion
ratio. However, the higher resin cost of branched material would limit its content in the
blends, which is very depending on the application and the required foam properties, too.
The effect of using the blends of branched and linear propylene poIyrners on the
ceIl density is shown in Figures 5.9 (c) and 5.9 (dl as a function of the die temperature,
and the arnount of Iinear materials in the blend, respectively. Tt was observed that the cell
density increased as the amount of branched mnterial in the blend increased. The cell
density For pure linear milteriai was about 3x10' cells/cm3 and by increasing the branched
material content to 20, 50, 80 and 100 WC%, the cell density increased to 8x105, 2 x 1 0 ~ ~
5x106, and 8 x 1 0 ~ cells/cm3. respectively. It was not clear whether the trend of the
increasing cet1 density with increasing in branched resin was due to the inherently better
cell nucleability of the branched resin or due to a Iower degree of cell coalescence. Park
et d ' s studies [4,541 on ceIl nucleation of Iinear and branched polypropylene resins
showed that the linear resin had better cell nucleability, which represents a reverse trend
to the currently observed results. Therefoce, one may not be able to conclusively state chat
a branched resin is better or worsc than a linear resin in terms of ceII nucleability. On the
other hand, there is a high probability that ceIl coalescence in the ce11 growth stage of the
Iinear resin was severe, and consequent1 y, the resuI tant ce11 density was degraded despite
a high cet1 nuclei density obtained in the stage of ce11 nuclerition. A further study is
required to clarify this issue. Figure 5.9 (e) shows that the die pressure increased as the
arnount of Iinear material in the blend increased- This was must likely due to the high
viscosity of the linear materials.
It was expecced chat blending of a branched propylene resin with a linear
propylene resin will enhance the fomability because of the enhanced rnelt strength and
significantly increased meit extensibility (see Figure 5.8). However, proper dispersion of
the branched min in the Linear poIypropyIene matrix may not be accomplished through
blcnding in a regulu foaming extruder or a compounder. Severe mixing action of
branched material to achieve even dispersion, will cause disentanglement of branched
molecules, whereas weak mixing action wilI resuit in pour dispersion of branched
material in the linear resin, if the ceII-to-cell distance is rnuch Iarger than the striation
thickness of the linear and bmched materials, then the melt strength of the ce11 wall will
be almost proportional to the weight fraction of the branched materiai in the blend.
However, if the cel1-fo-ce11 distance is almost the same or srnalier than the striation
thickness, the blend cm not be considered homogeneous with respect to the cell structure.
In this case, the ceIIs are formed in the Iinear region as well as in the branched region.
Then, the cells in the linear region will be easily ruptured and severe ceII coalescence will
occur. Since the striation thickness of the linear and branched polypropylene blends
cmnot be small, a fine-celled structure may not be easily developed in the blends of
linear and branched propyiene resins. The difficulty of homogeneously dispersing the
branched resins in the linear matrix and resuitant deterioration of ceil structure of blends
wiII be severer as the content of branched resin decreased in the blends.
55.6 Effect of Using Re-Extruded Polypropylene Materials on the Volume
Expansion Ratio and Cell Density
The effect of using re-extruded branched propylene resins on the volume
expansion behavior is investigated in this study. The materials used were branched
polypropylene mateflds, and 2 grades of re-extruded branched polypropyiene materiaIs.
They are denoted in this study as Branched Pl , Branched P3 and Branched P3,
respectively. Figures 5-10 (a), and 5.1 L (a) show the volume expansion behaviors of
Branched P l and Branched P2 for 5, 10, 15 and 20 wt 9% injected butane. It was observed
chat the maximum achievable volume expansion decreased in the case of Branched Pl
compared to the onginal material (Figure 5.3 (a)). Moreover, much lower volume
expansion was obtained in the case of Branched P2, which has a higher meIt flow rate (8
dmn) than the original materiai (2.3 ddrnin). This is believed to be due to the breakage
of long chain branching andor to segmentation of branched moIecuIes. When the
materiai is re-extruded, the melt strength is decreased and hence the voiume expansion is
reduced. Figures 5.10 (b), and 5. t i (b) show the effect of using re-extruded branched
materiais on the blowing agent eficiency. The increased high shear rate viscosities of
these recycled rnatetials played an important roIe in the expansion behavior of these
rnaterials. What happens to these materials during the recycling process needs to be
clarified through further research.
Figures 5-10 (c), 5. IO (d) and Figures 5.1 1 (c), 5.1 1 (d) show the effect of using
the re-extmded branched materials on the ce11 density for 5, 10, 15 and 20 wt % injected
butane. Tt was obsewed that the ce11 density did not change in the case of Branched Pl ,
however, the cell density decreased in the case of Branched P2. This effect is due to the
breakage of long chain branching when the rnaterial is re-extruded, which decreased the
rnelt strength and hence facilitated lower ce11 densities. Figures 5.10 (e), and 5.1 1 (e)
show the effect of using re-extmded branched materials on the die pressure. The Figures
show the platicizing effects when increasing the amount of butane content as discussed
earlier.
5.5.7 Effects of Die Geometry on the Volume Expansion Ratio and Cell Density
The effects of die geometry on the volume expansion behavior of branched
polypropylene were investigated in this section. A lower pressure drop rate die (denoted
as Die B) of 0.77 mm diameter and 12.7 mm length [17] was chosen in this study and its
results is compared with the previously used die (denoted as Die A).
Figures 5.12 and 5.13 show the foaming behavior of branched polypropylene with
butane and COI as blowing agents, respectively, using Die B. The injected arnount of
butane was 5, 10, 15 and 20 wt 8, while the injected amount of COz was 2,s 10 and 15
wt %. Figures 5.12 (a) and 5-13 (a) depict the expansion ratio versus the die temperature.
Equivalently, the blowing agent effectiveness is plotted against the die temperature in
Figures 5.12 (b) and 5.13 (b). The volume expansion ratios of Die B were overdl lower
than those of Die A (see Figure 5.5 (a)). This means that a high pressure drop rate is
favorable to the expansion ratio. The lower pressure drop rate die has a higher chance of
developing an initiai hump because of the longer ceIl growth time allowed in the die after
the nucleation point, and therefore, gas loss will be promoted more easily. In the case of
CO?, sirnilar trends were obtained and lower volume expansion was obtained using Die B
compared to Die A (see Figure 5.5 (a)).
The effects of pressure drop rate or die geometry on the ce11 population density
were also investigated. Figures 5-12 (c) and 5.12 (d) show the ce11 nucleation behavior
using butane as a blowing agent for die B. It was observed that at al1 butane contents, the
ce11 population density obtained from Die A (Figure 5.3 (c) and 5.3 (d)) was much higher
than that of Die B. This is due to the fact that the ce11 population density is a sensitive
function of pressure drop rate achieved in the die 1171. Sirnilarly, Figures 5.13 (c) and
5.13 (d) show the ce11 nucleation density as a function of die temperatures, and butane
contents, respectively, for branched polypropylene with 2, 5, 10 and 15 wt% CO2
contents using Die B. Figures 5.12 (e) and 5.13 (e) descnbe the die pressure as a function
of die geometry for Die B using butane and CO; as blowing agent, respectively.
