Extrusion Processing for Manufacture of Low-Density, Fine-Celled Polypropylene Foams ~ani E. Naguib A thesis submittd in conformity with the requirements for the Degree of Doctor of Philosophy Department of Mechanical and lndustrial Engineering university OF Toronto @Copyright by B. E. Naguib 2001
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Extrusion Processing for Manufacture of Low-Density, Fine-Celled Polypropylene Foams
~ a n i E. Naguib
A thesis submittd in conformity with the requirements for the Degree of Doctor of Philosophy
Department of Mechanical and lndustrial Engineering university OF Toronto
@Copyright by B. E. Naguib 2001
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Extrusion Processing for Manufacture of Low-Density, Fine-Celled Polypropylene Foam
Hani E. Naguib Degree of DacCor of Philosophy, 2001
Department of Mechanical and Industrial Engineering University of Toronto
A continuous extrusion process for the manufacture of low-density, fine-celled
poiypropylene foams is presented. Due to its outstanding Functiond characteristics and low
materiai cost, polypropylene foarns have been considered as a substitute for other
thennoplastic foams in industriai appIications. However, only limited research has been
conducted on the production of polypropylene foarns because of the weak melt strength, and
no research has been conducted to investigate the mechanisms that govern the expandabiiity
of poIypropylene foams. This thesis presents the effective scrategies for increasing the
volume expansion ratio as welI as the mechanisms goveming the foam density of
polypropylene foams. The basic strategies taken in this study for the promotion of a large
volume expansion ratio of polypropylene foams are: (a) to use a branched materiai For
preventing ceIl coalescence; (b) to use a long-chah blowing agent with low diffusivity; (c) to
lower the melt temperature for decreasing gas Ioss during expansion; and (d) to optirnize the
processing conditions in the die for avoiding premature crysrdlization. The effects of
processing and materiais parameters on the foam morphoiogies of poIypropyIene materiais
were thoroughly studied using a single-screw tandem foam extrusion system. A careFuI
analysis of extended experimentd resuIts obtained at various processing conditions indicates
that the final votume expansion ratio of the extruded polypropylene foams blown with butane
is governed either by Ioss of blowing agent or by crystdlization of the polyrner matrix. By
tailoring the processing conditions in the die, ultra low-density, fine-celied polypropylene
foarns with very high expansion ratio up to 90-fold were successfuIly produced from the
branched poiypropyiene resins. Fundamental snidies have also been conducted to investisate
the effect of various processing and materials parameters on the thermodynamic, thermal and
meit fracture behaviors of polypropyIene melts with foaming additives that influence the ceIl
morphology of poiypropylene foams.
ln memory of my Aunt Aida, who passed away during my Ph.0. program.
1 know you still hear me and pray for me. 1 thank you for al1 the pars you took care of me, for your love, support and encouragement. 1 could not make it without you. You will aiways be in my
heart and my thoughts
In memory of my Mother who was and still is my inspiration in my life
70 my great Father to whom 1 owe al1 that 1 am and al1 that 1 am going to ber
and al1 that 1 can ever hope to be
Acknowledgments
1 would like to start by giving thanks to the beneficent and mercifuI God, my Lord
and Savior Jesus Christ, for he has covered me, supported me, preserved me, accepted me
ont0 him, had compassion on me, sustained me, and brought me to this hour.
1 would like to express my sincere gratitude to my supervisor Professor Chu1 B. Park
for providing guidance and encouragement throughout my iesearch. I will never forget the
support he gave at my most difficult times.
1 would Iike also to thank my Ph. D. thesis committee: Prof. Steve Bdke from
Department of Chemical Engineering, Prof Shaker Meguid, Prof Beno Benhabib €rom the
department of Mechanical and Industrial Engineering.
Special thanks to Prof. O. A, Aziz, Prof. M. Al Gammal, Prof. A. AbdeI Messih and
Mr. Albert Mickail for their guidance throughout my career and the time and effort spent.
My gratitude is extended to the Department of Mechanical and Industnal Engineering
at the University of Toronto for providing the University of Toronto Open Fellowships, the
Ontario Graduate Scholarship, as well as, the NSERC Postdoctoral fellowship. Also, i would
like to thank Boredis AG. Company, in Austria for their funding and support in this project.
I would also like to thank my colleagues in the Microcellular Plastics Manuhcturing
Laboratory for their help and Friendship over the p s t four years. They include Simon Park,
Dmitry Ladin, Dr. Valentina Padareva, Anthony Yeung, Amir Behravesh, Ghaus Rizvi,
Esther Richards, Patrick Lee, Linda Lin, Remon Pop-iiiev, Deepak Fernandes, Dr. Chris
Song, Dr. Sang Mae Lee, Dr. Yuejian Liu, Minhee Lee, Xiang Xu, Donglei KU, Wanlin Chen,
Chris Ozolins, Anjan, Xioyang Guan, Rehan Khan, Gumjan Guo, and Haiou Zhang.
Especially, sincere gratitude poes to Simon and Dmitry who shared common thoughts. 1
would like to thank al1 the undergraduate students that worked with me Erin Youn, Young fi,
David Allas, Suzan Oh, Won Park, Carri Li, Joyce Lam, Fames Koo, Brandon Lee, Michael
Lee, Ibrahim Abu Eisha, and others. Also, 1 wish to acknowledge the professiond technical
support from Mike Smith, Len Rooseman, Jeff Sansome and Dave Eisdaile in the Machine
TooI Laboratory, from Wendy and Amanda in the generai office, from Mary Rose, Teresa,
Dan and Sheila in the purchasing department, and from Oscar in the computer support. A
special thanks to Esther Richards for proof reading my thesis and for Dimtry Ladin,
Mohammed Attia, DongIei Xu, Linda Lin, Ghaus Rizvi, Xiang Xu, WanIin chen and Lianne
Ing for the time they spent for helping me at the final stage of the thesis.
iv
1 would like to thank my family al1 over the world; from Egypt: 1 would like to thank
my wife for her continuous efforts to provide me with her love and support, my great father,
my beloved brother Rami, his wife Mary and Daniel, to my beloved cousins: Sarni, Nagi and
their families, Shadi, and to a11 my aunts and uncles especidly Aunt Mary, Uncle Soria] and
Uncle Atef for their continuous support and encouragement. My family in Canada: my God
father Uncle Esmat, to my beloved cousins: Tarek, Kim, Mark, Magdy, Samer, Rosemary,
Chris and Sally. A special thanks to my parents in Canada Uncle Said and Aunt Mona for
their continuous care and love. My family in the States: my God mother Nadia for her
continuous love and support, my uncles and riunts: Osiris, Nela, Menes, Madelaine, Kamilia,
and my cousins: Mona. Marc, Sergio, Franco, and their families, Gehan and James.
1 would like to thank father Reuiss, father Georgeos, father Makari from St. Georges
Church in Toronto, father Amonios and father Messaeil form St. Mark church in Toronto. A
special thanks to the youth group in my church, who supported me with their prayers and
provide me with more love to rny Lord Jesus Christ.
1 also would like to thank many friends ihab and Lydia, Essam and Marianne, Adel and
Sally, Bassem and Dalia, Sameh and Gehan, Roger and Mira, Jules and Mary, Joseph
Armanios, Talaat and Mona, Nagui and Dalia, Dimitry Saad, and Alfred Mobayad.
Finally, 1 would like to praise my Lord with David the prophet and king by saying
I I O LORD. yorr have searched me and you knoiv me. YOU knoiv ivhen C sir and ivhen I rise; You perceive my rfroirgl~ts from afar. Yori discent my going otrr and my king down:
Yoir are faniriliur ivith al1 niy ivays. Before a word 13 on Iny tongrle You know ir cumplrrely. O LORD. Yori hem me in--bellinriund before: Yori have laid yoiir hand iipon me.
Such knoivledge is too wondefil fur me. tao lofty for me to airain. Where can I go fiom p u r Spirit? Wliere can Iflreji-om yuurpresence?
I f I p up to the heavens. You are there; iff make tny bed in the depths You are there. IfI rise on the wings of the daivn, iflsettle on rhe fur side of the sea, even tlrere your hand ivill giride me. your nghr hand will hold me fast.
IfI say, "Siirely the darkness will hide me and the light become nighr urotitrd me," even the darkness will nor br dark ro Yoti; the nigfrt will shine like the duy,
for darkness is as tighr to You. For You creared my inmost being: You knit me togerher in my mother's womb.
I praire You because I am fearfitlly and ivonderfulfy made; Your ivorks are rvonderfuf. I knorv rharfiill ivefi.
My frame ivas not hiddenfrorn Yori when 1 was made in the secret place. When I ivas ivoven together in the depths of the earrh. Your qves saw my rinformed body. Al1 the duys ordainecifor me iuere wiften in Your book before one of rhem came to be.
Hoiv precioiw to me are Your rhorighrs, O God! Hoiv vast fs the strm of them! Were 1 to cortnr them. they would ournumber the grains of sand
When I nwuke, I am $tilt with you "- Psalm 139 (1-19)
The molar heat of sorption, AH,, can be a negarive or a positive d u e depending on the
pol ymer-gas system.
Equations (2.1) and (2.2) allow us to determine the solubility of a blowing agent in
the polymer at the pracessin; pressure and temperature, and estimation of the soiubility of
CO2 in some polymers has been given in the literature [17,18]. However the solubility of
butane in polyrners was not available in the literature. In the acrual ~xtrusion foarn
processing, the ratio of gas [O polymer weight is maintained below the solubility Iimit, by
controlling the flow rate of the polymer and gas with injection amount.
The initia1 dope of the curve in Figure 2.1 corresponds to the diffusivity of the
blowing agent into the polymer matrix, and can be used to calculate the diffusivity using the
following equation 11051:
where is the value of (tfl?) at MJM- =1/2. By rearranging Equation (7.4). the time
required for the compietion of absorption cm be approximated from Equation (2.5) as given
below [107]:
Equation (2.5) indicates that the time of absorption is inversely proponional to the diffusivity
(D) and proportional to the square of the diffusion distance (ha).
The diffusivity D is mainly a function of temperature, and its influence cm be
explained by the following equation [105,106]:
where Do= diffusivity coefficient constant (cm%),
EF activation energy for diffusion (J).
Dissolution
During the batch process, the gas absorption process into the polyrner rnatrix will
automaticatly terminate when the concentration of the bIowing agent dissolved into the
polyrner matrix and on its surface has reached an equilibrium value. Thus, large voids are not
generated in the batch proccss since it is impossible for any extra gas to be absorbed by the
polymer above its solubility limit. However, in a continuous process, there is the possitiility
of retaining an undissolved amount of jas in the polyrner matrix, if excess gas is injected,
and therefore it is critical to ensure that the amount of gas injected must be below the
solubility limit at the processing conditions. The main advantage of utilizing the extrusion
foarning process, apart from it being the most cost-effective method among nurnerous
foaming processes, is the reduction of dissolution time due to higher gas diffusivity D at the
high processing temperature.
Although the proper amount of gas injection is emphasized, it does not necessririIy
guarantee formation of a uniform solution because the time required for completion of gas
diffusion into the polymer is equally important. If the required gas difision time is ronger
than the residence time of the melt, Le., the time between gas injection and nucleation, it is
obvious that a uniform solution would not be achieved.
Park et al. [18,108] studied the diffusion phenomenon in an extrusion process
containing a mixing screw. It was observed chat shear mixing promotes convective diffusion,
In convective diffusion, a high gas concentration region (gas bubble) is brought into contact
with a 1ow gas concentration region (polyrner melt). Moreover, due to stretching of gas
bubbles, in the shear field generated by the screw motion, the diffusion process is enhanced
as the interfacial area is increased. The tirne for completion of dissolution was estimated.
