Extrusion Processing for Manufacture of Low-Density, Fine-Celled Polypropylene Foams
~aniE. 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
@Copyrightby B. E. Naguib 2001 National Library Bibliothèque nationale 1*1 .,,da du Canada Acquisitions and Acquisitions et Bibliographie Services senrices bibliographiques 395 WeU'iglon Street 395. nie Wellington Ollawa ON K1A ON4 ûttawa ON K1A ON4 Canada Canada
The author has granted a non- L'auteur a accordé une licence non exclusive licence allowing the exclusive permettant à la National Library of Canada to Bibliothèque nationale du Canada de reproduce, loan, distribute or seii reproduire, prêter, distribuer ou copies of this thesis in microfom, vendre des copies de cette thèse sous paper or elecironic formats. la forme de microfiche/f'ilm, de reproduction sur papier ou sur format électronique.
The author retains ownership of the L'auteur conserve ta propriété du copyright in this thesis. Neither the droit d'auteur qui protège cette thèse. thesis nor substantiai extracts fiom it Ni la thèse ni des exûaits substantiels may be printed or otherwise de celle-ci ne doivent être imprimes reproduced without the author's ou autrement reproduits sans son pemksion. autorisation. 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 pst 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
II O LORD. yorr have searched me and you knoiv me. YOUknoiv ivhen C sir and ivhen I rise; You perceive my rfroirgl~tsfrom 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 loftyfor me to airain. Where can I go fiom pur Spirit? Wliere can Iflreji-om yuurpresence? IfIp 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 ivonderfulfymade; 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 wiftenin 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 1.2 TherrnopIastics Foarns ...... 7- 1.3 Low-Density, Fine-Celled Foam Processing ...... 3- 1A 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 froperties ...... 13 3.2.2 Crystallization* * Behavior ...... 14 2.2.3 Melt Fracture ...... 16 2.3 Theoretical Background ...... 18 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 ...... 41 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
vii 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 55 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 Camera ...... 143 6.2.2.1 System Setup ...... 143 6-2-22Expenmental 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 Propenies of the polypmpylene resins used in determining the onset of melt fracture ...... 76
Tabie 5.1 Properties of the poiypropylene resins used in the foaming experiments ...... 116
Table 6.1 Observed number of celis per cross section at various temperatures ...... 151
Table A .L Independent Variable Settings (Single Constraints) ...... 173
Table A.2 Sample Prepararion Sheet ...... 174
Table A.3 Response Data ...... 175
Table AA Overall Error Statistics: Tnnsformed Volume Expansion ...... 176
Table AS GeneraI Regression Statistics: Tnnsformed Volume Expansion ...... 176
Table A.6 Regression ANOVA Statistics: Transformed Volume Expansion ...... 176
Table A.7 Model Coefficients: Trmsformed VoIume Expansion ...... 177
Table A.8 Model Terrn Ranking: Transfonned Volume Expansion ...... 177
Table A.9 Mode[ Term Ranking: Transforrned Volume Expansion ...... 178
Table A .t O Residuals Data: Transformed Volume Expansion ...... 179 LISTOF FIGURES
Figure 2.1 Solubility of Gas into polymers ...... 34
Figure 2.2 Effect of the variation ofenergy on the bubble growth ...... 34
Figure 2.3 Model of a nucleated cell inside a polymer matrix ...... 35
Figure 3.1 A schematic of the tandem extrusion system ...... 48
Figure 3.2 A schematic of the rnixing section ...... 48
Figure 3.3. A schematic of the gear pump ...... 49
Figure 3.4 A schematic of the heat Exchanger ...... 50
Figure 3.5 A schematic of the filment die ...... 50
Figure 3.6 A schematic of a cooling sleeve ...... 51
Figure 4 .L Experimental setup for the PVT measurements of Polypropylene/butane solutions ...... 77
Figure 4.2 Degassing oven for the PVT measurements ...... 78
Figure 4.3 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) Effect of the processing temperature on the specific volume for linear polypropylene materials ...... 79
Figure 4.4.(c) Effect of the processing pressure on the specific volume for linear polypropylene materials ...... 80
Figure 4.5.(a) Effect of the dissolved butane on the specific volume for branched polypropylene materials ...... 8 1
Figure 4.5.(b) Effect of the processing temperature on the specific volume
for branched polypropylene materiaIs ...... ,.. 8 1
Figure 4.5.(c) Effect of the processing pressure on the specific volume for branched polypropylene mitterials ...... 82
Figure 4.6 Design of cooling system for the high-pressure DSC ceIl ...... 83
Figure 4.7 DSC thermograms for linear and branched polypropylene materials ...... 83
Figure 4.8.(a) Effect of talc on the crystallization behaviors of linear and branched polypropylene resins ...... 84
Figure 4.8.(b) Effect of talc on the degrees of crystdlinity of linear and branched polypropylene resins ...... 84
Figure 4.9.(a) Effect of GMS on the crystalIization behaviors of iinear and branched polypropylene resins ...... 85
Figure 4.9.(b) 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 wenta ...... 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 CO2as blowing agent ...... 123
Figure 5.5.(b) The blowing agent efficiency versus the melt temperature for branched poIypropylene materials using COlas blowing agent ...... 123
Figure 5.54~) The cell density versus the melt temperature for bnnched polypropylene materials using COras 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 CO1as
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) The die pressure versus the melt temperature for branched polypropylene materials using butane as blowing agent for various talc contents ...... 128
Figure 5.8 The melt strength and melt extensibility of linear and branched poIypropyIene materials biend ...... 139
Figure 5.944 The expansion ratio versus the melt temperature for linear and branched polypropylene materials blend using butane as blowing agent ...... , 130
Figure 5.9.(b) The blowing agent efficiency versus the melt temperature for linear and branched polypropylene materials blend using butane as blowing agent ...... 130
Figurc 5.94~) The ce11 density versus the melt temperature for Iinear and branched polypropyiene materials blend using butane as blowing agent ...... 13 1
Figure 5.9.(d) The ce11 density versus linear polypropylene percentage for Iinear and branched polypropylene materials blend using butane as blowing agent...... 13 L
Figure 5.9.(e) 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 ~olypropyIenernateriaIs 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 polypropylene materiais using CO2as blowing agent for die B ...... 139
Figure 6.1 Effect of processing temperature on the extrudate shape ...... 152
Figure 6.2 Effect of gris Ioss and crystallization on the volume expansion ...... 153
Figure 6.3 Fundamental volume expansion mechanism of polypropylene
foams ...... ,...... 154
Figure 6.4 Images of the foarn extrudate coming out of the die ...... 155
Figure 6.5 Effect of processing temperature on the initial expansion rate ...... 156
Figure 6.6 Effect of processing temperature on the maximum and final diameters ...... 156
Figure 6.7 Extrudate diameter as a function of the distance frorn the die ...... 157
Figure 6.8 Description of ceIl shape mode1 ...... 158
Figure 6.9 (a) The caicuiated volume expansion ratio based on Mode1 1...... 159
Figure 6.9 (b) The caiculated volume expansion ratio based on Model 2 ...... 159
Figure 6.10 Caiculated average ce11 size ...... 160
Figure 6 .I I 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 = 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 (ds)
= 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)