5.6 Statistical Analysis of the results
StatisticaI design and anaiysis for the volume expansion of branched
polypropylene foam were developed using CARD-PRO design of experirnents software
[195]. The set of variables selected for the experimental design was: die temperature,
blowing agent percent, additive percent, mass flow rate and die geometry. VariabIesi
ranges and restrictions were specified and an experiment design was generated based on
these values. Experiments were then checked according to this proposed design and the
results data were analyzed. The experimental error statistics including the error
percentage and the confidence level were computed based on the experimental data.
Regression statistics for the volume expansion results were conducted using N O V A
method. An empirical mode1 and its variables coefficients were determined; also the
mode1 term ranking for the various variables was genented in Pareto charts. The detailed
statistical analysis results are shown in Appendix 2.
5.7 Summary and Conclusions
Experimentai studies were carried out to manufacture low-density polypropylene
foams in extrusion using a long-chain blowing agent. The experiments conducted in this
study Iead to the following conclusions:
1. The basic strategies employed for the promotion of a large volume expansion ratio
with polypropylene materiais was identified.
2. An extremdy large expansion ratio up to 90 times was successfulIy obtained from
the branched polypropylene materials by tailoring the proçessing conditions.
The effects of processing parameters such as the temperature and pressure, the
materials parameters such as the blowing, nucleating agents, and long chain
branching, and the die geometry on the production of low-density, fine-ceiled
polypropylene foms were investigated,
There exists an optimum temperature for achieving the maximum expansion ratio of
polypropylene foam with butane. This optimum temperature decreased as the
amount of blowing agent increased because of the plasticizing effect.
Despite its high flammability, butane is a very effective blowing agent for low-
density polypropyfene foams.
A higher ceIl density and lower volume expansion ratio were obtained by increasing
the amount of CO2 when using a blend of CO: and butane as blowing agents.
The ce11 density increased proportionally when the amount of talc increased from O
to 2.4 %. However, there was an optimum talc amount for achieving the maximum
volume expansion ratio.
The expansion ratio obtained from the linear polypropylene materials was much
lower than that frorn the branched polypropylene materials because of severe ce11
coalescence. The expenmental results indicate that branched polypropylene
materials are effective for low-density foam application because of the reduced
degree of cell coalescence.
The use of a blend of linear and branched polypropylene materials revealed that the
ceII density in the foamed material increased as the amount of branched
polypropylene was increased in the blends. Despite the fact that there were no clear
governing mechanisms observed for the ce11 density of linear and branched
polypropylene blends, the ceIl coalescence could obviously be reduced by the
addition of branched HMS material, giving the blend higher melt suength and rneIt
extensibility.
IO. The voiume expansion increased when the amount of branched materid in the blend
increased. This cm be explained in the same way. The optimum temperature for
producing the maximum volume expansion ratio decreased as the branched material
content was increased. This is believed to be due to the high viscosity of the linear
propylene together with the extensional characteristics of the added KMS
component.
I l . The use of re-extruded branched materials resulted in lower ceIl density and volume
expansion ratio, This was attributed to the breaking of the long chain branching of
the high rnelt strength polypropylene, which led to cell coalescence in the foam
structure,
12. The maximum volume expansion ratio was decreased when using a lower pressure
drop rate die, However, the expansion ratio was higher with the low pressure drop
rate die in the high temperature range. The cell density was higher in the case of
higher pressure drop rate. Further studies are needed to clarify the effects of die
geometry on the volume expansion behavior.
Figure 5.1: Melt strength and meir extensibility of polypropylene rnelts (Goettfert Rheotens: 230 OC; L/D=30/2 mm; 120 mrn/s2 acceleration)
Figure 5.2 (a): Complex viscosity of polypropylene melts (Rheometrics DSR2ûûû; platelplate configuration; frequency sweep a& 230 O C )
100 I l
1 10 100 1000
Apparent shear rate (radis)
Figure 5.2 (b): Apparent viscosity oFpoIypropyIene melts (Goettfert Rheograph 200 1; 230 OC)
" 100 120 140 160 1110
Ternpntun (C)
Figure 53 (a): the expansion ratio versus the butane content for branched polypropylene materials using butane as biowing agent
Figure 5.3 (b): the biowing agent efficiency versus the rnei t temperature for branched polypropylene materials using butane as biowing agent
Figure 5.3 (c): the ce11 density versus the meIt temperature for branched polypropylene materiais using burane as blowing agent
Figure 5.3 (d): the ce11 density versus butane contents for branched polypropylene materids
Figure 5.3 (el: the die pressure versus the meIt temperature for branched polypropylene materials using butane as blowing agent
- t 4 50 C
g 40 W
30
20
10
O
100 120 140 160 110
Tempantun (C)
Figure 5.4 (a): the expansion ratio versus the melt temperature for linear polypropylene materials using butane as blowing agent
100 120 11P 1W 180
Ternpmtun (C)
Figure 5.4 (b): the blowing agent efficiency versus the melt temperature for linear polypropyIene materials using butane as blowing agent
1W 120 110 160 180
Tmnptalun (C)
Figure 5.4 (c): the ceii density versus the meIt temperature for linear polypropyiene materials using butane as blowing agent
O 5 10 16 N ZS
Butina &nmi(%l
Figure 5.4 (d): the cell density venus butane contents for Iinear polypropylene materials
Figure 5.4 (e): the die pressure versus the meh temperature for linear polypropylene materials using butane as blowing agent
= ~~r-~7- -tlO%COI - - - +i% COI
2 44% COI
-. 100 120 140 160 110
T i m p i n i u n (C)
Figure 5.5 (a) the expansion ratio versus the melt temperature for branched polypropylene materials using CO, as blowing agent
Figure 5.5 (b) the blowing agent efficiency versus the meIt temperature For branched polypropylene materials using CO, - as bIowing agent
Figure 5.5 (c) the cell density versus the melt temperature for branched polypropylene materials using CO, - as blowing agent
Figure 5.5 (d) the ce11 density versus CO2 contents for branched polypropylene materials
Figure 5.5 (e) the die pressure versus the melt temperature for branched polypropylene materials using CO, as biowing agent
Temperabin (C)
Figure 5.6 (a) the expansion ratio versus the melt temperature for branched polypropylene materiais using a blend of butane and CO, as blowing agent -
Figure 5.6 (b) the blowing agent efficiency versus the melt temperature for branched po1ypropy:ene materials usinp a biend of butane and CO, as blowing agent
Figure 5.6 (c) the cell density versus the melt temperature for branched polypropylene materials using a blend of butane and CO, - as blowing agent
lf*0 O 20 40 60 80 100
RihUvr CO2 Contint h W n (?&)
Figure 5.6 (d) the cell density versus CO2 relative amount for branched polypropylene materials using a blend of butane and CO, as biowing agent -
Figure 5.6 (e) the die pressure versus the meIt temperature for branched potypropylene materials using a blend of butane and CO, as bIowing agent
t O d X Talc