They also proposed use of a dissolution enhancing device containing static mixers in the
extrusion system and claimed that it would promote the dissolution process by generating
shear fields as the mixing elements repeatedly reorient the melt dong the flow direction, thus
promote solution formation.
2.3.2 Nucleation
Nucleation is a critical step in the fine-celled foaming process because a large number
of cells must be generated in order to achieve smdl-size celIs. Nucleation of bubbles in the
polyrner cm be modeled by the cIassica1 nucleation theory [109,110], which was originally
developed for a single-component system where the second phase is created by evaporation
of the iiquid when superheated. Blander and Katz [IO91 extended the classical theory to a
diffusion system in which one component is volatile and foms bubbles.
A diffusion process of the dissolved component into the nucleacion site _ooverns
bubble nucleation, and fine cells can be nucleated either randomly throughout the polymer
matrix that is known as homogenous nuckation, or prefenbly at certain sites that is known as
heterogeneous nucleation. Heterogeneous nucleation, which can be promoted by using
additives in the liquid, is more prefernble, since less energy is required for nucleation
compared with homogenous nucleation and bubbles are more likely to nucleate at the
preferred sites provided by the additive particles.
Homogenous Nucleation
Colton and Suh [ i l 1,1121 modeled the nucleation behavior in microcellular foaming
using the classical nucleation theory. According to this theory, the work required to genente
a bubble of radius r in a body of a liquid can be given by (1 111:
where the first term (y,, A,) is the work required to create a bubble with a surface area Ab
and a surface tension ~pb, and the second term (MV,) is the work done by che expansion of
gas inside a bubble of volume Vb. The difference between the gas pressures inside the bubble
and the surrounding matrix. AP, is approximated to be the saturation pressure [ I l l]. By
replacing Ab and Vb with the area and volume of a sphere nuclei, Equation (2.7) becornes:
Figure 2.2 depicts the variation of energy W with radius r. it is secn that there exists a
maximum energy, which must be overcome to forrn a nucleus that continues to grow
spontaneously. if the energy induced is less than this maximum energy (or free energy
barrier), the bubbles (with radius rce) will collapse. This activation energy for
homogeneous cell nucleation AG,:,,, is calculated by differentiating W with respect to r
which yieIds [ I 10,1111:
The nudeation rate is then represented by [110,112]
where C,, = concentration of gas molecules in solution (#lm3),
f;, = frequency factor of gas molecules joining the nucleus (Ils),
k = BoItzman constant (JK).
The classical nucleation theory for a homogenous system predicts that the higher the
saturation pressure AP, the greater the number of ceIIs nucleated because of the lower
activation energy for ceIl nucleation Il?]. This effect has been experimentally verified in the
batch process [19,111- 1 151. The saturation pressure in a batch process corresponds to the gas
concentration in the polymer as given by Henry's law (Equation (2.3)). Thus, the effect of
pressure can be intetpreted as its influence on the amount of gas absorbed into the polyrner.
When the amount of gas increases, the chance of nucleation is higher and ri larger number of
nucleated cells is achieved.
Park et al. studied the effect of the pressure drop rate in an extrusion process and
indicated that the pressure drop rate is influencing the number of cells nucleated [17]. They
induced various pressure drop rates by using viirious nucleation dies and examined the final
foam structures. Their study shows that the higher the pressure drop rate, the greater the
number of cells nucleated, which may be explained by the mechanism of cell
growthhucleation cornpetition [17]. When the pressure drop rate is high, the poiymer-gas
system experiences a certain pressure drop in a shorter tirne period. During this shorter time
period, the already nucleated cells grow Iess because less time is available for gas diffusion
to the cells. Therefore, less gas is consumed for ce11 growth, and more gas is rivailable for
further nucleation in the polymer matrix (Equation 2.t0, Co is higher). As a result, the final
foam structure contains a Iarger number of cells, and the cells are smaller since Iess gas is
consumed for each ceIl. Hence, it is essentid that a higher pressure drop rate be induced to
achieve a fine-celled structure in the foaming process.
Stewart's work on elastomers [Il31 shows that the number of nucleated cells
increases with increasing the temperature. Similarly, the results obtained by Kumar et ai.
[116] on PVC also show an increasing effect of temperature on ceIl nucleation, and Ramesh
et al. [117] demonstntd that in a PS-CO: system, increasing the temperature increases the
cell density. Goel and Beckman [ L 141 studied the nucleation behavior of a PMMA-COI
system and demonstrated that increasing the foaming temperature tends to decrease the
number of nucleated cells. Baldwin et ai. (1 181 show that iincrerising the foaming temperanice
increases the ceIl density in amorphous PET and CPET below 100°C. and produces no
further change above 100°C. On the other hand, they found that temperature does not
significantly affect the cell density in semicrystailine PET and CPET. They suggested that
nucleation occurs rnainly during the pressure release andor in the early stage of heating.
In surnmary, generating a large number of bubbles by dissolving a large amount of
gas in the polymer is a critical step to obtain a fine ce11 structure for polypropylene foarns-
However, ensuring the amount of gas dissotved in the polyrner below the solubility limit
according the processing conditions is equaily important. Nevertheless, a sufficiently high
pressure is required, as stated in Henry's Iaw, to maintain a large amount of gas dissolved in
the polymer.
Heterogeneous Nucleation
Bubble nucleation is heterogeneous when it initiated at some preferred sites by
mixing the polymer with an additive. in generaI, nucleation tends to occur at the boundary of
the matrix and additive rather than inside the polyrner matrix as with hornogenous nucleation.
At the boundaries of the matrix and additive, the fee energy barrier for nucleation is lower
than that in hornogenous nucleation (Figure 2.3); therefore, nucleation is more likely to occur
heterogeneously rather than hornogeneously. By controlling the amount of additive, one cm
generate the desired number of bubbles [lI9-1211. However, in general, it is difficuit to
generate a large number of cells of micron size using additives due to poor dispersion [Il91
and agglorneration [121] of the additive particIes. Using an appropriate nucleating acgent that
consists of very srnall particles (less than a micron) well dispersed in the polyrner matrix
without agglomeration, one could produce a large number of cells for fine-celled foam
application [119].
The shear force is aiso
system such as extmsion [122,
bubbles increases. Lee [122,123
affecting the heterogeneous nucleation rüte in a dynarnic
1231; as the shear force increases, the number of nucleated
1 developed a lump cavity nucleation model, which indicated
that the cavities on the rough surfaces of the tiny nucleating particles form potentiai sites for
bubble nucleation. M e n the gas phase in the cavity grows and rnatured by difision of the
dissolved blowing agent into the cavity or by a pressure drop, the applied shear force
enhances the chance of detaching it from the cavity to generate a bubble.
2.3.3 CeIl Gmwth
Ce11 growth is due to the continuous difision of gas from the polymer matrix into the
nucleated cells since the soiubility of grts in the polymer is decreasing with the pressure drop
(Henry's law). Concurrently, because the pressures inside the cell is greater than the pressure
in the surrounding matrix, the cell tends to grow to rninimize this difference [i l . The
viscosity, the diffusion coefficient, the gas concentration, and the number of nucleated
bubbles are goveming the growth dynamics. Foc instance, when the polymer viscosity
decreiises via a temperature increase, the rate of growth increases due co the decrease in
resistance against cell growth. Conversely, ceIl growth terminates when al1 the gas ciissolved
in the polymer rnatrix is depleted or the matrix is too stilf to allow further growth. Figure 2.3
rnodels a nucleated ceIl inside a polymer matrix charged with a biowing agent (such as gas).
When the ceil is nucleated, the gas concentration around the cell decreases. This generares a
gradient of $as concentration around the cell, which is responsible for funher cell growth.
The dynamics of cell growth have been extensiveiy studied [t24-132,137-1391,
Arefmanesh et al. El301 modeled the cell growth in a finite sea of liquid flowing in a
mold by considering the pressure drop along it. They described the effect of various
parameters on the dynamics of the growth and simuiated the effects of the diffusivity and
viscosity on cet1 growth. It was shown that the higher the diffusivity, the more rapid the cell
growth. In contrast, increasing the viscosity retarded the growth process. The results showed
that in the initial stage of growth, the viscosity hm a greater effect. The effect becomes
weaker as growth proceeds so that eventuaily the rate becomes the same at different
viscosities.
Han and Villarnizar [13 11 investigated the growth of bubbles in a foam sheet
extrusion process. They studied the pressure profile of foam along the die and found that
when the fiow rate increases, the point at which foaming becomes visible moves closer to the
die exit. On the other hand, a higher gas concentration and higher temperature tend to move
the foarning point away from the die exit. This observation is important since early ce11
growth causes deformation of ceIIs to a non-sphericd shape due to the presence of a shear
field dong the die, which mrty promote ce11 codescence- At a higher flow rate the residence
time is shorter, and hus the already nucleated cells grow less. As a resuIt, it is Iess tikeiy that
adjacent celts codesce, which would cause foam degradation.
Villamizar and Han [132] studied ceII growth in an injection molding process. They
studied the effect of the rnelt temperature, injection pressure and blowing agent concentration
on the final bubble size. It was observed that at Iower rnelt temperatures, bubble growth is
slowed down, and also fewer bubbles were detected by the naked eye. This was due to the
increase in viscosity which opposed the rate of ce11 growth and the lower diffusivity
(Equation (2.7)), which means that a less blowing agent was supplied ta the bubble. Their
results also showed that as the arnount of blowing agent increased, the final ce11 size became
greater. it is interesting to note that when the blowing agent concentration exceeded the
solubility limit, a small number of large bubbles appeared. Therefore, the arnount of injected
blowing agent must be maintained below the solubility limit. Their research also showed that
an increase in the injection pressure (thus, a reduction in the filling tirne) resulted in a
decrease in the ceIl size and a uniforrn cell distribution. This effect is similar to that of the
flow rate, which was explained earlier* As the filling time decreases, the amount of ce11
growth diminishes and the ce11 size rernains smaller. Although the final number of cells is not
mentioned, it is believed that the cell number increased as the injection pressure increased.
This corresponds to the effect of pressure drop rate as explained in Section 2.3.2, A shorter
filling time at a higher injection pressure results in a higher rate of pressure drop, which
favors the nucleation of a larger number of cells.
Ce11 growth has uusally been modeled wiihout the assumption of possible Ioss of
blowing agent. The loss of blowing agent or gas escape is significant in a batch process
although the gris loss is not easily noticed. [n a batch process, the sample skin has the highest
temperature and this prornotes the gas escape to the environment [133]. This is why the
microcellular foarns produced in a batch process have a typically low volume expansion ratio
in the range of 1.5 to 10 times although the maximum possible volume expansion ratio is
about 35 times for 5 wt% CO2 concentration. However, the loss of blowing agent in a batch
process can be minimized by Iowering the foming temperature and thereby Iowering the
diffusion of gas. if the foaming temperature increases too much, the amount of gas escape
becomes significant so that the final foarn expansion will be dtamaticaily reduced 11341.
On the other hand, for foaming in a molding process [131,132], the poIymer in the
mold is mechanicdIy constrained by the mold wails. Therefore, the volume expansion ratio
is determined by the mold cavity to shot size ratio regardIess of the gas escape through the
foam skin.
Lee and Ramesh [l20,135] scudied the effect of gas loss theoretically and
experimentally and showed that the amount of the Iost gas is influenced by the number of
cells nucleated. The lost gas amount decreases with an increase in the number of cells up to
500 cells/cm3, and then becomes almost insensitive above 500 to 2000 ce11slcm3. They
attributed this effect to the fact that less time is available for cell growth and gas escape. As
the number of cells increases, the wali thickness between the cells decreases, and thus the
distance decreases for gas diffusing from the macrix to the cell. As a result, the rate of ceII
growth is faster when the number of cells is greater. Moreover, because a large number of
cells are nucleated, the final ceII size should be smaller. Thus, a cell must grow faster to a
size which is smaller, and so the growth time is much shorter. A shoner growth time allows
less time for gas escape from the sheet surface and chus the expansion increases. Therefore,
the final foam density decreases as the number of ceI1s increases.