Y," = 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 cm be manufactured by a variety of processes, depending on the application. TypicaI processing methods include extrusion, injection molding, blow molding, rotations! molding, themoforming, etc. [Il. 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 ha 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 cm 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 cm 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 cm 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 Il4 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 HMS 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 dichIorotetrafluoroethaneas 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~/III~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 2.5~10~poise (at 190" C and a shear rate of LOO0 sec-'). They clairned that this composition of polypropylene resins would enhance the fom 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 PWproperties 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 CO2cm 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 ~luorocacbonelastomer 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 due 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 diffusiondistance (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 [il 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- 1151. 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 affecting the heterogeneous nucleation rüte in a dynarnic system such as extmsion [122, 1231; as the shear force increases, the number of nucleated bubbles increases. Lee [122,123 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. Men 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 [il. 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 ~lationshipbecween 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 ~canningCalorimeter 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) cm 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 AL2should be zero. EIernents AL3 and A14should 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 AZ4is 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 A3Jshould 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, hIshould 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 W1laboratory 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~ingsection 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 cm 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 1GlJI 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 me 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 cm be detennined using the following equation [171]: where I = laminar factor (from the manufacturer), and AP = pressure drop, psi. With a flow rate (Q) oF0.303 m3/min (~.~LcxIo-~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 (40psi). 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 &loCwas 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). cooling oil cooling channel static mixer \ oil pressure transducer seat hot polymer cotd polyrner melt melt- +
high ternphanire thermocoup1e cooiing O-ring seat oil
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 ~oamingAgents
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 [60,611, 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(,,ighZhighP, 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 mus flow rate of the propylendburane solution was dererrnined at high experimentnl temperatures and pressures using Equation 4.3.
- mp+,v(IiighT. hjgh PI - mplhighT. 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 mas 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~.injcctionP)
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 cm then be determined by Equation 3, using the sepmteIy measured gas and polymer mas 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 hn 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: Pm-TAY1- 4w-
4.2.1.2 Experimenîai Materiais A linear and a branched polypropylene material with MFRs (ISO 1133, 230 OC/2.16 kg) of Il 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 gar 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 OCthroughout 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 OC, and 15% butane content. The result is depicted graphically in Figure 4.3. It was observed that a npid decrease in the mmoccurred 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 of 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~butanesolution. 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-~crn3/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 OC. 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 ol 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 OC 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 CO2on 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 CO2pressures 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 fee 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 foaing 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. IO 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 (MW, 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~irfacemelt 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 fom 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 - k[FR MW Mn LCB degree Material (gllornin) (Wmol) (KgImol) (n/lûûûc)
Branched P 1 2.3 418 3 1.7 0.2 1
Branched PZ 4.8 416 48 0.17
Branched P3 3.1 327 46.5 O. 13
Linear P 1 II 360 70 O
Linear P2 2.8 500 97 O I I I I 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
I 85 123456789tO 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 Pl (onst) +ûranched Pl (peak) +Linear Pl (onal) +Linear Pl (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 Pl (onsht) . -a-Branched Pl (ptak)
+Linear Pl (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 ioo 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 Pl (onat) +Branched Pl (peak) :
I i -e- Linear Pl (onmt) 1 +Linear Pl (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 Pl (peak) - --I --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 is 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 fom 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 min and high meIt strensth (HMS) branched polypropylene min. 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 WC % 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 las chan 7 pm was used as the nucieating agent. The talc was supplied by the in-house facilities of Bo~aIisAG.
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 ~he 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 (ex 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 OC, 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 of 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 IO 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 CO2acts 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 cm 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 fom 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, 2x10~~ 5x106, and 8x10~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 consequent1y, 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 Pl 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 NOVA 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. Il. 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 OC)
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 Timpiniun (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 Wn(?&)
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 tOdXTalc 10 -1.6% Talc -cZ.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 OM1 1.5225 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 Pl 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 ioa 120 140 160 180 trmpriuin (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 CO2contents 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 fom 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 wdl 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 fom 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, =vp= constant at any point in the cross section at the distance x, the continuity equation for polyrner melt is:
pp Apvp= fip , (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 in Case 1. This is shown in the 1st part of Figure 6.8a.