10 -1.6% Talc - c Z . 4 % Talc
"- 60 à 1 : 50 - I
g 40
30
20
10
O 100 120 140 160 110
Temperature (C)
Figure 5.7 (a) the expansion ratio versus the melt temperature for branched polypropylene materials using butane as blowing agent for various talc contents
Figure 5.7 (b) the blowing agent efficiency versus the melt temperature for branched polypropylene materials using butane as blowing agent for various talc contents
100 120 140 160 110
Timprnhirr (C)
Figure 5.7 (c) the cell density versus the melt temperature for branched polypropylene materials using butane as blowing agent for various talc contents
1.E4 ! l.E+OQ 4
O M 1 1 . 5 2 2 5
Talc Conion! ( X )
Figure 5.7 (d) the ceil density versus the melt temperature for branched polypropylene materials using butane as blowing agent for various talc contents
Ttmprnlurr (C)
Figure 5.7 (e) the die pressure versus the melt temperature for branched polypropylene materiais using butane as blowing agent for various talc contents
Drawdown velocity in mmb
Figure 5.8 the melt strength and melt extensibility of linear and branched polypropylene materials blend
Figure 5.9 (a) the expansion ratio versus the melt temperature for linear and branched polypropylene materiais biend using butane as blowing agent
Figure 5.9 (b) the blowing agent efficiency versus the meIt temperature For linear and branched polypropylene matecïds blend using butane as blowing agent
Figure 5.9 (c) the ce11 density versus the melt temperature for linear and branched polypropylene materials blend using butane as blowing agent
Figure 5.9 (d) the ce11 density versus linear polypropylene content for Iinear and branched poIypropylene materials bIend using butane as blowing agent
Figure 5.9 (e) the die pressure versus the rnelt temperature for linear and branched polypropylene materiais blend using butane as blowing agent
Figure 5.10 (a) the expansion ratio versus the melt temperature for re-extruded Brancheci PI polypropylene materials using butane as blowing agent
Figure S. 10 (b) the blowing agent eficiency versus the melt temperature for re-extruded Brmched P l polypropylene materials using butane as blowing agent
Figure 5.10 (c) the ceIl density versus the melt temperature for re-extruded Branched Pl polypropylene materials using butane as blowing agent
Figure 5.10 (d) the ce11 density versus the butane contents for re-extmded Branched P 1 polypropylene matends
Figure S. 10 (e) the die pressure versus the melt temperature for re-extruded Branched P 1 poIypropylene mateciais using butane as blowing agent
100 120 140 160
Ternpwrlurr (C)
Figure 5.1 1 (a) the expansion ratio versus the melt temperature for re-extruded Branched P2 polypropylene materials using butane as blowing agent
Figure 5.11 (b) the bIowing agent efficiency versus the melt temperature for R-extruded Branched P2 polypropylene materiaIs using butane as blowing agent
i oa 120 140 160 180 t r m p r i u i n (C)
Figure 5.1 I (c) the cell density versus the melt temperature for ce-extruded Branched P2 polypropylene materials using butane as blowing agent
Figure 5.1 1 (d) the ce11 density versus butane contents for re-extruded Branched P2 polypropylene materials using butane ris blowing agent
O J 100 120 110 160 110
Trmptnmn (C)
Figure 5.1 1 (e) the die pressure versus the melt temperature for re-extruded Branched P2 polypropylene mater& using butane as blowing agent
- . 100 120 140 160 110
Ternperiluie (C)
Figure 5.12 (ri) the expansion ratio versus the melt temperature for branched polypropyicne materials using butane as blowing agent for Die B
- ,
100 120 140 160 180
Timp.ritur~ (C)
Figure 5.12 (b) the blowing agent eficiency versus the melt temperature for branched polypropylene materials using butane as blowing agent for Die B
Figure 5.12 (c) the cell density versus the melt temperature for branched polypropyiene materials using butane as bIowing agent for Die B
Figure 5.12 (d) the ce11 density versus butane contents for branched polypropylene materials for Die B
40 T
Figure 5-12 (e) the die pressure versus the rnelt temperature for branched polypropylene materials using butane as blowing agent for Die B
Figure 5.13 (a) the expansion ratio versus the melt temperature for branched polypropylene materials using CO, - as blowing agent for Die B
Figure 5-13 (b) the blowing agent efficiency versus the melt temperature for branched polypropylene materials using CO, as blowing agent for Die B
Figure 5.13 (c) the cell density versus the melt temperature for branched polypropylene matends using CO, - as blowing agent for Die B
Figure 5.13 (d) the cell density versus CO2 contents for branched polypropylene materials for Die B
Figure 5-13 (e) the die pressure versus the melt temperature for branched polypropylene materials using CO, - as blowing agent for Die B
Chapter 6
Fundamental Mechanisms of Volume Expansion Behavior of Polypropylene Foam Filaments
6.1 Introduction
In the previous chapter, the effects of processing and materials parameters on the
foaming of polypropylene materials were thoroughly investigated, The results show that
the volume expansion ratio was govemed by the processing and materials parameters. It
was shown that the processing temperature, the blowing agent arnount and type, the
nucleating agent amount and type, the materials branching, and the die geometry are
affecting the finai volume expansion ratio of polypropylene foains. in order to improve
the expandability of polypropylene foams, we need to have a clear understanding of the
fundamental mechanisms that govern the volume expansion ratio. Despite the recent
snidies that addressed the production of low density, fine-celled polypropylene foams, no
research has been conducted to investigate the mechanisms that govern the expandability
of polypropylene foams.
In this chapter, a qualitative modeling for the volume expansion behavior was
described. initially, the fundamental mechanisms goveming the volume expansion of
polypropylene foams were identified and the volume expansion phenomena were
described based on our expetimental observation. The system setup used for monitoring
the expansion phenomena of extmded foam is then described. The procedure for
monitoring the expansion mechanisms and the image analysis will be elucidated.
Consequently, the effects of processing parameters on the volume expansion behavior
were depicted based on the extrudate images. A theoreticai mode1 for calculating the
voIume expansion ratio, ce11 size, and ceIl wail thickness was developed as a hnction of
extrudate diameter. The mode1 was used to describe the volume expansion behavior at
various distances from the die using the observed CCD images of extruded foams.
6.2 Qualitative Modeling of Volume Expansion Behavior
6.2.1 Fundamental Mechanisms Governing Volume Expansion of Polypropylene
Foams
A careful analysis of extended experirnental results obtained at various processing
conditions indicates chat the final volume expansion ratio of the extruded polypropylcne
foarns blown with butane is governed either by Ioss of blowing agent through the foam
skin or by crystailization of the porymer matrix. In general, upon exiting the die, the
foaming extrudate exhibits one of the shapes depicted in Figure 6.1, depending on the
temperature of the extrudate. At higher temperanires, rhe cross section of the extrudate
expands suddenly, Le., it has a higher initiai angle 0, and this iingie decreases as the
processing temperature decreases. Below an optimum processing temperature, the
foaming of extrudate is inhibited and the initial angie is substantially reduced. The
mechanisms governing the volume expansion of polypropylene foams are explained in
the following sections.