Lee et al. [136] studied the effects of foam sheet thickness and nucleation on
thennoplastic foam sheet extrusion. The results show that a decrease in sheet thickness
reduces the final cell size and increases che foam density (Le., decreases the amount of
expansion). They attributed this effect to the influence of heat transfer and also to the
possible gas escape from the surface of the sheet to the atmosphere rit the time of growth. Lee
et al. [136] claimed that at a Iarge thickness, the sheet behaves like an insulator thus
enhancing the bubble growth. As the thickness decreases below 0.5 mm, the amount of final
sheet thickness (or volume expansion) decreases significantly. This effect can be further
described by a surface-to-volume ratio concept- Decreasing the sheet thickness increased the
surface-to-volume ratio of the foam sheet. Since the gas loss occurs at the foam surface, this
Ioss is promoted by increasing the surface-to-volume ratio. Thus, the amount of the lost gas
is greater resultïng in a lower expansion (or a higher foam density) in a thin sheet.
Shafi et al, [137] developed a mode1 that combines nucleation with bubble growth,
They studied the effects of processing variables such as the pressure and the dissolved gas
concentration on bubble growth dynamics, nucleacion and final bubble size distribution
during free expansion polymer processing. The mode1 was based on an expansion influence
volume where each bubble is assigned a region depending on the nucleation threshold, and
this specified volume expands as the bubble grows. They found out that the growth rate
increased with an increase in the solubiIity of giis in polymer. They also claimed that the final
bubble size distribution depends on both nucleation rate and bubble growth dynamics, and
that lowering surface tension significantly increased nucleation rate and resulted in a much
narrower bubble size, a higher bubbIe density, and smaller bubble sizes.
Shimoda et al. [138] conducted a numerical simulation for polymeric foaming
extrusion processes. They combined the classicai nucleation rate and bubble growth rnodels
with a non-Newtonian fluid model of a flow in order to simulate bubble growth and
nucleation in a flow field. They examined the effects of influence volume region and the
initial equilibrium vapor pressure of bubbles on bubble size and number density calculation.
They claimed that the processing parameters such as diffusion viscosity is strongly affecting
the simulation of the ceIl nucleation and growth,
Ramesh et al. [139] developed a new bubble growth model that includes the effect of
blowing agent concentration, temperature effects on physical properties during foam
formation. Experimental data confirmed well the mode1 predictions. Gas loss, blowing agent,
and transient cooling effects are shown to be the most important factors. As found by
Rarnesh et al.. predictions of old foam models differ significantly from experimentally
observed values, since they ignored inff uence of blowing agent concentration effects in the
binary system [139].
in conclusion, ce11 growth is governed by many parameters such as the final foarn
density, ce11 size, and distribution, and can be controlled by the temperature which influences
the diffisivity and melt viscosity. Tempenture control should be performed before ce11
nucleation since the rate of cell growth is much higher in the initial stage where the cells are
small [Il. in the foaming process, it is important to consider the gas loss from the extrudate
dthough the effect of gas Ioss is generally not dnmatic [120,133,136]. However, at a high
melt temperature, where the diffisivity is high and the viscosity is low, gas escape couid be
vigorous. This could result in a dramatic drop in foarn expansion, resulting in an undesirably
high foam density. Moreover, when smailer cells are present, the wall thickness separating
the two cells becomes weaker, and thereby the rate of growth is faster, which may cause
rupture in the ce11 wall and ce11 coalescence [135J.
2.3.4 Crystallization kinetics
Semicrystailine polymers are polymers in which crystallites are dispersed into an
arnorphous matrix. The fraction of the polymer that is fully crystalline is known as the
crystallinity (or the degree of crystallinity). When a polymer is super-cooled by lowering the
temperature below the equiIibrium melting temperature, crystallization takes place and
normally proceeds in two stages: nucleation, and growth. When a long chain molecule srarts
to fold back on itself repetitively, it forms chain-folded lamellae radiating from a nucleus and
nucleation is said to have occurred. The folding length of the crystal is temperature
dependent and increases with the crystallization temperature [140,141]. The broad range of
melting temperature observed in semicrystalline polymers depends on the crystallinity
distribution, and is attributed to the distribution of chain folding length.
During nucleation, the crystallites are nucleated from the melt at a definite range of
temperature and subsequently grow during the growth phase to forrn three-dimensional
aggregates of crystallites known as sphenilites. In generai, the spherulite growth rate is faster
than the nucleation rate because the free energy requirement of the former is Iower [L42].
Nucleation is therefore the rate-determining step of polymer crystallization. When
crystallization occurs under stress, as in the extrusion process, the overall crystallization rate
will increase because the stress could cause orientation of the molecular chains thus making
them more packable [142].
Nucleation can occur in one of two ways; when nuclei appear instantaneously at the
beginning of the process, then athermal or instantaneous nucleation has occurred and it is
assumed to depend only on temperature and to be independent of time and cooling rate. The
grown crystais will therefore be of approximately equai sizes. if nuclei appear in the liquid
phase during the process, thermal or sporadic nucleation has occurred, and the activated
nuclei appear at a constant rate per unit volume.
Nucleation and Growth Rates
In gnerai, two types of nucleation are defined for polymer crystallization: primary
and secondary nucleation. During primary nucleation, when a potentiai nucleus reaches a
critical size, cqstai growth occurs quickly and spontaneously and a three dimensional crystal
is genented, the rate of which depends on temperature. Pnmary nucleation is caiied
homogeneous nucleation if no preformed nuclei or foreign surfaces Xe present. Secondary
nucleation then follows, where chah segments are dded to the existing crystal surface. The
essential difference between primary and secondary nucleation is the energy of formation of
a nucleus of critical size or the Gibbs free energy,
The classical nucleation theory developed by Gibbs [143-1451 is based on an
assumption that energy fluctuations in the supercooled phase can overcome the nucleation
barrier caused by the surface of the crystal. Based on this assumption, Turnbull and Fisher
[146] derived an expression to determine the primary nucleation rate as a function of the
crystallization temperature, using the Williams-Lmdel-Fe;erry (WLF) Il471 equation which
universally describes the temperature dependence of pdymer melt viscosity:
Based on the surface or secondary nucleation theory, Lauritzen and Hoffman [148-
1501 formulated a linear growth rate equation, which incorporates fold surface energy, lateral
surface energy, heat of fusion, and lameIlar thickness ternis into the Gibbs free energy to
describe the linear growth rate of spherulites.
Crystullizution Regimes
The relative rates of nucleation and deposition of chain segments on the crystd
surface will affect both the crystal growth rate and the sphemlite size. Hoffman [151.1521
ctassified several crystailization regimes, which describe the relacive nucleation and growth
rates. A regirne transition occurs when the ~lationship becween growth rate, G , and the
surface nucleation nte, i, undergoes a change. In regirne I, the highest temperature regime.
the chain segment deposition rate on one surface nucleus is so fast, that the crystal unit
growing on an existing crystd face is complete before the next layer is nucleated, in other
words, G varies as i. In regime il, the nucleation rate is fast compared to the growth rate,
: M ir. G = i . consequently ihe nucleation of new crystal layers occurs before deposirion on
existing Iayers is complete. This results in a downward break in the growth rate curve as one
passes through the regime 1 to regime II transition. Finaily, in the Iowesc temperature regirne,
Oregime ID, the mean separation of the nuclei approaches the width of rhe molecular stems,
and G = i again, such that at the regime II to regime III transition, an upward break in the
growth rate curve occurs. Other kinetic theories of crystailization were fomiulated by Frank
and Tosi [ 1531, Sanchez and DiMarzio [154], Sadler [155,156], and Lauritzen, DiMarzio and
Passaglia 1157-1591.
Avrami's ïïleory
Knowledge of the primary nucleation rate and the linear growth rate is usually
sufficient to calculate the overall crystallization rate. Many rneasurernents of crystallization
invoIve the macroscopic determination of crystallinity as a hnction of time. The first efforts
to quantitacivefy describe the rnacroscopic development of crystallinity in tems of nucieation
and linear crystal growth were made by Kolmogoroff 11601, Johnson and Mehl [161], and
Evans [162]. However, the classical theory of Avrami [163-1651 for phase transformation
kinetics is the most widely quoted mode1 For the analysis of isothermal nuclerition and
crystallization in polymer processing.
Despite 'its wide use, the Avrami mode1 suffers from some limitations caused by
simplified assurnptions. These assumptions are as follows: (i) there is no volume change
during crystallization; (ii) the sampie is completely trrinsformed; (iii) there is constant linear
growth rate; (iv) the nuclei have constant shape during growth; and (v) there is no secondary
crystallization occurs,
The implications and complications caused by these assumptions are dealt with in
detail in the literature [166]. It is sufficient at this point to state that macroscopic observation
of the increase in crystallinity cannot adequately descnbe the rnicroscopic mechanisms of
crystal gowth. The Avrami analysis sirnply provides a convenient representation of the
macroscopic data.
Diflerential ~canning Calorimeter
DSC is a powerful tool in the determination of the Avrami parameters as desccïbed in
his theory. Determination of polyrner crystallinities using the heat of fusion is often based on
rneasuring the area of DSC rnelting peaks above a chosen baseline. During the
crystdlization of a sernicrystdline polymeric material, heat is rejected as the rnotecules
become ordered and form crystallites. The heat evoIved (Mldt) is rneasured and recorded as
a function of time by the DSC, and the weight fraction XJt) of materiai crystallized aiter
rime t cm be cdculated from the relation:
where X,,(t) is the absolute crystallinity at crystallization time t, and X, is the ulcimate
crystallinity for t==. For isothermal crystallization experiments, the heat evolved is
evaluated while the polymer is in the isothermal condition, consequently, XJt) c m be
physicaliy obtained as the area under the crystailization peak in a plot of heat tlow versus
cime. Since the Avrami equation is expressed in terms of the volume fraction, it is necessary
to transform the weight fraction measured by DSC into a volume fraction. This can be done
using the following relation:
where p, is the morphous region density, and p , is the crystdline density.
in conclusion, the Avrami ana1ysis is a useful tool for the representation of the
macroscopic data of polymer crystallizrition. Based on this analysis, the main parameters
governing the crystdlization kinetics under isothermal conditions are the primary nucleation
rate and the spherulitic growth rate. Physical investigation of the crystallization phenomena
can also be conducted using DSC method by obiaining the heat evolved as a function of time
under various isothemd conditions. The resulting crystaltization kinetics cm be used as a
bais for establishing strategies for the production of low-density, fine-celied potypropylene
foarns.
Figure 2.1: Sorption isotherm for general gadpolymer system
A
Free Energy,
AG
+ Bubble Radius, r
Figure 2.2: Effect of the variation of energy on the bubble growth for homogenous nucleation
polymerlgas solution
Figure 2.3: Model of a nucleated ceIl inside a polymer matrix
Chapter 3
Design & Construction of the Experimental Equipment
3.1 Conceptual Design of Tandem Extrusion System
3.1.1 Introduction
An axiomatic design method developed by Suh at the Massachusetts Institute of
Technology [167] was employed to create a conceptual design of the tandem extmsion
system. in brief, this design method starts with the definition of Functional Requirements
(FRs) from the perceived needs; this defines the design problem. With the FRs, one can
generate a number of possible physical entities corresponding to each FR. These physical
entities are called Design Parameters (DPs). The retationship between the FRs and the DPs
can be expressed by a matrix equation:
{ FRs } = [A] { DPs} , (3-1)
The elernents in the A matrix cm be either "X", which denotes a suong relationship between
the corresponding FR and DP, or "O", which denotes a weak or absent relationship. The A
matrix presents the cornplete retationships between the FRs and the DPs; it can help the
designer better understand and improve the design.