Modell
If the ce11 is represented by ri sphere of diameter D (Cases 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 ha 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
Die Temperature Number of celIs per Cd1 Density I (Cl 1 cross secrion (n) 1 IN) I
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 1O 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 Ok(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, foaing 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 palymer/gas solutions. This system will facilitate the control of extruders RPM. pressure and temperature conditions and the amount of blowing agent injected. D. Klernpner, and K. C. Frish, Hnndbook of Polvmeric Foams and Foam Technolow, Hanser, N. Y. (1991). C. Maier and T. Calafut, Polypropylene: The Definitive User's Guide and Data book, Norwitch, N.Y. (1998). R. D. Leaversuch, Modem Plastics (1996). C. B. Park and L. K. Cheung 'X Stirdy of Ce11 Niicleation in the Ertrusion of Polypropylene Foams, " Polym. Eng. Sci., 37, 1 ( 1997). J. Throne, Thenno~lasticFoams, Sherwood Publishers, Ohio (1996). V. Kurnar and J E, Weller, in: Polvmeric Foams, K. C. Khernani, Amencan Chernical Society, Washington, DC, 10 1 ( 1997). A. Nojiri, T, Sawasaki, 1. Noguchi, S. Onobori and J. Sasazirna, US Patent 4421867 (1983) C. Jacob and S. K. Dey, "lnert Gases as Alternative Blowing Agnets for Extruded Low Density Polystyrene Foarns," J. Cell. Plast., 31,38 (1995). S. K. Dey, P. Natarajan, and M. J. Xanthos, "Use of Inert Gases in Extruded Medium Density PoIypropylene Foarns," J. Vinyl Add. Tech., 2,339 (1996). T. Shimano, K. Orimo, S. Yamamoto, and M. Azurna, US Patent 398 1649 (1976). K. Orimo, M. Azurna, T. Shimano, and S. Yamamoto, US Patent 3988404 (1976). J. Martini-Vvedensky, N. P. Suh, and F. Waldman, US Patent 4473665 (1984). J. Colton, and N. P. Suh, US Patent 5 160674 (1992). S. Cha, N. P. Suh, D. F. Baldwin, and C. B. Park, US Patent 5 158986 (1992). D. F. Baldwin, N. P. Suh, C. B. Park, and S. Cha, US Patent 5334356 (1994). C. B. Park, N. P. Suh, and D. F. Baldwin, US Patent 5866053 (1999). C. B. Park, D. F. Baldwin, and N. P. Suh, "Effect of the Pressure Drop Rate on Cell Nucleation in Continuous Processing of Microcellular Polyrners," Polym. Eng. Sci., 35,432 (1995). C. B. Park and N. P. Suh, "Filarnentary Extrusion of Microcellular Polyrners Using a Rapid Decompressive Element," Polym. Eng. Sci, 36,34 (1996). D. F. Baldwin, C. B. Park and N. P Suh, "An Extrusion System for the Processing of Microceltular Polymer Sheets: Shaping and Ce11 Growth Control," Polym. Eng. Sci., 36, 1425 (1996). C. B. Park, A. M. Behravesh, and R. D. Venter, "Low-Density, Microcellular Foam Processing in Extrusion Using C02," Polym, Eng. Sci., 38, 18 12 (1998). A. H. Behravesh, C. B. Park, and R. D. Venter, "Challenge to the Production of Low- Density, Fine-Cell HDPE Foams Using C02," Celldar Polym., 17,309 (1998). U. Panzer, Personnel communication, Borealis GrnbH, (1998). D.Y. Kwok, L.K. Cheung, C.B. Park, and A.W. Neumann, "Study on the Surface Tensions of Polymer Melts Using Axisymmetic Drop Shape Analysis," Polym, Eng. Sci., 38,757 (1998). J, Mulrooney, M.A.Sc. Thesis, University of Toronto (1995). M. Lee, C.B. Park, and C. Tzoganakis, "Measurements and Modeling of PSISupercntical CO2 Solution Viscosities," Polym. Eng. Sci., 39.99 (1999). M. Lee, C.B. Park, C. Tzoganakis, and H. E. Naguib, "Cornparison of Various Viscosity Models for PSICO2 Solutions," SPE ANTEC Technicd Pupers, 45, 1241 ( t 999). D. Ladin, CcB. Park, S. S. Park, H+E. Naguib, and S. W. Cha, "Study of Shear and Extensionai Viscosities of Biodegradable PBSICO2 Solutions," J. of Cel. Piast., 37, (2ûû 1). A. Nojiri, T. Sawasaki, and T. Koreeda, U. S. Patent 4,424,293 (1984).
A. Nojiri, T. Sawasaki, and T. Konishi, S. Kudo, and S. Onobori, " Crosslinking of Poiypropylene and Application to Foaming," Frukawa Rev., 2,34 (1982). Y. D. Lee and L. F. Wang, "Properties of Polypropyiene Serucnirai Foams Crosslinked by Vinyletrirnethoxy Silane," J. AppL Polym. Sci., 32,4639 (1986). S- Kitagawa, T. Nakayama, and M. Isono, U. S. Patent 4,766,159 (1987)- K. Feichtinger, U, S. Patent 5929 129 (1999). M. B. Bradley, and E. M. Phillips, "Novel Polypropylene for Foaming on Conventiond Equipment," SPEANTEC Tech. Papers, 36,717 (1990). M. Van Caister, " High Melt Strength PP Foams for the Automotive and Packaging Industries," Foamplas 97, 149 (1997). V. P. Bavaro, "A New Farnily of High Melt Strength Polypropylene Copoiymers for Extruded Low-Density Foam Applications," Sch. and Polyolefins Conf., (1995). U. Panzer, "Development of High MeIt Strength Polypropylene Foams," SPO 98, Houston, Texas (1998). M. Ratzsch, U. Panzer, A. Hesse, and H. Bucka, "Developments in High Melt Strength PP: Technology, Properties, Applications and Markets," SPE ANTEC Tecli. Papers, 45,207 1 ( 1999). J. Thoen, I. Zhao, P. Ansema, and K. Hughes, "New Developments in inspire Performance Polymers for Foarns," Foams 2000, Parsippany, New Jersey, 5 (2000). P. Park, "Unique Reactor Made High Melt Strength PP Resins for Microcellular Foamed Applications," Foamplas 2000,35 (2000). R. Holland and J. Fellers, U. S. Patent 3,996.17 1 (1975). N. Fukushima, Y. Kitagriwa, T, Okurnura, and K. Sakakura, U. S. Patent 4,522,955 (1982). J. J. Park, L Katz, and N. G. Gaylord, U. S. Patent 5,149,579 (1991). J. J* Park, L. Katz, and N. G. Gaylord, U, S. Patent 5.1 16.88 1 (1992). C. P. Park, U. S. Patent 5,527,573 (1994). C. P. Park, U. S. Patent 5,567,742 (1 994). C. P. Park, U. S. Patent 5,348,795 (1994). G. Wilkers, J. Stimler, K. Bly, H. Dunbar, and E. Uhl, U. S. Patent 5,817,705 (1996). E. Andreassen. K. Bore, K. Rommetveit, and K. Redford, "Foaming of PolypropyIene in Extrusion Processes," SPE, ANTEC Technicd Papers, 45,2104, ( 1999). J. Alteepping and J. P. Nebe, U. S. Patent 4,940,736 (1990). A. A. Ahmadi and P. R. Homsby, "Moulding and Characterization Studies with Polypropylene Structural Foarn, Part 1: S tructure-Property Interrelationships," Plus. Rrrbber Proc. Appl., 5,35 (1985). A. A. Ahmadi and P. R. Homsby, "Moulding and Characterization Studies with Polypropylene Structural Foam, Pan 2: The Influence of Processing Conditions on Structure and Properties," Plas. Rrtbber Proc. Appl., 5,51 (1985). J. S. Colton and N. P. Suh, U, S. Patent 4,922,082 (1990).