62.1.1 Gas los
The gas ioss phenornena which occur during f o m processing cm be correlated
with the melt temperature. The diffusivity of bIowing agents at elevated temperatures is
very high, therefore, if the processing temperature is too high, the gas cm easily escape
from the extruded foam because of its higher diffusivity at eIevated tempennires. In
addition, as the ceII expansion increases, the cell w d l thickness decreases and the
resulting rate of gas diffusion between cells increases. Consequently, the rate of gas
escape from the foam to the environment increases. Gas escape through the thin cell
wails decreases the amount of gas availabie for the growth of cells resulting in lowered
expansion. Moreover, if the cells do not freeze quickly enough, they tend to sht-hk due to
Ioss gas through foam skin causing overail f o m contraction. This mechanism is
schematicdly shown in Figure 6.2.
This phenornenon of the gas escape at hi& temperatures and the resultant foam
contraction during the ceIl growth stage at elevated temperatures was observed during
experimentation, It was observed that the extruded foam initially expanded as the
nucleated cells grew very fast and then eventually contracted. The maximum expansion
occurred very close to the die exit with the diameter of the initially expanded foam being
larger than that of the final foam extrudate. This indicates that the volume expansion ratio
of the initially expanded state was considerably high, and therefore, the shapes of
expanded cells at this state were not spherical but polyhedral with thin cell walls (Figure
6.2).
6.2.1.2 Crystallization
The crystallization behavior of semisrystalline materials is another critical factor
that affects the maximum expansion ratio in plastic foam processing. For semicrystalline
poiymers, the polymer melt solidifies at the moment of crystallization during cooling.
Therefore, in the foam processing of polypropylene, the foam structure "freezes" at the
crystallization temperature during the foaming process. If the crystallization occurs in the
primitive stage of foaming, i.e., before the dissolved blowing agent fully diffused out of
the plastic matrix and in the nucIeated cells, then the foam cannot be fully expanded.
Therefore, in order to achieve the maximum volume expansion ratio from the
polypropylene foarn, the crystallization (or solidification) should not occur before al1 the
dissolved gas diffuses out into the nucleated cetls. Upon exiting the die, the temperature
of melt decreases due to external cooling outside the die and the cooling effect due to
isentropic expansion of the blowing gases. Thus, the processing temperature at the die
determines the time after which the polymer melt solidifies. Therefore, in order to give
enough time for the gas to diffuse into the polymer matrix. the processing temperature
should be high enough. U should be noted that if the processing temperature is too close
to the crystallization temperature, the poIyrner melt would be solidified too quickly
before the foam is expanded fully, as shown in the initial section of Figure 6.2.
On the other hand, if the tempenture is too high, then the solidification time
might be too long and the gas that has difised out of the plastic melt to the nucleated
cells might escape out of the foam as discussed above. This indicates that there is an
optimum processing temperature for achieving maximum expansion as shown in the
middle section of Figure 6.2. if the melt temperature (Le., the processing tempenture) is
too high, then the maximum volume expansion ratio is governed by gas loss and the
volume expansion ratio will increase as the processing temperature decreases. if the
processing temperature is too Iow, then the volume expansion ratio is govemed by the
soIidification (Le., the crystailization) of polypropylene, and the volume expansion ratio
will increase as the temperature increases.
In addition to the effects of the processing parameters on the crystailization, the
foaming additives and materials parameters can also contribute to changes in
crystailization temperatures, The effect of these parameters on the volume expansion
ratio is shown in Figure 6.3.
6.2.2 Visualization of Expansion Behavior Using a CCD Carnera
6.2.2.1 System setup
For verifying the mode1 described above, photographie images of the extmdate
were taken and analyzed for each processing and material parameter. The system setup
consisted of the tandem extrusion system already described in Chapter 3. In addition, an
on-line progressive scan imaging system was mounted at the die exit, to capture images
of the extrudate coming out of the die. using a frame gnbber and image processing
software, The progressive scan imaging system consists of a CCD camera (CV MlO),
which has a very high shutter speed up to 1180000 second, with a magnifjing lens
(Navitar), a frarne grabber (PC vision), and image processing softwrire (Sherlock).
6.2.2.2 Experirnental Procedure
The foamed extmdate was monitored while changing the processing temperatures.
From the captured images the angle of initial swelling of foamed polymer and the
extrudate diameter as a function of the distance from the die exit were measuted, The
images of the foamed extrudate were analyzed and the effects of gas ioss and
crystailization were extracted. Some sample images are shown in Figure 6-4. Sample
results for branched polypropylene materid with 15% butane concentration ai hree
different processing temperatures are elucidated in the following sections.
6.2.2.3 Effect of Processing Temperature on the Initial Expansion Rate and Final
Diameter of the Extudate
Figure 6.5 shows the effect of processing temperatures on the initial volume
expansion behavior characterized by the initiai angle 0 of the extrudate. The figure shows
that at the higher temperatures, the initial expansion rate, and consequently 8, was quite
significant. However, since gas Ioss was accelerated at this high temperature, the final
diameter of the extnided foam was smdl. As the temperature decreased. the initial
expansion rate (8) was decreased and the final diameter of the extnided foams was
increased due to the decreased gas loss amount. There existed an optimum temperature
to achieve the maximum diameter of the extrudate. When the temperature was further
decreased, and because of the earIy crystallization, the foamed extrudate was frozen
before the extrudate was fully expanded. As a consequence, the final diameter d was
small. The results are summarized in Figure 6.6.
Figure 6.7 shows the changes occurring in the extrudate diameter, as the distance
from the die exit increased, At the optimum die temperature of 130°C, the extrudate
diameter reached its maximum value of 5 mm at a distance of 5 mm away from the die
exit, At this point the foam extrudate reached the crystallization temperature and retained
its shape due to solidification of the polymer. At the lower temperature of 120°C, the
polymer reached the crystallization tempenture earlier, about 3 mm away from the die
exit and freezed at this point. Thus, at the same gas content of 15%, the maximum
diameter achieved was about 4.6 mm instead of 5 mm which would result in sacrifice of
the expansion ratio. in the meanwhiIe, at a higher temperature of 200°C, the effect of gas
loss was more pronounced. First of dl, the gas at a higher temperature caused a sharper
initia1 expansion ratio and the rate of diameter growth was greater. Simultaneously, the
gas loss due to the increased diffusivity of the gas resuited in a lowered maximum
diameter in the extrudate, which was about 3.6 mm and was observed about 2.5 mm away
from the die exit. As the temperature of the extrudate was stiIl above the crystdlization
temperature of the polymer, it continued to experience hrther gas loss and the diameter
continued to decrease beyond the range of the photographic equipment. This decrease in
diameter could attnbuted either to contraction due to cooling or due to additional gas loss
by diffusion through the foam skin.