3.1.2 Analysis of the Tandem Extrusion Systern for Foam Processing
The four functionai requirements (FRs) identified for producing low-density, fine-
celled polyrner foam are as follows:
RI = plasticarion of polymer;
FR2 = formation of a polymer/gas solution;
FR3 - - nucleation; and
Flt - - expansion.
To satisfy the above FRs, the following design parameters (DPs) are proposed:
DPI = heat provided by the plasticating screw motion in an extruder and
externally mounted band heaters;
DP2 = gas injection, diffusion system and gear pump;
DP3 = thermodynamic instability created by a nuclerition die: and
DP4 = a cooling system that can cool the polyrner melt without decreasing
its pressure.
in response to FRi, a single-screw extruder is chosen to plasticate the polymer. There are
two heat sources in the system. The primary source is the frictional heat generrtced by the
motion of plasticating screw. The secondary source is the externaily mounted band heaters.
The band heaters are important in the start-up period but they are not suitable as the only heat
source for this process due to the low thermal conductivity of the polymer.
FR2 is satisfied by using a gas injection system with a metering device, and a
diffusion enhancing device to produce a homogenous polymerlgas solution. The amount of
gas is metered by a gas injection pump, which supplies the gas under high pressure into the
polymer melt in the extruder. The resuIt is a two-phase polymerigas mixture, The shear field
generated by the motion of plasticating screw stretches the gas bubbles and increases the
interfacial surface area when the polymer is conveyed in the extruder barrer. As a result, the
gas can diffuse into the polymer matnx more quickly. To further assist the diffusion process
and homogenize the polymer/gas solution, a diffusion enhancing device consisting of static
mixers is employed.
in response to FR,, a rapid pressure drop is chosen to generate thermodynamic
instability and thereby promote high nucleation density. The thermodynamic instability is
created by a sudden drop in gas solubility in the polymer/ gas solution. As mentioned in
chapter 2, the gas solubility in polyrner is proportional to the pressure of the polymer melt. A
rapid drop in the pressure results in a rapid decrease in the solubility of gas in polymer melt.
Ttierefore, a high nuclei density can be created by dropping the pressure of the polymer/gas
solution rapidly.
A coaling system is employed to satisfy FR+ The most dominant parameter that
affects the growth of the nucleated cells is the melt temperature. As the gas from the
polyrner melt diffuses into the nucleated cells, the concentration of the dissolved gas
decreases in the region near the cells. This results in a concentration gradient in the melt,
which drives the dissolved gas towards the cell. The diffusion of gas into the cells causes the
cells to grow. The gas diffusion rate increases with temperature, and therefore, if the melt
temperature is too high, the gas cm easily escape from the polymer to the environment
instead of contributing to cell growth. A high temperature can also promote ceIl coalescence
since the melt strength decreases when the temperature increases. As the cells grow, the
walls between them will be stretched and could easily be broken due to the weak melt
strength. Thus, adjacent cells will join together, and the ceIl structure and the cell-population
density wilI deteriorate. Since gas diffusion and ce11 coalescence can be controlled by
lowering the temperature, a cooling system is chosen for satisfying m. A cooling system
consisting of a second extruder is introduced.
The FRs-DPs relationships cm be described in the following matrix:
The diagonal elements Aii of the design matrix are al1 "X", simpIy because each DPi
is chosen directly to accomplish the corresponding FRi- An examination of al1 the non-
diagonal elements of the matrix is required in order to determine the effects of each DPi on
the ather FRs.
Since gas is injected in the polymer after the polymer is cornpletely plasticated, it has
no effect on the plastication process. As a result, element AL2 should be zero. EIernents AL3
and A14 should aiso be zero because nucleation, expansion and shaping take place further
down strearn and therefore, will have no effect on the plastication stage. The shear field
generated by the plasticating screw motion affects the polyrnerlgas mixing. Therefore,
elernent should be non-zero. Element A3 should be zero because nucleation takes place
after formation of the polymerlgas solution. The amount of gas in the polymer melt is
dependent on De4. The addition of a gear pump decouples FR3 frorn these DPs. The gear
pump ensures a constant polymer flow rate as long as the gear speed and the extruder speed
are not varied. This feature implies that the amount of gas injected into the polyrner meIt,
once set, will becorne independent of the temperature variations and the die exit settings.
Therefore, elernent AZ4 is zero.
The spindle speed of the plasticating screw affects the systern pressure and therefore
affects the nucleation rate. As a result, element Afl should be non-zero. Element A3? should
be non-zero too, because the arnount of gas in the polymerlgas solution affects the nucleation
density. Elernent A3J should be zero because the cooling system of a second extruder can
cool the polymer rnelt without sacrificing its pressure.
As desciibed previously, the temperature of the polymerlgas solution affects the
volume expansion ratio. Thus, hI should be non-zero because the shear heat generated from
the piasticating stage affects the temperature of the polymer flow, Element should be
non-zero since the arnount of gas injected affects the expansion ratio. Elernent should be
non-zero because gas loss to the environment is localized if the cell-population density is
high.
From the above considerations, Equation (3.3) can be written as:
X X X O DP, X X X X - DP,
Equation (3.4) is a lower triangular rnatrix, which irnpties that the functional requirernents
can be achieved if the design parameters are implemented in the proper sequence.
Having arrived at a satisfactory decoupled design at the parent levet, it is necessary to
decompose these FRs and DPs to lower level hierarchies and zigzag between the functional
domain and the physical domain to arrive at an appropriate design for satisfying the initial
FRS.
3.1.3 Detailed Analysis and Further Decomposition of the FRs & DPs
Decomposition of FR4 & DP, (Second Level)
The cooling systern is required to perform two main functions. FirstIy, the melt
temperature has to be uniformly lowered to the optimum level, without losing the processing
pressure, in order to prornote ceIl growth without coalescence. Secondly, the surface
temperature of the extrudate has to be further reduced to prevent the gas from escaping
prernaturely to the atrnosphere through a weak surface layer. Therefore, the two FRs for this
syseem are:
FR,, = Cool the polyrner melt homogeneously to the optimum level
without losing the processing pressure;
FR,? = Cool the surface of the extrudate to prevent escape of gas to the
atmosphere.
In response to FR4,, a cooling extruder and a cooling section are incorporateci. The
shear mixing in the extruder can hornogenize the rnelt temperature whiIe cooling it, and the
thrust force provided by the screw motion can maintain the processing pressure. The cooling
section has a sleeve through which a cooling fluid flows to cool the polymer funher. The
static mixers in the cooling section ensure the homogeneity of temperature in the melt.
Although there is sorne pressure loss due to the cooIing in the cooling section, the high
pressure provided by the cooling extruder can compensate for this loss. To cool the surface
of the extrudate, an appropriate system is chosen. The DPs are then given as:
DP,, = Cooling extruder and cooling section;
DP,? = Polyrner surface cooling system.
The design equation can be represented as:
A,? is zero because the surface cooling cannot affect the inside temperature of the melt due CO
the low thermal conductivity. However, if the melt temperature is highec than the desired
temperature, the cooling of the surface cooling system will have to be increased ro achieve
the required surface temperature; this leads to the conclusion that A,, is non-zero. Equation
(3.6) wiIl become:
Therefore, the cooling system design is decoupled.
The polymer surface cooling system can be further decomposed to the third Ievel in
order to arrive at a proper design.
3.2 Detailed Design of the Tandem Extrusion Line
3.2.1 Ovewiew of the System
Based on the conceptual design of the overall system for low-density, fine-celled
foam extrusion in the previous section, the design of the detailed components is carried out in
this section. The tandem extrusion system consists of two single-screw extruders, gas
injection equipment, a gear pump, a diffusion enhancing device, a heat exchanger, and a
fitament die. The first extruder is used for plasticating the poIymer min, the gas injecrion
equipment is used for injecting gas into the polymer melt, while the second extruder provides
mixing and initial cooling for the polymer mek The gear pump controls the polymer melt
fIow rate, independent of temperature and pressure changes, the diffusion enhancing device
ensures the homogeneity of the polymer/blowing agent soIution and the heat exchanger
provides further cooling for the polymer melt to suppress cell coalescence. Shaping and ceIl
nucleation are accomplished in the die. Figure 3.1 shows a schematic of the tandem extrusion
system. The detailed design of each component is presented in the following sections.
3.2.2 The First Extruder in the Tandem Line
The extruder used in the single screw extrusion system consists of a W1 laboratory
extruder (Brabender: 05-25-000) with a 5 hp variable speed drive unit (Bnbender: Prep
Centre, Model D52T). The screw is a single stage mixing screw (Brabender: 05-00-144)
with a 30:l UD ratio. The purpose of the mixing stage is to enhance the mixing of the
blowing agent and the polymer melt. A schematic of the mi~ ing section is shown in Figure
3.2.
3.2.3 Gas Injection Equipment
The two main components of the gas injection equipment are a positive displacement
syringe pump and an in-house designed gas injection port. The pump is capable of injecting
the blowing agent into the polymer melt at a maximum pressure of 51.7 MPa (7500 psi) with
a wide range of flow rates ( fom O.OI mllmin to IO7 mI/min, depending on the pressure
condition).
At the hem of the gas injection port is a flow restrictor. Because of the compressibility of
gas, the pumping of gas into the barre1 is readiIy affected by the pressure fluctuation in the
barrel. Variations in the injection rate could affect the consistency of foaming significantly.
One solution is to maintain a high pressure difference between the gas injection pump and
the barre[. The choice of restrictor depends on the required pressure difference and the
desired gûs flow rate.
3.2.4 The second Extruder in the tandem Line
The second extruder in the tandem extrusion system consists of a 1%" extruder
(KiIlion: KN-150) with ri buïit-in 15 hp variable speed drive unit. The screw motion c m
generate shear heat in the polymer melt, however. since the second extruder is intended for
cooling, it is necessary to keep the generated shear heat to a minimum. In this context, a
smdl iength/diarneter (Ml) ratio of the screw (18:l) was chosen in the screw design.
Moreover, a compression ratio of 1: 1 was chosen since the second extruder is intended for
rnaintaining the pressure of the polymer melt.
Polymer rnelt is fed into the barre1 of this extruder from the firsc extruder through an
injection port. Typically, the injection pressure of the melt is around 27.58 MPa (4000 psi).
Because of the high pressure, it is necessary to have a dynamic sealing in the second extruder
to prevent the melt Leaking backward into the motor assernbiy through the clearance between
the screw and the barrel. The proposed solution was to miike extra flights on the screw
Iocated before the injection port. As the screw rotates, the thrust force generrited by the
screw motion of these extra fiights would push the polymer melt forward preventing the
leakage to the back. A detailed design of the second extruder was proposed by Young [I68].
3.2.5 Gear Pump
A gear pump (Zenith: PEP-I l), with a '/2 hp motor drive (Pacific Scientific: Model
SR), a speed control unit (Zenith: ZeDrive) and a temperature controller (Eurothem Controls:
Mode1 94), is used in the extrusion system to control the polymer solution flow me. The
gear purnp consiscs of two closely intermeshing gem that rotate in a counter-rotating manner
to convey the polymer melt from one end to the other. A schematic of the gear purnp is
shown in Figure 3.3.