S. K. Dey, P. Natarajan and M. Xanthos, " Use of inert Gsises in Extruded Medium
Density Polypropylene Foams," SPE MECTech. Papers, 42, 1955 (1996). C. B. Park, L. K. Cheung, and S.-W.Song, "The Effect of Talc on Ce11 Nucleation in Extrusion Foam Processing of Polypropylene with CO2 and isopentane," Cell. Polym., 17,32 1 ( 1998). G.N. Foster, N. Waldman, and R.G. Grisley, Polym, Eng. Sci., 6, 13 1 ( 1966). P. Zoller, J. of Appl. Pobm. Sci., 23, LOS 1- 1056 (1979)- P. Zoiler and D. Walsh, Standard Pressirre-Voltirne-Tem~crnt~dreDaru for Polvniers. Technomic PubIishing Company, Lancaster, PA (1995). Y. Sato, Y. Yamasaki, S. Takishima, and H. Masuoka, " Precise Measurements of PVT of Polypropylene and Polycarbonate," J. Appl. Polym. Sei.. 66, 141-150 ( 1997). T. Hicose, K. Mizoguchi, and Y. Karniya, J. Polym. Sci, Polyrn. Ph-., 24, 7107 ( 1986). S.S. Park, C.B. Park, D. Ladin, H.E. NaguibJ. and C. Tzoganakis, "On-Line Measurement of Poiymer Meir Specijlc Volrime Using a Gear Primp, " SPE, ANTEC Technical Papers, 45,247 1-2478, ( 1999). S. S. Park, C. B. Park, D. Ladin, H. E. Naguib, and C. Tzoganakis, "Measurement of the PVT Properties of a PoIymerGas Solution Using a Foaming Extruder and a Gear Purnp", ASME .l.Manif. Sc!. and Eng., accepted (2000). J. Menczel and J. Varga, " influence of Nucleating Agents on Crystallization of Polypropylene. 1. Talc as a Nucleating Agent, " Thermal Anal., 28, 16 1 (1983). J. 1. Velasco. I. A. de Saja, and A. B. Martinez, "Crystallization Behavior of Polypropylene FilIed with Surface-Modified Talc," J. Appl. Polym. Sci., 61, 125 ( t 996). M. Cook and J. F. Hqr," Influence of Magnesium Hydroxide Morphology on the Crystdlinity and Properties of FilIed Polypropylene," Adv. Pol_vm. Tech., 17,53 (1998). fi. N. Beck, J. Appl. Polym. Sci., 11,673 (1967). K. Mitsuishi, S. Ueno, S. Kodama, and H. Kawasaki, "CrystaIlization Behavior of PolypropyIene Filled with Surface-Modified Calcium Carbonate," J. Appl. Polym. Sci., 43,2043 (199L). Khare, A. Mitra, and S. Radhakrishnan, " Effect of CaC03 on the Crystallization Behaviour of Polypropylene," J. Muter. Sci., 31,569 1 (1996). S. K. Dey, C. Jacob, and J. A. Biesenberger, "Effect of Physical Blowing Agents on Crystallization Temperature of Polymer Melts," SPE MECTech. Papers. 40, 2 197 (1994). Z. Zhang and Y. P. Handa, "In SSu Study of Plasticization of Polymers by high- pressure Gases," J. Polym. Sci., Pt. B: Polym. Phys., 36,977 (1998). J. S. Chiou, J. W. Barlow, and D. R. Paul, "Polymer Crystallization Induced by Sorption of CO2 Gas," J. Appl. Polym. Sci., 30,391 1 (1985). Y. P. Handa, S. Lampron, and M. O'Neill, " On the plasticization of poly (2.6-imethyl phenylene oxide) by C02," J. Polym. Sci., Pt. B: Polym. Phys., 32,2549 ( 1994). Y. P. Handa, P. Kruus, and M. O'Neill, " High-pressure Calorimetric Study of Plasticization of Poly(methy1 methacrylate) by Methane, Ethylene, and Carbon Dioxide," J. Polym. Sci., Pt. B: Polym. Plzys., 34,2635 (1996). Y.P. Handa, 2. Zhang, and B. Wong, "Effect of Cornpressed CO2 on Phase Transitions and Polymorphism in Syndiotactic Polystyrene," Mizcromolecules, 26, 8499 (1997). P. Handa, 2. Zhang, and B. Wong, "Effect of Supercritical CO2 on the Polymorphism in Syndiotactic Polystyrene," SPE ANTEC Tech. Papers, 44, 14 1 1 ( 1998). i. He and P. Zoller, "Crystallization of Polypropylene, Nylon-66 and Poly (ethylene terephthalate) at Pressures to 200 MPa: Kinetics and Characterization of Products," J. Polym. Sci., Pt. B: Polym. Phys., 32, 1049 (1994). 2. Zhang and Y. P. Handa, "CO2-Assisted Melting of Semicrystalline Polymers," Mncromoleciiles, 30,8505 ( 1997). K. Mizoguchi, T Hirose, Y. Naito, and Y. Karniya, "COZInduced Crystallization of Poly(ethy1ene terephthalate)," Polymer, 28, 1298 (1987). E. Bechan and R. S. Porter, "Crystdlization of Bisphenol a Polycarbonate hduced by Supercritical Carbon Dioxide," J. Polym. Sci., Pt. B: Polym. Phys., 25, 15 1 L (1987). R. M. Patel, and J.E. Spmiell, "Crystallization Kinetics During Polymer Processing- Anaiysis of Avaiiable Approaches for Process Modeling," Polym. Eng. Sci., 31, 730 (1991). H. D. Keith, F. J. Padden, J. Applied Phys., 30, 1479 (1959)- J. H, Griffith, B. G. Ranby, 3. Poi-m. Sci., 38, 107 (1959). L. Marker, P. M. Hay, G. P. TiiIey, R. M. Early, O. J. Sweeting, J. Polym. Sci., 38,33 (1959). F. J. Limbert, E. Baer, 3. Polym. Sci., Part A, 1,33 17 (1963). S. Hoshino, E. Meinecke, 1. Powers, R- S. Stein, S. Newman, J. Polym. Sci., Part A, 3,3041 (1965). F. L. Binsbergen, B. G. M. DeLange, Polymer, 11,309 (1970). L. Goldfarb, "Heterogeneous Nucleation of Isotactic Polypropylene," Mukromol. Chem.. 179,2297 (1978). C. C. Lin, 'The Rate of Crystdlization of Poly(Ethy1ene Terephthalate) by Differential Scanning Calorimeter," Polym. Eng. and Sci., 23, 113 (1983). W. Y. Chiu, C, C. Fan-Chiang, K. H. Hsieh, L. W. Chen, "Crystallization of Polypropylene: 1. isothermal Kinetics," Marerials Chemistry and Physics, 34, 46 (1993). G. B. A. Lim, and D. R. Lloyd, "Isothermai Crystallization of Tsotactic Polypropylene in Dotiacontane. 