6.3 Theoretical Mode1 for Calculating the Expansion Ratio, Cell Size, and Cell Wall
Thickness from the Observed Foam Profile
6.3.1 Development of a Theoretical Mode1
A theoretical mode1 was developed for describing the expansion ratio in terms of
distance from the die exit at each processing temperature. The rnodel takes into
consideration gris loss at high temperatures and solidification due to the early
crystallization at low temperatures.
In order to determine the instantaneous diameter of the extrudate upon leaving the
die exit, a first order mode1 based on the conservation of mass was considered. The
continuity equation rnay be written in the following form:
riz = p Av =constam. (6.1)
where m is the rnass flow rate, p is density, v is speed, of a polymerlgas solution
expenencing flow through the channel having cross-sectional area A (see Figure 6.8). By
definingx-axis as the flow direction, Eq. 6.1 at any given point becomes
Px Axvx = mpr * (6.2)
where m,, is the combined flow rate of both polymer melt and gas. Since
Eq. 6.2 becomes
where @, is the volume expansion ratio at the distance x, ppx = p, = the polymerfgas
solution density, p, is the polymer melt density, and d, is the extrudate diameter.
Assumingv, = v p = constant at any point in the cross section at the distance x, the
continuity equation for polyrner melt is:
p p Apvp = f i p , (6.5)
where m, is the flow rate of polymer melt. and Ap is the polymer area which is a
hnction of the ce11 shape. The ce11 shape depends on the processing condition as follow:
Case 1: when x is very small, the ceIls are not fully grown and are generally of
spherical shape, separated from each other and uniformly distributed (Figure 6.8a).
Case 2: when x is at the maximum expansion point (Figure 6.8b). The cells are
hlly-grown and are polyhedral in cross-section with thin polymer walls separating them.
As a first approximation for this model, it is assurned that the cells are cubic in shape.
Case 3: when x is large, there are two possible cases. First, the polymer reaches
the crystallization tempenture before x became large, and freezes in the state that it was
in corresponding to either case 1 or case 2. Second, the temperature was above the
crystallization limit. In this case the gas will diffuse out through the thin ce11 walls and
outer skin and the cells will shrink. in this situation, the approximation of spherical cells
can again be applied and we can use the model described
1s t part of Figure 6.8a.
Modell
If the ce11 is represented by ri sphere of diameter D (Cases
in Case 1. This is shown in the
1 and 3, then
(6.6)
whiIe n is the number of cells per cross-section (see Figure 6.8). Substitution of Ap into
Eq. 6.5 yields the following:
The volume expansion ratio, 4,. is defined as
where V, and V, are the occupied volume of gas and polymer melt in the foarn.
respectively. Since the ceil density is defined to be the number of cells per unit volume of
polymer:
Combination of Eq.'s 6.9 and 6.10 results in:
Using Eq. 6. i 1, Eq. 6.8 is then rransfotmed into:
Since the flow is assumed unidirectionai. and v, = v , =v,, Eq. 6.1 1 may be substitured
into Eq. 6.4:
when $, is too Iarge, equation 6.14 becomes
If the ceIl is represented by a L-sized cube (Case 2), then
Eq, 6.6 could be rewntten as:
and therefore
For a cubical cell, Eq. 6. IO becomes:
Keeping the definition introduced in Eq. 6.9, an equation for L may be derived:
Combined with Eq. 6.1 8, this equation yields:
Then combination with Eq. 6.4 produces the following formula:
when 4,is too large, equation 6.23 becomes
Figures 6.9 (a) and 6.9 (b) show the calculated values of the volume expansion
ratio using Models 1 and 2, respectively, for a fixed ceIl density (N) of 8.7 x106 cells/cm3
while varying the number of cells per cross sectional area (n) from 200 to 800. It can be
seen that the expansion ratios predicted by both models are close to each other. However,
Model 1 would show yet better prediction in the low expansion region in which the cells
have not met each other and therefore maintain a spherical shape. On the other hand,
Model 3 would show better prediction in the high expansion region where the ce11 shape
can not be spherical but rather polyhedral due to the contact of the cells. The predicted
average ceIl sizes based on both rnodek are also close to each other within 25% (see
Figure 6.10). Since Model 1 predicts a negative cell wall thickness for the large
expansion region (see Figure 6.1 1), Model 2 seems to be a better model overall to
describe the expansion ratio, cell size, and cell wall thickness. Therefore, Model 2 has
been selected to estimate the expansion ratio, ce11 size, and cell wall thickness of the
extmded polypropylene foams from the observed extrudate profile as shown in the next
section. However, it should be mentioned that this model is based on the observed outer
shripe of the extrudate, the finally observed number of cells in the cross-section, and the
ce11 density of obtained foam. in order to predict the shape of the extrudate as a function
of time, a set of equations that describe the dynamics of the polymer-gas system
including the flow of polymer melt, diffusion of gas, and heat transfer needs to be solved
simultaneously- Further research is required for this work.
6.3.2 Determination of Expansion Ratio, Ce11 Size and Cell Wall Thickness of
Extruded Foams from the Observed Profiles
Figures 6.12, 6.13, and 6.14 show the calculated volume expansion ratio, ceIl size
and celt wall thickness of the extmded foarn, respectively, as a function of the distance
from the die exit at the melt temperature of 120°C. 130°C, and 200°C. Table 6.1 shows
the number of cells per cross section area and the ce11 density for each temperature.
Figure 6.12 shows that for the lower temperatures of 120°C and 130°C, the expansion
ratio initiaily increased rather rapidly with increasing extrudate diameter. This trend
continued up to 60-fold expansion. At the point of maximum diameter, rit a distance of 5
mm away from the die exit, the maximum volume expansion predicted was 90-fold, for
the optimum processing temperature of 130°C. On the other hand, at the processing
temperature of 120°C, the polymer reached the crystallization ternperature before
attaining the maximum expansion, and consequently the achieved maximum expansion
ratio was only about 78-fold. At the higher temperature of 30O0C, the plot shows that the
volume expansion ratio was very low. The excessive gas loss, due to increased gas
diffusion at the higher temperature resulted in only a very low maximum volume
expansion ratio of 13-fold. On the other hand, it is worth of noting that the ceil sizes of
120°C and L3O0C were very close to each other (Figure 6-13), whereas the cell wall
thickness were quite different (Figure 6.14). This demonstrates chat the lower temperature
prohibited the ceII walls from being seretched and thinned during expansion becnuse of
the eariy crystailization,
6.4 Summary and Conclusions
The fundamental mechanisms governing the volume expansion behavior of
polypropylene foams were determined based on the experimentai results shown in
Chapter 5. Tt turned out that either gas loss or polymer crystallization governs the
expansion behavior of polypropylene foams. A progressive image scanning setup was
configured to capture the images of the foamed extrudace coming out of the die on a PC.