3.2.6 Diffusion Enhancing Device
A diffusion enhancing device is used to ensure the polymer melt and the blowing
agent are mixed hornogeneously. It was originally designed by Park [169j and modified by
Behravesh [170], It consists of an in-house design, rniId steei body, a static mixer (Omega:
FMX844 lS) enctosed in a mild steel case, and two band heaters (Omega: MB 1GlJ I A t &
MBIG2A1 Al). The rationale behind this design is to use the static mixer to promote shex
mixing, and to maintain a high melt temperature to promote a high difision rate of the
blowing agent into the polymer rnelt.
in order to determine the required number of static mixer elements, a calculation was
done based on an equation provided by the manufacturer's technical brochure [L? LI:
where Re = Reynolds number,
Q = flow rate, gallmin
S = specific ,mvity
p = viscosity, cP
d = inside pipe diameter, in
A typical flow rate (Q) of the system is in the range of 5-12 cm3/min (1.3-3.2xl0-' gallmin).
The specific gravity (S) is approximately 0.9 for polypropylene, and the inside pipe diameter
(6) is 0.0127m (0.5 in).
The viscosity is influenced by the shtar m e and temperature. The apparent shear rate
( j,,, ) in a circular channel can be calculated by the following equation [17 11:
With a channel diarneter (4 of 0.07 m (0.28 in) and a flow rate (Q) of 10 cm3/rnin (2 .64~10~~
gdrnin), the apparent shear rate is approximately equal to 4.95 11s. From the manufacturer's
data sheet (Borealis), the apparent viscosity of bmched polypropylene is 916.108 Pa's
(916108 cP). With these values, the Reynolds nurnùer can be cakulated using Equation
(3.7):
The pressure arop across the static mixer c m be detennined using the following equation
[171]:
where I = laminar factor (from the manufacturer), and
A P = pressure drop, psi.
With a flow rate (Q) oF0.303 m3/min ( ~ . ~ L c x I o - ~ gal/min), a viscosity (p) of 916108 cP. and a
larninür factor (0 of 0.05, the vaIue of head loss in the static mixer is:
AF = 120.92 psi (0.8 17 MPa) (3. L 1)
This value is reasonable when compared to a typical system opemting pressure of 27.58 MPa
( 4 0 psi). Based on the calculated results, it can be concluded that a static mixer with six
elements cm fulfill the rnixing requirernents (1701, Additional mixing of the polymer melt is
provided by the mixing stage of the extruder's screw.
3.2.7 Heat Exchanger
As mentioned in Chapter 2, cooling the polymer meIt will increiise the melt strength
and suppress ce11 coalescence. It is important to cool the polymer melt homogeneous1y
because non-uniforrn temperature distribution could induce in-homogeneous ce11 growch,
resuIting in a irregular cell structure. The heat exchanger utiiized in this research was
designed by Behravesh [170]. It consists of a static mixer (Labcore: H-04669-12) encased in
a mild steel body with embedded cooling channels as shown in Figure 3.4. This static mixer
is stntcmrally different €rom the one used in the diffusion enhancing device. This static mixer
used in the herit exchanger promoted polymer transport in the radiai direction such that the
core material is constmtly exchanging with the boundary materiai, whereas the mixer of the
diffusion enhancing device does not.
Since the temperature cm directly affect the foaming behaviour of the polymer, it is
important to control the temperature of the heat exchanger precisely. Using band heaters
alone with temperature controllers do not provide adequaie control because of lack of a
cooting source, consequently, a high pressure gas is introduced CO provide cooIing for the
heat exchanger. The temperature controller controls the air flow through the cooling channel
in the heat exchanger using a solenoid valve based on the polymer melt temperature. When
the temperature is above the set point, the solenoid value opens to let the high pressure air
expmd and flow through the chmnel in the heat exchanger. Because of the cooling induced
during isentropic expansion [172], the temperature of air is reduced significantly and the cold
air removes the heat from the heat exchanger. When the temperature is below the set point,
the band heaters are switched on to provide the necessary heating. Using this band
heatedhigh pressure air arrangement, temperature control of &loC was easily maintained,
3.2.8 Filament Die
The principal function of the filament die is to induce a thenodynamic instability,
which promotes nucleation of a large number of cells. The section of the die, which induces a
rapid pressure drop, is characterized by a small diamcter and is referred to as the nucleation
section or capillciry section. In this case, two nucleation dies were designed to induce a rapid
pressure drop. For polypropylene/butane solutions, it was found that a die of 0.46 mm (0.018
in.) diameter and 7.62 mm (0.3 in.) length, and a larger die of 0.76 mm (0.030 in.) diarneter
and 12.6 mm (0.5 in.) length, resulted in a desirable processing pressure in the processing
temperatures that are of interest. These filament dies were therefore manufactured since the
foaming behavior of the polyrner rneIt solution at different temperatures was not known.
The effects of die diameter and length on the pressure drop and the pressure drop rate
have been explriined in details in the literature [169,170]. It is clear that a change in diarneter
affects the pressure drop rate, while a change in diameter does not influence the pressure
drop rate. However, both the die diameter and length affect the pressure drop in the die. For
instance, by decreasing the die length without changing the die diarneter, the pressure drop
decreases while the pressure drop rate remains the same. Thus, a shorter die may be used
when a high-pressure drop is experienced during processinp. Figure 3.5 show a schematic of
the filament dies used.
3.2.9 Cooling Sleeve
As discussed earlier, the strategy for preventing the gas escape is to cool the extrudate
surface whiIe it flows into the nucleation section of the filament die, This cm bbe achieved
through cooling of the die using high pressure air, to cool the entire nucleation section of the
die. Thus, there is a need to provide a seded channel around the die to circulate the
pressurized air. It would be inefficient to make a channel for each die. especiaIly when a
large number of dies are to be exiunined. An alternative design is to make a sleeve wirh a
large circumferentiai grwve, as shown in Figure 3.6. When the sleeve is mounted on the die
and sealed at each end using two O-rings, the cooling air can be circuhted througti the
channel.
Figure 3.1: Schematic of the tandem extrusion system
Figure 3.2: Schematic of the extruder mixing section
Figure 3.3: Schernatic of the gear pump (counesy of Zenith).
hot polymer melt -
cooling oil cooling
static mixer channel \ oil pressure transducer
seat
cotd polyrner melt
+
cooiing oil
high ternphanire thermocoup1e O-ring seat
Figure 3.4: Schematic of the heat exchanger (courtesy of Behravesh) 11701
Figure 3.5: A schematic of the filament die
Heat exchanging air grooves for
nucleation ,' nozzle \ - -
- -
circulation
Figure 3.6: A schematic of the cooling sleeve to be mounted on the nucleation nozzle
Investigation of Fundamental Properties of Polypropylene Materials with ~oaming Agents
4.1 Introduction
This chapter presents the investigation of various processing and materials parameters
on the themodynamic, thermal and melt fracture behaviors of polypropylene melts with
foaming additives under different processing conditions. This will aIlow us to develop the
strategies for achieving low-density, fine-celled polypropylene foams and to identify the
fundamentai mechanisms governing the volume expansion ratio of polypropylene foarns in
the next two chapters.
4.2 Measurements of PVT Properties of Polypropylene Materials
In this section, the effect of dissolved butane, processing temperature and pressures,
and materials bnnching on the PVT relationships of polypropylene matenals in a moIten
state are investigaced. The basic principle involved in the measurement of PVT properties of
polyrner/gas solutions is to measure the specific voIume of the solution by deterrnining the
mass and volume flow rates of the polymerlgas solution at ditierent temperatures and
pressures. In this context a dilatometer based on a foaming extruder with a new degassing
oven to facilitate the measurernent of the mas flow rate of polymedgas solutions from the
foaming extruder is presented [173]-
The system utiIized for the measurement of the PVT properties of propylenehutane
solutions is based on the previous system developed by Park et al. [60,61]. The previous
system was found to be inadequate for experiments involving butane. Since, the diffusivity
of COz in the polyrner is much higher than butane, the degassing of COz can be more easily
accomplished, For the case of butane, the gas does not emily go out of the polymer becriuse
of the low diffusivity of butane. An effective degassing system was cherefore designed and
implemented to achieve more accurate degassing of the butme from the propylene foams.
Xthough detailed information about the system cm be found in the previous pubrications
[60,6 11, the overall function of the foaming extruder-based dilatometer, and the procedure to
determine the specific density of propylenehutane solutions are briefly described below.
An important element in the measurement of the PVT relationships of the
propylenehutane solutions is the formation of a homogeneous and uniform single-phase
polymerlgas solution in a themodynamically well-defined pressure and temperature
condition. This is achieved through the use of a tandem foarning extruder shown in Figure
4.1. Melting of the propylene is achieved in the first extruder through the frictional shear
generated by the screw motion and through the use of extemally mounted band heaters.
While in the molten condition, gas is injected into the polymer in metered amounts using a
syringe pump. As the propylenehutane mixture is conveyed in the extruder barrel. the shear
fields generriteà by the plasticating screw and irregular mixing blades stretch the gas bubbles
under high pressure and result in dissolution of the gas into the polymer matrix [17,18].
After formation of the single-phase polymerlgas solution. the temperature of the solution is
controlled in the second extruder and further homogenized using a heat exchanger consisting
of a static mixer. A positive-displacement gear pump is then used to control the volume tlow
rate of the polymerlgas solution. A variable resistance valve is installed after the gear pump
to control the pressure of the solution flowing through the gear pump. The solution then
exits throush a small filament die, which facilitates collection of samples for analysis.
The specific volume of the propylene/butane solution foned at each temperature and
pressure was then determined using Equrition 4.1.
where subscripts p and g refer to the polyrner and gas respectively, and Q ~ ~ ~ ~ , , ~ ~ ~ ~ , , ~ ~ ~ ~ , and
mpçx(,,ighZ highP, refers to the volume and mass fiow rates of the propylenehutane soIution,
mesured in cm31g and glmin, respectively. The volume flow rate of the solution is
determined by:
The RPM of the gear pump was set constant for each series of experiments, whiIe the
throughput per unit revolution of the gear pump was determined from calibration
experiments performed by Park et al. [60]. The m u s flow rate of the propylendburane
solution was dererrnined at high experimentnl temperatures and pressures using Equation 4.3.
- mp+,v (IiighT. hjgh PI - mp lhighT. high PI
' ',g(higfr T. high P)
where rhp(highT.highP) represenrs the mass flow rate of the polymer in the polymerlgas
solution, determined with the aid of the degrissing system descnbed. The m a s flow rate of
the gas in the polyrner/gas solution, mX(,lighT,highP), was computed as:
- high T. high P) - mg.k;ispump(mom T.injcction P)
- - Qg.gu pvmp (momï'.injeciion PI
Pi(rwm~.injcction P)
The volume flow rate of gas in the :as pump.~,,pu,p ,,,,,,, lionPl, was collected for each
experiment by reriding the gris injection pump, whiie the gas density at room temperature and
gas injection pressure, pi (momZinjcnianP), was obtained from the thermodynamic tables for
butane [174]. The validity of equiition 4 holds only when the equilibrium point is reached.
The system was considered to be at equilibrium when there was no fluctuations in the gas
flow rate readings of the gas pump.
Since gas loss from the extruded foarn is unavoidable [20,2I], simpIy weighing the
collected extnrdde foam btown with butane woutd not provide an accurate measmement of
the mass Bow rate of the polyrner mett or the poIymerhutane solution. The rnass flow rate of
the solution c m then be determined by Equation 3, using the sepmteIy measured gas and
polymer m a s flow rates. Therefore, it is proposed that the mass flow rate of polymer melt
only be measured by substantially degassing the extruded polymerlbutane solution insrmd of
rneasuring the mass flow rate of the solution.