1: Effect of Nucleating Agent Addition on Overall Crystallization Kinetics," Polym. Eng. and Sci., 33.5 13 ( t 993). T. W. Chan, G. D. Shyu, A. 1. Isayev, "Master Curve Approach to Polymer Crystallization Kinetics," SPE ANTEC Papers, 40, 1480 (1994). P. Supaphol, P.J. Phillips, and J.E. Spruiell, "Crystallization Kinetics of High Density Polyethylene at Very High CooIing Rates," SPEANTEC Papers, 43, 1546 (1997). O. Verhoyen, F- Dupret, and R- Legras, "Isothermai and Non-isothermal Crystailization Kinetics of PoIypropyIene Terephthdate: Mathematical ModeIing and Experimental Measurements," Polym. Eng. and Sci, 38, 1594 (1998). P. Supaphol, "Kinetics of kothemal Crystdlization of Syndiotactic Polypropylene: Ami,Tobin, and Maikin Analysis," SPE AN7EC Papers, 46, 163 1 (2000). R. L. Sammler, R. J. Koopmans, M. A. Mangnus, and C. P. Bosnyak. "On the Melt Fracture of Polypropylene," SPE ANTEC Tech. Papers, 44,957 (1998). J. J. Baik and C. Tzoganakis, "A Study of Extrudate Distortion in Controlled Rheology Polypropylenes," Polym. Eng. Sci., 38,274 (1998). 1. B. Kazatchkov, S. G. Hatzikiriakos, and C. W. Stewart, "Melt Fracture Behavior of Molten Polypropylene," Polym. Eng. Sci., 35, 1864 (1 995). S. G. Hatzikiriakos and J. M. Dealy, "The Effect of Interface Conditions on Wall Slip and Melt Fracture of High Density Polyethylene," SPE ANTEC Tech. Papers, 37, 231 1 (1991). S. G. Hatzikiriakos, "The Onset of Wall Slip and Sharkskin Melt Fracture in Capillary Flow," Polym. Eng. Sci., 34, 1414 (1994). D. G. Baird, "Surface Melt Fracture Behavior of Linear Low Density Polyethylene," SPE ANTEC Tech. Papers, 41, 1095 ( 1995). R. S. Jeyaseelan, A. J. Giacomin, P. K. Dhori, and J. C. Slattery, "Polymer Melt instability in Large Amplitude Oscillatory Shear Between Sliding Plates," SPE ANTEC Tech. Papers, 41, 1 106 (1995). A. V. Ramamurthy, "Melt Fracture - Part 1 Capillary Rheometer Studies: Part It Product Fabrication Studies," SPE ANTEC Tech. Papers, 35,717 (1989). E. E. Rosenbaum, S. K. Randa, S. G. Hatzikiriakos, C. W Stewart, D. L. Henry, and M. D. Buckmaster, "A New Processing Additive Eliminating Surface and Gross Melt Fracture in the Extrusion of Polyolephins and Fiuoropolymers," SPE ANTEC Tech. Pupers. 44,952 (1998). E. E. Rosenbaum, S. G. Hatzikiriakos, and C. W Stewart, '"The Melt Fracture Behavior of Teflon Resins in Capillary Extrusion," SPE ANTEC Tech. Papen, 41. 11 11 (1995). Y. Oyanagi, "Melt Fracture and Sharkskin," Progress in Rribber and Platics Technology, 13,23 1 (1997). Crank, J. and Park, G.S., Difision in Polvmers, Academic Press Inc., New York, N.Y- (1968). D.W. Van Krevelen, Properfies of Poivmers, Elsevier Scientific Publishers Company (1990). G. Marom, In: Polvmer Permenbilib, Elsvier Applied Science Publishers, London (1985). CJ. Park, and N.P. Suh, "Rapid PoIymer/Gas Solution Formation for Continuous Processing of Microcellular Plastics," ASME Trans., J. Manrtf: Sci. Eng., 118, 639- 645 (1996). M. Blander, and J.L. Katz, "Bubble NucIeation in Liquids," AIChe. J., 21, No. 5,833- 848 (1975). J.H. Han, and D.H. Han, "Bubble Nucleation in Polymeric Liquids. II. Theoretical Consideration," J. Polym Sci: Part B: Polymer Physics, 28,743-76 1 (1990). J.S. Colton, and N.P. Suh, "The Nucleation of Microcellular Thermoplastic Foam with Additives: Part 1: Theoretical Considention," Polym. Eng, Sci., 27, 485-492 (1987). J.S. Colton, and N.P. Suh, "The Nucleation of Microcellular Thermoplastic Foxn with Additives: Part II: Experimental Results and Discussion," PoLym. Eng. Sci., 27. 493-499 (1987). C. W. Stewart, " Nucleation and Growth of Bubbles in Elastomers," J. Polwn. Sci: Part A-2,8,937-955 ( 1970). S. K. Goel, E. Beckman, "Generation of Microcellular Polymeric Forims Using Supercriticai Carbon Dioxide. I: Effect of Pressure and Temperature on Nucleation," Polym. Eng. Sci., 34, 1 137- 1 147 ( 1994). S. K. Goel, E. Beckman, "Generation of Microcellular Polymeric Foams Using Supercritical Carbon Dioxide. IL Ceil Growth and Skin Formation," Polym. Eng. Sci., 34, 1148-1 156 (1994). V. Kumar, J. E. Weller, and R. MontecilIo, "MicrocelIular PVC," SPE ANTEC Tech, Papers, 38, 1452- 1456 ( 1992). N. S. Rarnesh, D. H. Rasmussen, and G. A. Campbell, "The nucleation of Microcellular Foams in Polystyrene Containing Low Glas Transition Particies," SPE AMEC Tech. Pupers, 39, 1828- L 831 ( 1993). D. F. Baldwin, C. B. Park, and N. P. Suh, "A Microcellular Processing Study of PoIy(Ethy1ene Terephtfialate) in the Amorphous and Semicrystalline States. Paa 1: Microcell Nucleation," Polp. Eng. Sci., 36, 1437-1445 (1996). J. S. Colton, "Making MicrocelluIar Foams frorn Crystdline Pdymers," Plm. Eng., d4,53-55 (1988). S. T. Lee, and