The images captured were anaIyzed and the data were used to verify the proposed
mechanisms, A theoreticai mode1 h a been proposed, which relates the instantaneous
expansion ratio, ce11 size, and cell wai1 thickness to the exwdate diameter as the distance
from the die exit varies. The developed mode1 was effectively used to describe the
variation of the cell morphology in the extruded foarn from captured images. However,
the mode1 used the experimentaily observed parameters such as the ceil density, the
number of cells per extrudate cross-sectiond area, and the final volume expansion ratio.
Further studies are required to describe the gas Ioss phenomena over tirne.
Tabie 6.1 Observed number of cells per cross section at various temperatures
I (Cl 1 cross secrion (n) 1 IN) I Die Temperature Number of celIs per Cd1 Density
Crystallizalion ,,, Gus Loss 4-- >
Figure 6.2 Effect of gas Loss and crystallization on the volume expansion
Crystallization by gas loss I
Temperature
Figure 6.3 Fundamental volume expansion mechanism of polypropylene Foarns
Figure 6.4 images of the foarn extrudate coming out of the die
Figure 6.5 Effect of processing temperature on the initial expansion
Ternparatum (C)
Figure 6.6 Effect of Processing temperature on the initial diameter
0.00 2.00 4.00 6.00 8.00 10.00
Oistanci tom Oia (mm)
Figure 6.7 Extrudate diameter as a function of the distance from the die
die at a
Figure 6.8 Description of ce11 shape models
O ! t
O 1 2 3 4 5 6
Eztnrdate Dkmibr (mm)
6.9 (a) The calculated volume expansion ratio based on Model 1 for ceII density = 8 . 7 ~ IO6 ceIls/cm
0 ! O 1 2 3 4 5 6
Extrudita Dhmimr (mm)
6.9 (b) The caiculated volume expansion ratio based on Model 2 for ce11 density = 8 . 7 ~ 106 cells/cm
O 5 1 O 15 20 25 30 35 40
Expanrlan Rafio
Figure 6.10 Calculated average ceIl size for for cell density = 8.7~ IO6 cells/cm
Figure 6.1 1 Calculated ce11 wail thickness for ce11 density = 8.7~ IO6 ceIls/cm
1°- 0.00 ff . 0.00 200 4.00 6.00 8.00 10.00
Distance from Die (mm)
Figure 6.12 Calculated expansion ratios of extnided foams from the observed profiles
0.00 2.00 4.00 6.00 8.00 10.00
Distance (rom Die (mm)
Figure 6.13 CdcuIated average ce11 size of extmded foams from the observed profiles
0.00 200 4.00 6.00 8.00 10.00
Distanci fmm O k (mm)
Figure 6.14 Calculated ce11 wall thickness ofextnrded foams from the observed profiles
Chapter 7
Summary and Conclusions
7.1 Sumrnary
A continuous extrusion foaming process has been studied for the manufacture of large
volume expansion ratio with a fine-celled structure in extruded polypropylene foams. A
tandem-extrusion foaming system was designed and anaiyzed based on the axiomatic design
approach. Experiments were performed to check the functionaiity of the designed system.
The strategies for promoting ultra low-density polypropylene foarns were developed.
The effects of processing and materials parameters, including those of the processing
temperature, arnount of blowing agent, amount of nucleating agent, blending of blowing
agents, long-chain branching of polymer, blending of linear and branched materids, and die
dimensions on the finai foam properties were investigated in detail.
The fundamental mechanisms for the promotion of high volume expansion ratio of
polypropylene foams were elucidated. A qualitative modeling for the volume expansion
behavior was described at various processing temperatures A theoretical model for
calculating the volume expansion ratio, average ce11 size and ceIl wall thickness was
developed. This model was used to describe the volume expansion behavior at various
distances from the die using the observed CCD images of extnided foams.
Basic studies were also cam'ed out co investigate the effects of dissolved gas on the
PVT data, crystallization kinetics and meIt fracture behaviors of polypropylenelgas soIutions.
The obtained results were usehl in the understanding of the processing technology that
govems the production of low-density, fine-celled polypropylene foams.
7.2 Conclusions
The experimental work presented in this thesis leads to the following conclusions:
1. A tandem extrusion system for the production of low-density, fine-celled polypropylene
foams was designed and constmcted. This system will considerably enhance the
potentid to scale up the existing system into an industrial production system.
3. The basic strategies employed for the promotion of a large volume expansion ratio with
polypropylene materials were fourfold: to use a branched material for preventing ce11
coalescence; to use a long-chah blowing agent with low diffusivity; to lower the meIt
temperature for decreasing gas loss during expansion; and to optimize the processing
conditions in the die for avoiding too quick crystdlization.
3. Low-density, fine-celled polypropylene foams were successfully produced using the
designed system. The branched polypropylene foams have a maximum volume
expansion ratio in the range of 90 times, and a ce11 density of higher than 1o6 ce11s/cm3.
The results show the effectiveness of the fundamental strategies adopted to promote
large expansion.
4. The effect. of dissolved butane on the PVT relationships of linear and branched
propylene materials in a molten state were investigated. The specific volume of the
propylenehutane solution increased significantly with an increase in the percentage of
butane injected in both branched and linear propylene. At al1 experimental conditions
selected, the specific volume was found to be higher for the linear propylene than for the
branched propylene. When butane was dissolved in the propylene matrix, the sensitivity
of the specific volume with respect to pressure increased with the butane content for
both the linear and branched resins, whereas there was no significant changes in the
sensitivity with respect to the temperature.
5. A series of experiments were conducted to investigate the effects of materid branching,
foa ing additives, cooling rate, hydraulic pressure, and dissolved gas on the
crystallization behaviocs of potypropylene resins. Branching and foaming additives in
the polypropylene matrix caused a significant increase in the crystailization temperature.
The crystallization temperature was a sensitive function of the cooling rate and it
decreased as the cooling rate increased. Crystdlization of polypropylene materials was
enhanced as the hydraulic pressure increased. But the dissolved N2 and CO1 lowered the
crystailization temperatures of polypropylene resins.
Experirnental studies were carried out to investigate the effects of branching, processing
temperature, foaming additives, and blowing agent on the critical shear stresses of linear
and branched polypropylene resins. An on-line technique for detecting the onset of
surface melt fracture for extruded foam has been developed by visualization of the
extrudate using a CCD carnera. The long-chain branching of polypropylene matends
significantly decreased the critical shear stress of the resins. The critical shear stress was
insensitive to the die ternperature; however, the die temperature signitlcantly riffected
the critical shear rate at the onset of surface melt fracture. The foaming additives of talc
and GMS increased the critical shear stress of branched polypropylene rnaterials more
than that in linear rnaterials. The dissolved butane significantly increased the critical
shear stress'es of linear and branched polypropylene resins.
The effects of processing parameters such as the temperature, the materials parameters
such as the blowing, nucleating agents, and long chain branching, and the die geometry
on the foam density and ceIl density of polypropylene foarns were investigated.
Despite its high flamrnability, butane is a very effective blowing agent for low-density
polypropylene foams cornpared to COz. A higher ceIl density and lower volume
expansion was obtained by increasing the amount of CO2 when using a blend of CO2
ruid butane as blowing agent.