The principles of Axiomatic Design [167] were employed to design a special
degassing oven to remove al1 traces of the blowing agent from the polymerigas solution (see
Figure 4.2). The main design considerations were: (i) to devoid the polymerlgas solution
sample of gas and (ii) to determine the rnass of the sample. To achieve these objectives, a
heating mechanism and a weighing device were chosen, respectively. The main requiremenr
for the heating mechanism was to achieve and maintain high temperatures above 200°C. To
satisfy the requirement, a physical enclosure with appropriate insulation was designed ta
provide a finite space for heating, a heat source was chosen with the appropriate range, a
thermocouple and temperature controller were used to maintain the set temperature, and air
circulation was provided to maintain temperature uniformity. It was known that any
weighing device would be adversely affected by very high temperatures, and would result in
incorrect mus readings, Consequently, physical separation of these rwo functions was
desirable. A balance with the desircd resolution and capacity was chosen. The balance couId
take measurements in one of two modes: above or below the balance. It was determinrd that
physical separation would best be achieved by positioning the balance above the heating
mechanism. The overail design was then constructed to specifications, and tested for
functionality.
4.2.1 Experimentai Equipment and Procedure
4.2.1.1 Experimental Setup
The PVT measurement system for the propylenehutane solution consisted of a
tandem foaming extrusion line described in chapter 3 equipped with a variable resistance
valve, a filament die (see Figure 4.1), and a degassing system (see Figure 4.2)- An explicit
description of this setup can be found in the ceferences [60,6 11.
An integral part of this study involved using a degassing system to achieve maximum
gas loss from the foam sampIe. As shown in Figure 4.2, the degassing oven was divided into
two separate sections. The lower section was used for heating and degassing the propylene
foarns and consisted of a cenrnic heater (Omega: CRFP-126/120A). tnsulation was provided
for the walls in the form of glas wool fibers. The wall directly opposite the ceramic heater
was hinged to the frame to allow access to the heated chamber, and a circulation h n was
inserted through one wall. The upper portion of the oven was left exposed to the
environment with a simple platform on the top. A high-resolution balance (Scientech:
1 1 144- 12) equipped with beIow balance weighing function was mounted on this platform. A
weighing assembly was constructed of a small diarneter aluminum rod and thin duminum
plate. The rod section of this assembly was attached to the bottom of the scale to enable
below-balance weighing. The plate section was positioned in the oven to support the sample
during the degassing process. Finally, a Type J thennocouple was positioned in the lower
charnber and the temperature controlled using a temperature controller (Fuji: P m - T A Y 1-
4w-
4.2.1.2 Experimenîai Materiais
A linear and a branched polypropylene material with MFRs (ISO 1133, 230 OC/2.16
kg) of I l dg/min and 2.3 dg/min, respectively, were chosen as the poIymer materials for the
experimentation in the presect study. Both polypropylene materials were supplied by
Borealis AG Austria. The blowing agent used in this study was n-butane supplied by
Matheson Gas Company.
4.2.1.3 Experimental Procedure
The band heaters located dong the tandem foaming extmsion line equipped with a
gas injection system were turned on for approximately half an hour before running the
system to ensure proper plristication of the propylene materials that were charged in the
extruder through a hopper. Initially, the speed of the calibrated gear pump was set and the
system was started. While the system ran, butane was injected from the gas pump, and the
ga was mixed with and dissolved in the polymer melt in the first extruder. Then, the molten
polymerlgas mixture advanced through the second extruder, the heat exchanger, the g a r
pump, and exited the system through the filament die. Before the measurements, the system
was kept running for 20 minutes or more to reach a steady-state condition. The weight of
butane with respect to the weight of the polymer was maintained constant throughout each
experiment to ensure consistency. This required repeated measurements of the mass flow
rate of the polyrner as well as the mass flow rate of butane being injected by the gas pump.
The temperature of the filament die was maintained to be 220 O C throughout the experiments.
Since measurements had to be taken at a thermodynamically defined pressure and
remperature, control of these two parameters was essential to the measurement procedure.
The pressure upsueam of the gear pump was controlled by varying the rotationd speed of the
second extruder, while the downstream pressure was controlled using a variable resistance
valve 1601. The temperature of the rnelt was controlled in the second extruder using band
heaters placed around the barrel. The temperature was then Further homogenized unifonnly
as it passed through the static mixer, while sirnultaneously being cooled.
At the instant when the inlet and the outlet pressures of the gear pump were equalized
and no fluctuations in the pressure readings were observed, extruded polymerlgas solution
samples were collected for 1 minute at the die exit. Three separate sets of samples were
collected. After collection, the samples were placed in the degassing oven CO Facilitate
degassing of the sample. After 8 minutes (which represents the degassing time described in
the next section) the mas of each sample was recorded and the volumetric and mass flow
rates of the polymerigas solution determined. Finally, the specific volume of the
propylenelbutane solution was calculated based on Equation 1.
4.2.1.4 Calibration of the Degassing Oven
in order to reduce the error involved in merisuring the polymer tlow rate. the time
required to degas the polymerlgas solution sample was detennined frorn calibration
experiments. It should be noted that if the degassing tirne was too short, then the residud gas
would cause an error. The degassing oven was therefore caiibrated to detennine the
maximum time required for cornplete degassing of the samples. Extruded polymerigas
samples were collected at each experimentai condition and placed in the oven. The mass of
each sample was recorded at one-minute intervals. Of ail the cdibration experiments
perforrned, it was obsewed that the Iongest degassing times required occurred ar
expenmentd conditions of 170 O C , and 15% butane content. The result is depicted
graphically in Figure 4.3. It was observed that a npid decrease in the m m occurred beyond
a certain time. This is beiieved to be due to the degradation of the plastic materiai at the hi$
temperature. The threshold value was 8 minutes and this was selected as the maximum time
required for degassing during the experiments at ail conditions. It should be mentioned that a
significant amount of degassing of the propylene foam occurs immediately after exit from the
die, because of the small cross-section of the sample and the high temperature of the filament
die, which facilitates desorption of gas, however a foam structure is still evident. After
degassing in the oven, the foam structures disappeared, and only few bubbles of very small
sizes could be seen. Degassing was assumed complete, when the mass remained unchanged
for three consecutive readings.
42.2 Results and Discussion
4.2.2.1 Effect of dissolved butane on the specific volume
Figures 4.4 (a) and 4.5 (a) represent the effect of varying the amount of dissolved
butane in the solution on the specific volume Cor linear and bnnched propylene, respectively.
For both linear and branched propylene, the specific volume linearly increased as the amount
of injected butiuie increased, When the butane content was increased from O wt% to 15 wt%.
the specific volume increased by spproximately 02798 cm31g for linear propylene. This
increase was approximately constant over the temperature range 170°C to 210°C and the
pressure range 13.8 MPa to 37.6 MPa.
It is believed that, due to the high solubility of butane in propylene, al1 of the injected
butane was dissolved in the polymer, This reduces the solution density and consequently
increases the specific volume. For O wt% butane (pure propylene melt), the specific volume
increased by 0.0222 cm31g over the range of temperature from 170°C to 210°C. This was
found to be consistent with published resuIts 1601. With 15 wt% butane, the specific volume
increase was only 0.0045 cm3ig. This change seems to be reduced for temperature, however,
with respect to pressure, the sensitivity became higher. Over the temperature range 180°C to
220°C. Park et al. 1601 obtained a specific density increase of 0.050 cm31g when the CO2 was
increased from O wt% to 4 wt% for the PSICOl system. For the propylene/butane solution,
the specific volume increased approximately by 0.100 cm3tg when the butane content was
increased fom O wt% co 4 wt%. For the same wt% of gas, the swelling was therefore
observed to be higher for butane, This would indicate that butane has a higher plasticizing
effect than CO?.
4.2.2.2 Effect of branching on the specific volume
A significant difference between the specific volume behaviors o f linear and
branched propylene materials was observed (Figures 4.4 and 4.5). This cornparison reveaIs
that at rnost experimental conditions seiected, the specific volume was found to be higher for
the linex propylene material. This is of interest, because for polyethyiene, branched
materials are known EO have Iower density by occupying more volume. For the branched and
linear propylene materials of interest, the densities ac room temperature were almost the same
(0.91 @cm3). However, when the temperature was increased to the high range of 170°C to
190°C, the specific volume of the branched propylene was lower. When butane was
dissolved into the propylene matrix, the polymer swelled; it is believed that in branched
propylene, the long chain branching restricted the mobility of the molecular chahs. As the
butane dissohed in the propyIene matrix, the difference between the specific volumes of
linear and branched propylene resins increased, especially at high pressures. It was dso
noted that the branched propylene materiai showed a higher sensitivity with respect to the
change in pressure compared to the Iinear resin. The higher sensitivity of the branched resin
with respect to the pressure became more severe as the gas concentration increased. It is not
dear why the specific volume of branched resin shows a very high sensitivity with respect to
the pressure at hi& butane concentrations. A further study is required to clarZy this issue.
4.2.23 Effect of processing temperature and pressure on the specific volume
The graphs presented in Figures 4.4 (b) and 4.4 (c) itlustrate the effect of the
processing temperature and pressure respectively, on the specific volume of
propyIenehutane solutions for five different butane contents (Le., O%, 2%, 5%, IO%, and
15%) for linea. propylene. Sirnilar results for branched propylene are shown in Figures 4.5
(b) and 4.5 (c). These resuits reveal that at O wt% butane, the sensitivities of the specific
volume with respect to the temperature and pressure were similar to the results obtained by
others [57,58]. However, when butane was dissolved in the propylene melt, the specific
volume became more sensitive to a change in the pressure for both materials, whereas the
specific volume with respect to temperature was not changed much.
4.2.2.4 Error Analysis
Due to the dynamic nature of PVT measurements, a number of sources of error were
possible. These include the errors associated with the residual gas, and the controi and
measurement of temperature, pressure, and gas concentration. It is believed that the error
associated with the residual gas in the polymer matrix was negligible because of the newly
designed degassing system. Evidence in support of this assumption could be observed in the
degassed polymerlgas solution, as only very few gas bubbles of very small size could be seen
in the degassed propylene~butane solution. On the issue of temperature measurement and
control, examination of the results reveals that the average sensitivity of the specific volume
with respect to temperature was approxirnately 0.005 crn3/g0~. The error range associated
with the temperature control of our systern was detemined to be t 2°C [60,61]. This
corresponds to an error of t 0.0 10 cm31g in the specific volume, With regards to the pressure
control and measurement, our system error was determined to be approximately k 0.69 MPa.
From our results, the sensitivity of the specific volume with respect to pressure was
approximately 3.600~10-~ c r n 3 / g - ~ ~ a . The error range in the specific volume associated
with the sensitivity due to pressure is Cherelore 2.500~10-~ cm3/g. The average sensitivity of
the specific volume with respect to the butane content was found to be relatively high at
0.020 cm3/g-butane wt%. However, ctiere is no possible way of measuring the final butane
content in the extmded propylene foam CO determine our systern error with regards to the
butane content. Further research and development are therefore needed to quantify and
minirnize the error in the specific volume associated with the butane content. Since the
butane content of the foam is also inextricably Iinked to the butane flow rate. improvements
to the system must also be made in terms of accurate gas flow rate control.
4.2.3 Conclusions
In this study, the effect of dissolved butane on the PVT relationships of linear and
branched propylene materials in a molten state were investigated. Based on the experimental
results, the following conclusions can be drawn,
1. The specific volume of the propylenehutane solution increased significantly with an
increase in the percentage of butane injected in both branched and linear propylene. Both
the linear and branched propylene resins sweHed significantly as the butane permeated
into the resins.
2. At al1 experimental conditions selected, the specific volume was found to be higher for
the linear propylene than for the branched propylene.
3. When butane is dissolved in the propylene matrix, the sensitivity of the specific voIurne
with respect to pressure increased with the butane content for both the linear and
branched resins, whereas there were no significant changes in the sensitivity with respect
to the temperature. EspeciaIly, the branched resin exhibited a very high sensitivity with
respect to the pressure at a high concentration of butane.