N. S. Rarnesh, "Gas Loss Dunng Foam Sheet Formation," SPE ANTEC Tech. Papers, 41, 22 i7-2ZM (1995). A. H. Behravesh, C. B. Park, L. K. Cheung, and R. D. Venter, "Extrusion of Polypropylcne Foarns with Hydrocerol and Isopentane," J. Vinyl and Add. Tech,, 2, 349-357 ( 19%). S. T. Lee. "Shear Effects on ThermopIastic Foam Nucieation," Poiym. Eng, Sci., 33, 4 18-423 (1993). S. T. Lee, "More Experiments on Thennoplastic Foam NucIeation," SPE ANTEC TeeCi. Papers, 40, 1992- 1997 (1994). N. S. Rarnesh, H. Rasmussen, and G. A. Campbell, "Numerical and Experimentd Studies of Bubble Growth During the MicroceIIular Foam Process." Polym. Eng. Sci., 31, 1657-1664 (199 1). N. S. Rmesh, N. Dontula, H. Rasmussen, and G. A. Campbell, "Theoretical and Experimental Study of the Dynamics of Foam Growth in Thennoplastic Materiais," SPE ANTEC Tech. Papcrs, 37, 1292-1295 (199 1). R. D. Patel, "Bubble Growth in a viscous Newronian Liquid," Chem. Eng. Sci., 35, 2352-2356 (1980). C. D. Han, and H. J, Yoo, "Studies on Structural Foam Processing. IV. Bubble Growth During Mold Filling," Pofym, Eng. Sci., 21,s L8-533 (198 1). M. Amon, and C. D. Denson, "A Study of the Dynarnics of Fom Growth: Analysis of the Growth of Closely Spaced Sphericai Bubbles," Poiym. Eng. Sci., 24, 1026- 1034 (1984). M. Amon, and C. D. Denson, " A Study of the Dynamics of Foam Growth Simplified: Anaiysis and Experimentd Results For Bulk Density in Structurai Foam Molding," Potym Eng. Sci., 26,255-267 (1986). A. Arefmanesh, S. G. Advani, and E. E. MichaeIides, "A Numericd Study of BubbIe Growth During Low Pressure Structural Foam Molding Process," Polym. Eng. Sci., 30,1330-1337 (1990). C. D. Han, and C. A. Villamizar, "Studies on Structural Foam Processing 1. The Rheology of Foam Extrusion," Poiym. Eng. Sci., 18,687-698 (1978). C. A. Villamizar, and C. D. Han, "Studies on Structural Foam Processing II. Bubble Dynamics in Foam injection Molding," Polym. Eng. Sci., 18,699-7 10 (1978). A. H. Behravesh, C. B. Park, and R. D. Venter, "Extrusion of Low-Density, Microcellular HIPS Foams Using COZ" in: Cellular and Microcell~ilarMaterials, V. Kumar and K.A. Sealer, eds., ASME, N. Y., 47-64 (1996). L. Matuana-Malanda, C. B. Park, and J. J. Balatinecz, "Effect of Ce11 Morphology on the Properties of Microcellular Foarned PVCiWood-Fiber Composites," In: Cellular und Microcellrtlar Materiais, Kumar, V. and Sealer, KA., eds, ASME, N.Y., 1-16 ( 1 996). S. T, Lee, and N. S. Ramesh, "Study of Foam Sheet Formation Part üI: Effects of Foam Thickness and Ce11 Density," In: Cell~darand Microcellular Materials, V. Kumu and K.A. Sealer, eds., ASME, N. Y., 7 1-80 (1996). S. T. Lee, N. S. Ramesh, and G. A. Campbell, "Temperature Effects on Thennoplastic Foam Sheet Formation," SPE ANTEC Tech. Papers, 39, 3033-3036 (1993). M. A. ShaEi, J. G. Lee, and R. W. Flumerfeit, "Prediction of Cellular Structure in Free Expansion Polymer Foam Processing," Polym Eng Sci., 36, 1950, (1996). M. Shimoda, M. Tanigaki, and M. Oshima, Foams 2000,9 1 (2000). N. S. Ramesh in: Fonm Ertrrrsion Principles and Practice, ed. S. T. Lee (2000). W. A. Johnson, and R. F. Mehl, Trans. AIME, 135,416 ( 1939)- J. D. Hoffman, and J. J. Weeks, "Rate of Spherulitic Cryscallization with Chain Folds in PolychIorotrifluoroethylene," J. Chem. Phys., 37, 1723 (1962). A. Peterlin, "Plastic Deformation of Crystalline Polymers," Polym. Eng. Sci, 17. 183 (1977). J. W. Gibbs, Trans. Conn. Acad. 111,343 (1878). M. Volmer, and M. Marder, 2. Phys. Chem.. Abt. A, 15,97 (193 1). F. P. Price, Nucleation, New York ( L969). D. Tumbull, and I. C. Fisher, J. Chem. Phys., 17.71 (1949). M. L. Williams, R, F. Landel, J. D. Ferry, 3. Amer. Chem. Soc., 77,3701 (1955). J.D. Hoffman,and J.I. Lauritzen Jr., J. Res. Nat. Bru-. Std.-A., 65,297 (1961)- J. D. Hoffman, SPE Transactions, 4.3 15 (1964). J.I. Lauritzen Jr., and J.D. Hoffman, "Extension of Theory of Growth of Chain- Folded Polymer Crystals to Large Undercoolings," J. Appl. Phys., 44,4340 (1973). J.D. Hoffman, "Role of Reptation in the Rate of CrystalIization of PolyethyIene Fractions from the Melt," Polymer, 23,656 (1982). J. D. Hoffman, "Regime iII Crystallization in Melt-Crystailized Polymers: The Available Cluster Model of Chain Folding," Polymer, 24,3 (1983). F. C. Frank, and M. Tosi, Proc. R. Soc.: Ser. A., 263,323 (1961). 1. C. Sanchez, and E. A. DiMarzio, J. Chem. Phys., 55,893 (197 1). D. M. Sadler, and G. H. Gilmer, "Model for Chain Folding in Polyrner Crystals: Rough Growth Faces are Consistent with the Observed Growth Rates," Polymer, 25, L446 (1984). D. M. Sadler, "On the Growth of Two Dimensional Crystals: 2. Assessment of Kinetic Theories of Crystallization of Polymers," J. Chem. Phys. 87, 177 I (1987). J. 1. Lauritzen Jr., J. Chem. Phys,, 45,4444 ( 1966). J. 1. Lauritzen Jr., and E. Passaglia, J, Res. Nat, Bur. Std.