There exists an optimum temperature for achieving the maximum expansion ratio of
polypropylene foam with butane. This optimum ternperature decreased as the amount of
blowing agent increased because of the plasticizing effects.
IO. The cell dènsity increased proportionally when the amount of taIc increased frorn O to
2.4 5%- However, there was an optimum talc amount for achieving the maximum volume
expansion ratio.
11. The expansion ratio obtained from the linear polypropylene materials was rnuch lower
than that from the bnnched polypropylene materiais because of severe ce11 coalescence,
The expetimental results indicate that branched polypropylene rnatcrials are effective for
lowdensity foam application because of the reduced degree of ceII coalescence.
12. The use of a blend of linear and branched polypropylene materials revealed that the ce11
density in the foamed material increased as the amount of branched polypropylene was
increased. Despite that there was no clear goveming mechanisms observed for the ce11
density of linear and branched polypropylene blends, the ceil coalescence could
obviously be reduced by the addition of branched HMS material, giving the blend higher
melt strength and melt extensibility. On the other hand, the volume expansion increased
when the amount of branched rnaterial in the blend increased in the same way as the
previous phenomena. The optimum Lempenture for producing the maximum volume
expansion ratio decreased as the branched material content was increased. This is
believed to be due to the high viscosity of the linear propylene together with the
extensional characteristics of the added HMS component.
13. The use of re-extruded branched materials resulted in a lower ce11 density and a lower
volume expansion ratio- This was attributed to the breaking of the long-chain branching
of the high melt strength polypropylene, which led to ceIl colilescence in the foam
structure.
14. The maximum volume expansion ratio was decreased using a lower pressure drop rate
die. However, the expansion ratio was higher with a lower pressure drop rate die in the
high temperature range due to the reduced gas loss from the foam extrudate. The ceIl
density was higher in the case of higher pressure drop rate as per previous studies
[17,18].
15. A careful anaiysis of extended experimental results obtained at various processing
conditions indicates that the final volume expansion ntio of the extruded polypropylene
foams blown with butane is governed either by loss of blowing agent or by
crystallization of the polymer matrix. The CCD images were anaiyzed to illustrate both
these mechanisms of gas [oss and crystallization during foarning at various temperatures,
and it was observed that the maximum expansion ratio was achieved when the governing
mechanism was changed from one to the other. Therefore, it is highly recomrnended to
look for this transition point to maximize the expansion ntio for the low-density foam
applications.
Recomrnendations and Future Work
The following suggestions are made for the direction of future research on the
production of low-density, fine-celled polypropylene foam:
1. The extrusion process for the manufacture of low-density, fine-celled polypropylene
foarns filaments developed in this thesis cm provide useful insights to extend the
system to sheet extrusion and injection molding processes. Using the processing
conditions determined here, appropriate design strategies can be developed for other
foam processes,
3. This study was confined to branched and Iinear polypropylene homopolymer. It was
found that each material has particular processing conditions. However, other
thermoplastics materials with different properties could be examined using the
process developed to determine suitable processing conditions for various materials.
3. The die design needs to be modified to improve the surface melt fracture of the
exuudate. A number of parameters, including the Iength to diameter ratio, the die
exit angle, and the die material cm be altered to investigate how they influence the
surface melt fracture.
4. The use of a CCD canera in this study was confined to the image rinalysis of the
filament extrudate to describe the fundamental mechanisms for the volume expansion
ratio of polypropylene foams. However, the study of the ce11 growth mechanisms
based on computational mathematical modeling such as development of a finite
element anaiysis simulation software will be beneficial to determine the growth of
bubbles in any processing conditions.
5. Incorpcirating a completely automated computer control for the tandem extrusion
system for the production of low-density fine-celled plastics foams wiI1 be of a great
value in reducing the error bounds for the measured fundamental pmperties of
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Flow in Capillary Die
A schernatic of the flow in a capillary die is given in Figure Al, 1. For the usual
cylindticzii polar coordinate system, there is a complete circular geometry about the z-
mis, i.e., nothing depends on the 0-coordinate [196], Therefore, the 0-component of the
fluid veiocity, v ~ , becomes zero. The assumption that the tube and the fluid extend very
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cornponent of the fluid velocity, v,, is also zero. Considering that the flow is steady, Le.,
and that the density of the incompressible fluid, p, is constant, the equation of continuity
in cyfindrical coordinates [ 1971
Ieads to
Hence, v, does not change with the z-coordinate and the 8-coordinate (because of the
assumed cornpiete axisyrnrnetry of the flow), and is a function of r only:
v: = v: (r) (A 1.4)
This flow is, therefore, also fulIy developed since the velocity profile is the same
everywhere. Now one can apply the Navier-Stokes equations [198] in the r-direction
+n[a [ 1 a ( + - - - - I a l v a 2 a..] al- r ar n>r) c r ~ Je' azz rl 20
and the z-direction
where g is gravity and T\ is the shear viscosity. The Navier-Stokes equations become in
the case of steady flow
where Eqs. A1.7 (a) and (b) represent simplified Eqs. A1.5 and A1.6, respectively.
Applying the shear stress definition [198]
and defining a modified pressure as
P = p +pgh , (A 1.9)
one can rewrite Eq. A 1.7 using the following transformation d e [197]:
Hence Eq. A1.7 c m be simplified to yield
Eq. Al . l l (a) indicates that P is a function of z only (remembering the assumed
syrnmetry). A simple modification of Eq. A 1.1 L (b) can give
in this form, the right-hand side of the equation does not depend on z, while its left-hand
side does not depend on r. Thus, both sides are constant:
For the power law unidirectional flow, satisfying Eq. A1.4, the shear stress is given by
[197]:
av, The absolute value sign may now be dropped because --.- is always negative in tube
&
flow. Substitution of Eq. A1.14 in Eq. A1.13 results in
lntegrating this equation twice leads to the following:
where Ci and Cl are the integration constants. The boundary conditions are (see Figure
Al.1):
Afr=R*vZ =O, (A1.17 a)
av, A t r=O=&=O. (A1.17 b)
ar
Now one can easily resolve a simple system of two equations, established by substitution
of Eqs. A1.17 (a) and (b) in Eq. A1.16, to find expressions for Ci and Cz, which finaily
yield the velocity distribution:
The shear stress can be found by combining Egs. A 1-14 and A1 - 18:
and its maximum value at the wall is
where L is the tube length. The volumetric flow rate is obtained as [196]
R
Q = fv:(r) -2wdr. O
Again, Eq. A 1.18 is substituted to give
It should be noted that the value of hP = 4 - P, is negative due to pressure decrease in
the tube flow, and its absolute vaiue may be calculated from Eq. A1.22:
The shear rate is defined as
Using Eq. A1.18, the wall shear rate (Le., at r = R ) absolute value can be obtained:
The apparent shear rate
represents the wail shear rate for Newtonian fiuids (Le., ai n = 1 ).