4.3 Measurements of Thermal Behavior of Polypropylene Materials
The thermal behaviors of linear and branched poiypropylene with foaming additives
were investigated using a high-pressure differential scanning calorimeter (DSC).
Specifically, the effects of material branching, dispersed additives, and dissolved blowing
agent on the crystallization temperature of polypropylene resins were elucidated. The effect
of dissolved blowing agents was determined using the high pressure DSC cell with carbon
dioxide and nitrogen. The effect of hydraulic pressure was identified by performing DSC
study employing an inert gas such as helium, which has a very low solubiiity in the polymer
matrix. Foaming additives such as tdc and GMS as welI as processing parameters such as the
cooling rate dso played major roies during the crystallization process [175],
4.3.1 Experimental Equipment and Procedure
4.3.1.1 Experimental Setup
The crystallization experiments were perfonned w ith a DSC (TA Instruments, DSC
29 IO), A regular DSC ce11 and a high-pressure DSC cell were used in the experiments.
4.3.1.2 Experimental Materials
A linear and a branched propylene material were chosen as the polymer materials.
They are the same materials used in section 4.2. Talc (A7 with top-cut of 7 microns,
Naintsch) and glycerol monostearate (GMS, Pationic 909, PATCO Polymer Additives) were
used as foaming additives in our investigation. HeIiurn (BOC Gas 99.9 % purity), nitrogen
(BOC Gas 99.9 % purity) and carbon dioxide grises (BOC Gas, 99.5 % purity) were used in
high-pressure experiments.
4.3.1.3 Experimental Procedure
The regular DSC ce11 was used for investigating the effects of branching, additives
and cooling rates on the crystaliization behavior of polypropylene materials. The high-
pressure DSC cell was used for investigating the effects of dissolved gases on the
crystallization behavior of polypropylene materials. The calibration of both DSC cells was
done using indium. Samples for DSC experiments (typicai weight 3-4 mg) were taken from
the extrudate in a form of a very thin disk (typicai thickness 150-200 Pm) [176]. For the
nonisothermal experiments, the smples were heated up to 220 OC and kept at this condition
for 30 minutes to ense the thermal history [89], and then the samples were cooled down at
10 'Clmin (if not specified) to 50 O C . Next, the samples were heated at 10 'Clmin up to 200
OC. During the cooling and heating processes, the crystallization and melting patterns were
recorded- In the case of high pressure DSC ceII, the sarnples were pressurized to 1.37, 2.75,
4-13 and 5.5 1 MPa using helium or nitrogen or carbon dioxide.
4.3.1.4 Design of cooling system for the high pressure DSC ceIl
In order to investigate the effects of dissolved gas on the crystallization behavior of a
poiyrner, a high-pressure DSC ceIl (TA Instruments, DSC 2910) was used for this
experiment. Due to the poor cooling capability in the pressure DSC cell. only the heating
mode has been mainiy used in the high-pressure experiments.
A cooling system was developed as shown in Figure 4.6 to provide the existing high-
pressure ce11 with a cooling capability. The cooling systern enabled the measurernents of the
crystallization behavior of a polymer melt under high pressure by mounting a cooling coi1
connected to a liquid nitrogen dispensing unit which supplies liquid-nitrogen at a controlled
rate. Pnor to the cooling coi1 modification, the high pressure DSC could only attain non-
uniform cooling rates o l 10 "Clmin. Critical experiments were conducted to verify the
pressurizing and cooling functions of the modified pressure cell. The pressure was raised
inside the modified high-pressure ce11 from atmospheric pressure to 3.51 MPa, and at this
pressure a uniform cooling rate up to 30 "Clmin was successfully achieved. Also, the
crystallization behavior of polypropylene materials measured with the high-pressure DSC
ce11 was compared to that with a regular ce11 operating at atmospheric pressure. The
crystallization thermograrns of polypropylene materiais were determined at atmospheric
pressure using both the DSC cells and were found to be nearly identical at the same cooling
rates.
Finally, using this design, the measurements of crystallyzation kinetics under high
pressure was successfully carried out using the modified high-pressure DSC ceIl.
4.3.2 Regular DSC Celf Results
4.3.2.1 Effect of Branching
Figure 4.7 shows the cooling sections of the DSC thermograrns obtained for linear
and branched propyle - a materials without any additives at a cooling rate of IO 'Chin. It was
observed that branching of propylene chahs significantIy promoted the crystallization
kinetics of polypropyiene resins by increasing the crystaliization temperature about 20 OC.
This result supports the findings of previous studies [177,178], Figure 4.7 also shows that the
peak of linear polypropyiene materiais has a shouIder in the peak (or double peak) caused by
two-stage crystallization [66,179,180]. It can be understood that the interactive motion of the
polypropylene matrix layer at the particle surface changes the crystallization speed of
polypropylene matrix [66].
4.3.2.2 Effects of Foaming Additives
The thermal behaviors of linear and branched propylene resins with foaming
additives such as talc and GMS were also investigated as a part of this study. The
experiments were conducted at a cooling rate of 10 "Clmin. The concentrations of talc and
GMS were changed from O to 1.6 wt% and O to 1 .O wt%, respectively.
Figure 4.8 (a) shows the effect of talc amount on the crystallization temperature of
polypropylene materials. The effect of talc was more dominant in Iinetu polypropylene
material than in the branched one. After showing a sharp increase in the crystallization
tempenture as the talc concentration increased from O to 0.2 wt%, the crystallization
tempenture did not change much above 0.2 wt%. Figure 4.8 (a) shows that the crystallizacion
tempenture of branched polypropylene material is only 5-10 OC higher than that of the linear
material in the actual foam processing with the talc content ranging from 0.8-1.6 wt9.
However, without the addition of talc particles, the crystallization ternperiture of brrinched
polypropylene material is 10-20 "C higher than that of linear one. The degrees of crystallinity
of linear and branched propylene resins were also measured and the results are shown in Fig.
4.8 (b). It was observed that the degrees of crystallinity of branched and linear propylene
materiais increased moderately as the talc concentration increased.
Very sirnilar results were obtained for the case of GMS (Figs- 4.9 (a) and 4-9 (b)),
The GMS, used as an aging modifier, also increased the crystalIization temperatures and the
degrees of crystallinity of Iinear and branched propylene resins. However, because of its
strong tendency of migrate to the surface of the extrudate 118 11, the uniformity of GMS in
the polypropylene matrices could not be ensured. It is believed that the variations shown in
Fig. 4.9 (a) are due to the non-uniform distribution of GMS particles. The DSC samples for
this study were cut perpendicularly to the flow direction from a thin filament extrudate (about
4 mm in diameter) to minimize the effect of non-uniformity of GMS particles in the radiai
direction of filament extrudate.
4.3.2.3 Effect of Cooling Rate
The effect of cooling rate on the crystallization kinetics of linear and branched
propylene resins was also investigated. The cooling rate was varied from 0.1 Wmin to 50
"Clmin. Figure 4-10 shows the dependence of crystallization temperature on the cooling rate
for Linear and branched propylene materials. It was observed that the crystallization
temperatures of polypropylene resins were very sensitive to the change of cooling rate: the
crystallization temperatures decreased by 25-30 OC as the cooling rate increased from 0.1
"C/min to 50 'Clmin.
4.3.3 High-Pressure DSC Cell Results
Figures 4.1 1 and 4.12 show the effects of pressure on the crystallization behriviors of
linear polypropylene materials (Fig. 4.1 I ) and branched polypropylene materials (Fig. 4.12)
with various grises. The dependence of crystallization temperature on the pressure was quite
different for different gases. It was obvious that the gas at high pressure significantly affected
the crystallization kinetics of polypropylene materials. As a gas permeates into the pulymer
rnatrix under high pressure, the dissolved gas may change the rate of polymeric segmental
motions such as rearrangement into crystals. Furthemore, since the amount of gas dissolved
in the polymer increases with an increase in pressure [106], the magnitude of the change in
the plasticization and crystallization will be more pronounced at a higher pressure.
On the other hand, the crystallization kinetics of the polymer melt will aIso be
affected by the hydnulic pressure externally applied by the gas. When the solubility of gas in
the polymer is very low or negligible, the kinetics of crystallization will be governed by the
hydraulic pressure. But when the solubility of the gas in the polyrner is considerably high,
both the hydraulic pressure and dissolved gas will play individual roles during the
crystailization process in high-pressure experiments. The effect of dissolved gas on the
crystailization kinetics cm then be extracted by subtracting the effect of the hydraulic
pressure from the overall crystallization behavior of the material under high pressure.
4.3.3.1 Effect of Hydraulic Pressure
The effect of hydraulic pressure on the crystallization behaviors of polypropylene
materiais was determined from the high-pressure expeciments with helium (He). Since the
solubility of He in a polymer is very low, typically one order of magnitude lower than that of
NI and two orders of magnitude lower than that of CO? [106], the effect of the dissolved He
in a polymer melt would be negligible. Therefore, any change of the crystallization kinetics
under an elevated pressure of He can be considered as the effect of hydraulic pressure.
Figure 4.13 shows the onset and peak crystallization temperatures as a function of the
He pressure. It was noted that by increasing the pressure in the DSC cell, thc onset and peak
crystallization temperature increased for both linex and branched propylene resins: the onset
and crystallization peak temperatures for branched polypropylene material increased by
about 6 O C when the pressure was increased from the atmospheric pressure to 5.5 MPa, but
for linear polypropylene material, the temperatures increased by only about 3 OC for the same
amount of pressure change. As the hydraulic pressure in the DSC cell increases, the rnobility
of the polymer matrix molecules decreases, and hence accelerates the crystallization process,
resulting in a higher crystallization temperature. Figure 4.13 also shows that the increase of
the crystallization temperature for the branched polypropylene material was even more
pronounced than that of the linear material.
The crystaIlization temperature shown in Fig. 4.13 was not changed much above 2.75
iWa and no further effect of hydraulic pressure was observed at higher pressures. it was not
clear if the plateau region observed was due to the dissolved He in the polypropylene
materiais at elevated pressures, or if it reflected the actual effect of hydraulic pressure on the
crystallization behavior. Once the solubility of He in polypropylene resins is measured, this
issue will be clarified.
4.3.3.2 Effect of Dissolved N2
The effects of dissolved N2 on the crystallization behaviors of Iinear and branched
propylene materials were extncted by subtracting the hydraulic pressure effect (Fig. 4-13]
from the overail thermograrns (Figs. 4.1 1 and 4.12). Figure 4.14 shows the resuIting onset
and peak crystallization temperatures as a function of the Nz pressure. N2 has a relatively
higher solubility than He 11061, and therefore, the effect of dissolved gas with N? on the
crystallization kinetics of polypropylene materials will be more pronounced when compared
to the case of He.
It was observed that the onset and peak crystallization temperatures did not change
much at pressures below 1.4 MPa. Since the solubility of & in the polymer is very low at a
low pressure, the overall crystallization behavior in the low pressure range was mainly
governed by the hydraulic pressure effect as shown in Figs. 4.1 1 and 4.12. The effect of
dissolved N2 on the crystallization was more pronounced at a higher pressure (above 1.4
MPa), the onset and peak crystallization temperatures decreased moderately up to 5.5 1 MPa,
It should be noted that these compensated results shown in Fig. 4.14 may have some
errors due to the dissolved He in the polypropylene rnatrix at high pressures, causing an
incorrect hydraulic pressure effect. However, it is believed that the error range would be
srnail.
4.3.3.3 Effect of Dissolved COt
As in the case of N:, the effects of dissolved CO2 on the crystallization behaviors of
linear and branched propylene materials were extracted by subtracting the hydraulic pressure
effect (Fig. 4.13) from the overall thermograrns (Figs. 4.1 1 and 4-12). Figure 4. t 5 shows the
onset and peak crystallization ternperatures for the linear and branched polypropylene
materials as a function of the CO2 pressure. Tt was observed that the dissolved CO2
suppressed crystallization of the polypropylene materials significrtntly. The decreased
crystallization temperatures due to the dissolved CO2 was more pronounced in the
thennogram of branched polypropylene material cornpared co the linew one.