-A.,7 1,26 1 (1967). E. A. DiMarzio, J. Chem. Phys., 47,345 (1967). A. N. Kolmogoroff, Isvesr. Akud. Nartk SSSR Ser. Math, 1,335 ( 1937). W,A. Johnson, and R. T. Mehl, Trans. AIME, 135,416 (1939). U. R. Evans, Trczns. Faraday Soc., 4 1,365 ( 1945)- M. Avrami, "Kinetics of Phase Change 1: General Theory," J. Chm. Piys., 7, 1\03 (1939). M. Avrami, "Kinetics of Phase Change U: Transformation-Time Reiations of Random Distribution oCNuclei," J. Chem. Phys., 8,212 (1940). M. Avrami, "Granuiation, Phase Change, and Microstructure Kinetics of Phase Change III," J. Chem. Phys., 9, 177 (1941)- B. Wunderlich, Macrornotecular Physics, 2 (1976). N. P. Suh, The Princiules ofDeskn, Oxford, New York (1990). A- Yeung, Master thesis, University of Toronto (1998). C. B. Park, Ph.D. Thesis, Massachusetts Tnstitute of Technology, Cambridge. MA (1993). A. Behravesh, Ph.D. Thesis, University of Toronto (1998). Product Catalogue, The Flow and Level Handbook, Vol. 28, Omega Engineering, Inc. (1992). M. J. Moran, and H. N. Shapiro, Frindamentnls of En~ineerThennodvnamics, 2nd ed., John Wiley & Sons ( 1992). H, E. Naguib, E, V. Richards, R. Pop-Iliev, J. X. Xu, and C. B. Park, "Measurement of the Swelling of PP Melts with Dissolved Butane," SPE, ANTEC Technicd Pcrpers, Vol. 47, pp. 1537-1545, (2001) W. M. Haynes, and R. D. Goodwin, Thermophvsical Properties of Nomal Britme, National Bureau of Standards, Colorado. H. E. Naguib, S.-W. Song, C. B. Park, and Y. J. Byon, "Effect of Supercritical CO7 and N2 on Crystallization of Lineac and Branched Propylene Resins FiIled with Foarning Agents," SPE, ANTEC Technical Popers, 46, pp. 1867-1873, (2000). H. Janeschitz-Kriegl, H. WippeI, C. PauIik, and G. Edcr, "Polymer CrystaIlization Dynamics, as Reflected by the Differential Scanning Calorimetry. Part 1 On the Calibration of the Apparatus," Colloid & Polym. Sci., 271, 1107 (1993)- A, J. De Nicolas, A. F. Galombos, and M. D. Wolcowicz, Poly. Mat. Sci. and Eng., 67, 106, (1992). X. Wang, C, Tzoganakis, and G. L. Rempel, "Chemical Modification of Polypropylene with PeroxidelPentaerythritol Triacrylate by Reactive Extrusion," J. Appl. Polym. Sci., 61, 1395 (1996). J. Varga, I. Menczel, and A. SoIti, A. I. Themal Anulysis. 20,23 (198 1). A. GhijseIs, N. Groesbeek, and C. W. Yip. Pofymer, 23, 19 13 (1982). D, Dieckmann and B. Hottz, "Aging Modifiers for Extruded LDPE Foarn," Foams 99, NJ, 2 1 1 (1999). Y. Sato. M. Yarugi, K. Fujiwara, S. Takishima, and H. Masuoka, "Solubilities of Carbon Dioxide and Niuogen in PoIysîyrene under High Temperature and Pressure." Fllrid Phase Eqttil., 125, 129 (1 996). S. S. Park, MA.Sc. Thesis, University of Toronto (1999). H. E. Naguib, and C. B. Park, SPE, "A Study of Melt Fracture of Linear and Branched Propylene/Butane Solutions with Foaming Additives," SPE, AMEC Technical Papers, Vol. 16, pp. 1025-103 1, (2000). J. M. Dealy and K. F. Wissbrun, Melr Rheolocv and its Roie in Plastics pro ces sin^ Chapman and Hal1 (1995). N. Maltwitz and N. S. Ramesh, "Non-isothermal Model to Study the Influence of Blowing Agent Concentration on Polymer Viscosity and Gas Diffisivity in Thermoplastic Foam Extrusion," Foams 98, N. J., 243, (1998). H. E. Naguib, C. B. Park, A. Hesse, U. Panzer, and N. Reichelt, "Strategies for Achieving Fine-celIed Low-Density PP Foams," Blowing Agents and Foaming Processes, Frankfurt, Paper 7, (200 1). C. B. Park, G. Liu, and H. E. Naguib, "Chailenge to Forty-fold Expansion of Biodegradable Polyester Foams Using CO2 ," Cellrtlw Polym., 18, 367 (1999). W. R. Veith, Difiision in and throrrgh polvmers, Hanser Publishers (1990). G. Sims and W. Sirithongtaworn. Cell. Polym., 16,271 (1997). United Nations Environment Prognm, Handbook for the Montreal Protocoi for Substances thDedete the Ozone Laver,Montreal (199 1). W. C. Reynolds, Thermodvnamic Properties in SI, Stanford Department of Mechanical Engineering ( 1979). R. A. Gorski, R. B. Ramsey, and K. T. Dishart, J, Cell. Plast., 22,21 (1986). H. E. Naguib, J.X. Xu, C. B. Park, A. Hesse, U. Panzer, and N. ReicheIt, "Effects of Blending of Branched and Linear PoIypropylene Resins on the Foamability," SPE, ANTEC, Technical Papers, Vol. 47, pp. 1623- l630,XIO 1. R. Verseput, Desi.cned Erperiments and Information. S Matnx Corporation, ( 1995) J. ~lacho~oulos,Introducrion ro Plastics Processine McMaster University. Dept, Chem. Eng. (1997) C. W. Macosko, Rheoiow: Princides, Measuremenrs. and Ap~licnrions, VCH Publishers, New York (1994) C. D. Han, Rheolonv in Polvmer Processing Academic Press (1976) 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 far in the z-direction suggests that the flow is unidirectional, and therefore, the r- 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
+----I alv a 2 a..] +n[a [ 1 a ( al- r ar n>r) cr~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 de[197]:
Hence Eq. A1.7 cmbe simplified to yield