Figure AI.1: Flow in ~ h c Fully Drveiopcd Region ofa Circuhr Die
Appendix 2
Statistical A nalysis
Table A.l Independent Variable Settings (Single Constraints)
Variable Name
I I
(UD ratio)
Variabk
Units
Temperature (T)
Gris percentage (Gas %)
Talc percentage (Talc %)
Mass flow rate (M)
Capillary die
RnngeKevel (s)
OC
Gas %
Talc %
g m i n
1205T1180
5 < G a s % ~ 2 0
0.4 1 Talc % 1 2.4
lOIM5 12
Table A.2 Sample Preparation
1 T 1 Gas % Talc % 1 M 1 UD 1
Table A.3 Response Oata
Average Volume Expansion
(Times) 1 .O3
Expetimental Errot Statistics
Table A.4 Overall Error Statistics: Transformed Volume Expansion
Computed Error Statistic
Computed Error Statistic
Name Error %
Regression Statistics
Value 1 .go41 371 21
Experimental Error (It) 95% Confidence Limits (I) Adj. R Square Lower Lirnit
Table AS General Regression Statistics: T ransfoimed Volume Expansion
0.1 88680493 0.5238621 17 0.892833963
Table A6 Regression ANOVA Statistics: Transformed Volume Expansion
Regression -
Statktic Name R Square Adj. R Square
Standard Error (t) Observations Degrees Of Freedom
Lack-Of-Fit Lack-Of-Fit F-Ratio P-value
Regression Regression Statistic Value Target Name 0.974834597 Adj. R Square Upper Lirnit 0.9635101 66 Adj. R Square Mean
Adj. R Square Lower Lirnit 0.261 194306 Standard Error (3
Source Of Vanàtion
Regression Residual Total
Regression Target Value 0.99197221 1 0.980958629 0.892833963 0.188680493
30
16 2.1 45431 1 08 0.2402601 64
Degrees Freedom
- 9 20 29
Sum Of Squares
52.85483451 1.3644501 49 54.21 928466
Mean Square
5.87275939 0.068222507
F- Ratio 86.08243246
P-value 5.1 396E-14
Coefficients Table And Models
Table A.7 Model Coefficients: Transformed Volume Expansion
Coded Vanable Name
lntercept XI X2 X5
l2 (m2 XI *X2 X2'X4 x3*x4 X3*X5
- - -
Coefficient Value
3.945002965 -1 .O35280948 0.759955449 -0.1 3971 333 -1.3657491 n -0.898630631 -0.6516091 9 0.269371 443 -0.31 6477941 0.247939966
Coded Variable Name Key XI =T
Coefficient Standard
Limit 3.7421 37807 -1 -1 7205327 0.61 8222707 -0.248761 021 -1.606254766 -1.1 52489506 -0.824566242 0.1 1870523
-0.46173851 7 0.1 00995968
Upper 95% Confidence
iimit 4.1 470681 23 -0.898508626 0.90168819
-0.030665639 -1.1 25243588 -0.644771 757 -0.4786521 38 0.420037655 -0.1 71 21 7366 0.394883965
Coefficient Lower 95% Confidence
Model Term Ranking
Table A8 Model Term Ranking: Transformed VoIume Expansion
Model Term Name
X 1 XS
(X1 l2 X1 'X2 (x2)2
X3'X4 XTX4 X3'XS
X5
T ransformed Coded Variabk Model 1 ransforrned Natural Variable Model
Model Term Range
2 2 1 2
0.888888889 2 2 2 2
Coefficent Value
-1 .O35280948 0.759955449 -1.3657491 n -0.651 6091 9 -0.898630631 -0.31 6477941 0.269371 443 0.247939966 -0.1 3971 333
Volume Expansion = 3.945003 -1 .O35281 ' (Xt ) +0.75996 * (X2) -0.1 3971 3 (X5) -1.365749 ' (XI2) -0.898631 * (XZ2) -0.651609 ' (XI 'X2) +0.26937 ' (XZX4) -0.31 6478 ' (X3'X4) +0.24794 ' (X335)
Volume Expansion = -32.861 18 +0.45694 ' (T) +0.54005 ' (Gas %) +2.82463 ' (Talc
-0.005883 ' (M) -0.1 069% ' (UD) -0.001517 - ( 0 2 ) -0 .O1 5976 ((Gas %)2)
-0.002896 ' (T ' Gas %) +0.03592 ' (Gas % ' M) -0.31 6478 * (Taic % ' M) +O.OS449 ' (TaIc % ' UD)
Mode/ Term Effect
-2.070561 897 1 .51 991 0898 -1.3657491 n -1 30321 838 -0.798782783 -0.632955883 0.538742805 0.495879933 -0.27942666
Mode1 Term Rank 1 .O0 0.73 0.66 0.63 0.39 0.31 0.26 0.24 0.1 3
Model Term Ranking And Charts
Table A.9 Model Term Ranking: Transfoned Volume Expansion
Mode! T e n ( Mode! T e n ( Coefficient ( Range
2 2 1 2
0.888888889 2 2 2
Value TiziGG-1
Mode/ T e n Effect
-2.070561 897 Rank 1 .O0
Model Term Ranking Pareto Chart I
Model Term Ranking Pie Chart ):pI
Residuals Data And Plots
Table A.10 Residuals Data: Transformed Volume Expansion
Sample No. 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30
Transfomed Obsewed
Volume Expansion Predkted
Volume Expansion 0.457487276 4.31 9647868 1.247043736 3.084571 803 4.01 151 8226 3.658506099 4.01 151 8226 0.869579671 2.764039632 2.980492905 3.029741 423 3.86744792
0.779537644 3.348625228 3.077341 576 3-1 071 26075 0.82671 671 9 3.9721 35609 3.1 94424744 3.745207645 3.1 39594364 3.348625228 0.971755745 4-1 02520974 0.971 755745 0.55549238
0.598355332 3.658506099 4.00452387 4.325664831
Residuals 0.1 97438691 -0.461 131 13
-0.1 6693491 9 -0.1 483241 77 -0.02355321 7 0.3041 14959 0.01 5528456 -0.230360833 0.120481712 -0.01 0334883 0.442069738 0.04447508
0.072747758 0.1 7020703 0.298367043 0.344289023 -0.1 8749788 -0.052243645 -0.397448334 0.1 0270341 5 -0.31 2576644 0.1 552791 52 -0.07771 O622 -0.1 1641 9425 -0.07773 0622 0.1 09769597 0.040863506 -0.22791 1741 -O.lO67Ol214 0.1 78524125
Deleted Studentized Residuals
0.993830921 -2.1 2200639 0.869426698 0.607634754 ,0.104208128 1.51 597702
0.068692524 -1.262804666 0.60472089 ~0.048930984 2.049369029 0.21 9441 534 0.352831 699 0.724460757 1.482632655 1.5661 61 587 -1 .OZ5SOl9l9 .0.253394794 -1 -726274342 0.507268892 -1.342293477 0.659396785 -0.355485298 -0.497878826 -0.355485298 0.523408585 0.21 0352292 -1 -1 O7 140256 -0.482238332 0.848866591