Eq. Al.ll (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 Clare the integration constants. The boundary conditions are (see Figure Al.1): Afr=R*vZ =O, (A1.17 a)
At r=O=&=O.av, (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 ~hcFully Drveiopcd Region ofa Circuhr Die Appendix 2
Statistical A nalysis
Table A.l Independent Variable Settings (Single Constraints)
Variabk Variable Name RnngeKevel (s) Units I I Temperature (T) OC 1205T1180 Gris percentage (Gas %) Gas % 5 (UD ratio) 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 Computed Error Statistic Error Statistic Name Value Error % 1.go41 37121 Experimental Error (It) 0.1 88680493 95% Confidence Limits (I) 0.5238621 17 Adj. R Square Lower Lirnit 0.892833963 Regression Statistics Table AS General Regression Statistics: Transfoimed Volume Expansion Regression - Regression Regression Regression Statktic Name Statistic Value Target Name Target Value R Square 0.974834597 Adj. R Square Upper Lirnit 0.99197221 1 Adj. R Square 0.9635101 66 Adj. R Square Mean 0.980958629 Adj. R Square Lower Lirnit 0.892833963 Standard Error (t) 0.261 194306 Standard Error (3 0.188680493 Observations 30 Degrees Of Freedom Lack-Of-Fit 16 Lack-Of-Fit F-Ratio 2.1 45431 108 P-value 0.2402601 64 Table A6 Regression ANOVA Statistics: Transformed Volume Expansion Source Of Degrees Sum Of Mean Vanàtion Freedom Squares Square F- Ratio P-value Regression - 9 52.85483451 5.87275939 86.08243246 5.1 396E-14 Residual 20 1.364450149 0.068222507 Total 29 54.21 928466 Coefficients Table And Models Table A.7 Model Coefficients: Transformed Volume Expansion - -- Coded Coefficient Lower 95% Upper 95% Vanable Coefficient Standard Coefficient Confidence Confidence Name Value Limit iimit lntercept 3.945002965 3.7421 37807 4.1 47068123 XI -1 .O35280948 -1 -17205327 -0.898508626 X2 0.759955449 0.61 8222707 0.90168819 X5 -0.1 3971333 -0.248761 021 -0.030665639 l2 -1.3657491 n -1.606254766 -1.1 25243588 (m2 -0.898630631 -1.1 52489506 -0.644771757 XI*X2 -0.6516091 9 -0.824566242 -0.4786521 38 X2'X4 0.269371 443 0.1 1870523 0.420037655 x3*x4 -0.31 6477941 -0.46173851 7 -0.1 7121 7366 X3*X5 0.247939966 0.1 00995968 0.394883965 Coded Variable Name Key XI =T Model Term Ranking Table A8 Model Term Ranking: Transformed VoIume Expansion Model Term Model Term Coefficent Mode/ Term Mode1 Term Name Range Value Effect Rank X 1 2 -1 .O35280948 -2.070561 897 1.O0 XS 2 0.759955449 1.51 9910898 0.73 (X1l2 1 -1.3657491 n -1.3657491 n 0.66 X1 'X2 2 -0.651 60919 -1 30321838 0.63 (x2)2 0.888888889 -0.898630631 -0.798782783 0.39 X3'X4 2 -0.31 6477941 -0.632955883 0.31 XTX4 2 0.269371 443 0.538742805 0.26 X3'XS 2 0.247939966 0.495879933 0.24 X5 2 -0.1 3971333 -0.27942666 0.1 3 Transformed Coded Variabk Model 1ransforrned Natural Variable Model Volume Expansion = 3.945003 Volume Expansion = -32.861 18 -1 .O35281 ' (Xt ) +0.45694 ' (T) +0.75996 * (X2) +0.54005 '(Gas %) -0.13971 3 (X5) +2.82463 ' (Talc -1.365749 ' (XI2) -0.005883 ' (M) -0.898631 * (XZ2) -0.1 069% ' (UD) -0.651609 ' (XI 'X2) -0.001517 - (02) +0.26937 ' (XZX4) -0 .O1 5976 ((Gas %)2) -0.31 6478 ' (X3'X4) -0.002896 ' (T ' Gas %) +0.24794 ' (X335) +0.03592 ' (Gas % ' M) -0.31 6478 * (Taic % ' M) +O.OS449 ' (TaIc % ' UD) Model Term Ranking And Charts Table A.9 Model Term Ranking: Transfoned Volume Expansion Mode! Ten ( Mode! Ten ( Coefficient ( Mode/ Ten Range Value Effect Rank 2 TiziGG-1-2.070561 897 1.O0 2 1 2 0.888888889 2 2 2 Model Term Ranking Pareto Chart I Model Term Ranking Pie Chart ):pI Residuals Data And Plots Table A.10 Residuals Data: Transformed Volume Expansion Transfomed Deleted Obsewed Predkted Studentized Sample No. Volume Expansion Volume Expansion Residuals Residuals 1 0.457487276 0.1 97438691 0.993830921 2 4.31 9647868 -0.461 13113 -2.1 2200639 3 1.247043736 -0.1 66934919 0.869426698 4 3.084571 803 -0.1 483241 77 0.607634754 5 4.01 1518226 -0.02355321 7 ,0.104208128 6 3.658506099 0.3041 14959 1.51 597702 7 4.01 1518226 0.01 5528456 0.068692524 8 0.869579671 -0.230360833 -1.262804666 9 2.764039632 0.120481712 0.60472089 10 2.980492905 -0.01 0334883 ~0.048930984 11 3.029741 423 0.442069738 2.049369029 12 3.86744792 0.04447508 0.21 9441 534 13 0.779537644 0.072747758 0.352831 699 14 3.348625228 0.1 7020703 0.724460757 15 3.077341 576 0.298367043 1.482632655 16 3-1 071 26075 0.344289023 1.5661 61 587 17 0.82671 671 9 -0.1 8749788 -1.OZ5SOl9l9 18 3.9721 35609 -0.052243645 .0.253394794 19 3.1 94424744 -0.397448334 -1 -726274342 20 3.745207645 0.1 02703415 0.507268892 21 3.1 39594364 -0.31 2576644 -1.342293477 22 3.348625228 0.1 552791 52 0.659396785 23 0.971755745 -0.07771 O622 -0.355485298 24 4-102520974 -0.1 16419425 -0.497878826 25 0.971 755745 -0.07773 0622 -0.355485298 26 0.55549238 0.1 09769597 0.523408585 27 0.598355332 0.040863506 0.21 0352292 28 3.658506099 -0.22791 1741 -1 -1O7 140256 29 4.00452387 -O.lO67Ol214 -0.482238332 30 4.325664831 0.1 78524125 0.848866591