At high pressure (5.51 MPa), the increased crystallization temperature due to the
branched structure was compensated for by the dissolved CO? and the crystallization
temperature of branched polypropylene material becarne almost the same as that of the Iinear
one. This indicates that the effect of branching on the crystallization kinetics may be
neglected at elevated CO2 pressures above 5.5 1 MPa. However, because of the limitations of
our high-pressure equipment, experiments at pressure higher than 5.51 MPa coutd nor be
performed. On the other hand, the crystailization temperature of Linear Pt decreased Iineariy
as the pressure increased.
Tt is believed that the Iarger magnitude of change in che crystailization of
polypropyIene resins under a high pressure with CO: compared ta the case of NI_ wris due to
the solubility difference of both gases. Because of the higher solubility of CO? compared to
N2 [182], the effect of dissolved CO, was more pronounced even at a low pressure. A large
amount of dissoIved CO2 increases the f ee volume of the polymer [183], and enhances the
mobiIity of polyrner chah The resulting crystallization temperature of polypropylene
materials is therefore lowcnd.
4.3.4 Conclusions
A series of experiments were conducted to investigate the effects of materia1
branching, foarning additives, cooling rate, hydraul ic pressure, and dissolved gas on the
crystdlization behaviors of poiypropylene resins. Helium was empioyed to estimate the
hydraulic pressure effect on the crystallization behavior for the high-pressure experiments
with N2 and CO2 blowing agents. The f o a i n g additives considered in this study included
taIc ruid GMS. The experiments conducted in this study tead to the following conc1usions:
1. Branching in the poiypropylene rnatrix caused a significant increase in the
crystallization temperature.
2. The foaming additives such as talc and GMS increased the crystalIization temperature
of polypropylene materiais.
3. The crystallization temperature was a sensitive function of the cooling rate and it
decreased as the cooling rate increased.
4. Crystailization of polypropylene materials was enhanced as the hydraulic pressure
increased-
5. But the dissolved I$ and COz lowered the crystailization temperatures of
polypropylene resins. In particular, high-pressure COz decreased the crystailization
temperature significantiy because of the high sohbility.
4.4 Measurements of the Onset of Melt Fracture of Polypropytene Materials
The melt fracture behaviors of Iinear and branched polypropylene resins with
foaming additives were investigated. The effects of branching, processing temperature,
additives, and blowing agent on the surface melt fracture of polypropylene materiais were
thoroughly studied. A CCD carnera was installed at the die exit to precisely observe the onset
of surface melt fracture of extruded foams. The critical wall shear stress was determined for
various linear and branched polypropylene resins using a capillary die [184].
An experimental setup is designed to investigate the surface melt fracture behavior of
polypropylene materials melts and polypropylene/butanc solutions under various conditions,
Figure 4.16. This system is based on the tandem extrusion systern described in Chapter 3.
Due to the difficulties involved in determining the onset of melt fracture for polymer
foam, a visual approach is employed. A CCD camera is instalted at the die exit and the foam
extrudate is carefully monitored. In the case of foaming extrusion, the skin of extruded foam
is stretched as the expansion occurs. As a consequence, the foarn skin becornes shiny as
alortg as the expanded foarn does not contract due to gas loss [20]. Even if melt fracture
occurs and sharkskin developed on the extrudate surface, the foam expansion of extrudate
causes the foam skin to be stretched, and thereby, the trace of the sharkskin gets easily
removed, Therefore, it is vrry difficult to detect the onset of melt fracture by sirnply checking
the surface of the fully expanded foam. However, it is believed chat the onset of surface melt
fracture cm be observed by monitoring the early stage of extrudate using a CCD camera
before it gets expanded.
A capiIlary die of length L and radius R is used to calculate the critical shear stress at
which meIt fracture occurs [185]. A pressure transducer is mounted in the die More the
capiIIary section. in order to minimize the entrance effect, the UD ratio was chosen to be 35.
and the wall shear stress is calculated in terms of the die pressure, Pd, and the die geometry
[ISS]:
The corresponding wall shear rate, yUPP, is determined by 11851:
The polyrner flow rate is controlled by the rotational speed of the gear pump regardless of
the temperature and pressure fluctuations in the barrels. Appendix 1 shows a detailed analysis
of the Flow in capillary die.
4.4.1 Experimental Equipment and Procedure
4.4.1.1 Experimental Setup
Based on the above design, an experimental setup is constructed to study the effects
of branching, processing ternpenture, additives, and blowing agent on the surface rnelt
fracture behaviors of polypropylene materials melts and polypropylenehutane soIutions. The
setup consists of the tandem extmsion system described in Chapter 3, a ccipillary die of O. I O
cm diameter and 3.55 cm in length, and a cooling sleeve for the precise control of dit:
temperature. A CCD camera (Pulnix) is mounced at the exit of the die and connected to a
computer processor in order to accurately visudize and monitor the onset of rnelt fracture.
4.4.1.2 Experimental Materials
The materials used in this scudy were two linear standard polypropylene resins and
three high-rnelt-strength (HMS) branched polypropylene resins supplied by Borealis AG.
They are denoted in this study as Linear PI, Linear P2, Branched Pl, Branched P2, and
Branched P3, respectively. The materials properties, inciuding MFRs (ISO 1133,230 'Cf?. 16
kg), molecular weights and numbers ( M W , Mn), and the degrees of long chah branching per
1000 carbon atoms (LCB), are summarized in Table 4.1. The foaming additives used in this
study were talc and glycerol mono stearate (GMS) as the cell-nucleating agent and the aging
modifier. respectively [18 11. The blowing agent used in the experiments was n-butane, C.P.
(Matheson, 99.0%).
4.4.1.3 Experimental Procedure
The polypropylene resins were processed in the extrusion setup and the onset of me[t
fracture was investigated by obsewing the die pressure, the surface qua1it.y of the extrudate,
and the image of the extrudate €rom the CCD canera.
Firstly, experiments were conducted without gas injection for investigating the effects
of branching, melt temperature, and foarning additives on the melt fracture. Brrinched
materials with various degree of long chah branching were used to invescigate the effect of
brmching on the melt fracture. For investigating the effect of processing temperature, the
melt temperature was precisely controlled using the tandem extrusion system. Vanous
amounts of talc and GMS were added to the polypropylene materials to investigate the
effects of foaming additives on the surface melt fracture of the extrudate, The onset of
surface melt fracture was determined by directly obsewing the surface quality of the
extrudate and also by analyzing the CCD image of the extmdate. As the die pressure
increased by increasing the gear pump speed, the occurrence of melt fracture was detected.
When melt fracture occurred, the critical shear stress was calculated by reading the
corresponding die pressure.
Secondly, experiments were conducted with gas injection to study the effect of
dissolved gas on the melt fracture. A metered amount of blowing agent was injected and
dissolved into the polypropylene melt to form a single-phase polypropylene/butane solution.
The fomed single-phase polypropylene/gas solution entered the die and foaming was
allowed. The onset of melt fracture on the foamed surface was carefully observed using the
CCD camera as the die pressure was increased with a higher speed of gear pump.
4.4.1.4 Calibration of the CCD camera image for detecting the onset of surface melt
fracture
In order to precisely deterrnine the effect of dissolved gas on the onset of s~irface melt
fracture of polypropylene materials, the image of the extrudate captured with a CCD camera
For the early stage of foaming was carefully analyzed. It was obsewed that for a pure
polymer melt without any dissolved gas. and without melt fracture the extrudate was steadily
exiting the die. In this case, the extrudate shape was uniform and smooth. However, at the
onset of surface melt fracture, the extrudate corning out of the die started to oscillate
verticaiiy in a direction perpendicular to the flow. It was also confirmed that surface
irregularities started to appear on the surface of the extrudate at this moment. Since the
amplitude of oscillation was very smdl at the onset of surface melt fracture, the oscillation of
extrudate could be detected only by the magnified image of the extrudate from the CCD
camen.
On the other hand, when gas was dissolved in the polyrner and the expansion of
extrudate occurred due to foaming, the appearance of surface melt fracture on the foam
surface was not synchronized with the moment when the foam extrudate started to oscillate.
in other words, even when the minute-scale oscillation of the extrudate was detected from the
captured CCD image, there was no visible change on the extrudate surface and the degree of
straightness of foam extrudate. It was believed that surface melt fracture actually occurred at
the moment of oscillatory motion of the extrudate. However, the trace of melt fracture on the
foam surface must have been removed because of the stretching of foam skin during
expansion. The foam skin at this moment was typically very shiny. By contrat, when the
degree of oscillation of the extrudate became Iarger so as to be visible even to the naked eye,
the foam extrudate became wavy. However, the foam surface was still shiny and smooth.
In conclusion, the onset of surface melt fracture could be detected effectively using a
CCD camera by observing the image of the extrudate shape as it starts to oscillate for both
pure polyrner and Foam.
4.4.2 Results and Discussion
4.4.2.1 Effect of branching on the critical shear stress
The effect of branching on the surface melt fracture behavior of polypropylene resins
is shown in Figure 4.17, without the use of additives. It was observed that the critical shear
stress linearly decreased with an increase in the degree of long chah branching, in other
words, the surface melt fracture of polypropylene resin was promoted by the degree of
branching, According to the experimental results, the critical shear stress decreased
approximately by 0.0175 MPa with an increase of 0.1 nJ1000c in long chain branching The
decrease of criticaI shear stress may be attributed to the increase of melt elasticity as the
degree of long chah branching increased. An increase in melt elasticity can increase the slip
at the die wall, and hence promote surface melt fracture [6 11.
4.4.2.2 Effect of processing temperature on the critical shear stress
The effect of processing temperature on the surface melt fracture behavior of
polypropylene resins is shown in Figure 4.18 for al1 the linear and branched materiais, In this
experiment, the processing temperature was varied from 180 OC to 210 OC. It was observed
that the critical shear stresses of linear and branched polypropylene resins were aImost
insensitive to the processing temperature in the range of 180 OC to 210 OC.
The critical shear rate at the onset of surface melt fracture was also calculated as a
function of the processing temperature using Equation 4.6. Figure 4.19 shows that the
apparent shear rate at the onset of surface melt fracture linearly increased as the temperature
increased from 180 OC to 210 OC for al1 the polypropylene resins. The shear rates behavior
observed in this study conform to the behavior obsewed in other studies [95,96].
4.4.2.3 Effects of foaming additives on the critical shear stress
The effects of the dispersed foaming additives, Le., talc and GMS, on the critical
shear stress of polypropylene resins were also investigated. All the experiments were
conducted at 190 OC. Figures 4.20 and 4.2 1 depict the critical shear stress for Linear P 1 and
Brrinched P 1 materials as a function of the talc and GMS contents. respectively. For both talc
and GMS cases, it was observed that the critical shear stress of Branched PI increased
sharply as the amount of dispersed additive increased from O to 0.4 wt%. A further increse
in the arnaunt of dispersed additives only increased the critical shear stress slightly. By
contrast, the critical shear stress of Linear PI did not change much with the foaming
additives. For both talc and GMS, the critical shear stress for Linear Pl increased sli~htly
(less than 0.005 MPa) as the concentration of additives increased from O to 2.4 wt%.
It is interesting to note that the well-known lubricating effect of GMS was not
distinguished in this study [186]. On the other hand, the lubricating effect of talc, as a ceII
nucieating agent, for the branched polypropylene materials was also surprizing whereris the
Iubricating effect of boron nitride, which is another well-known ce11 nucleating agent in f om
processing, has been reported before [174]. It is speculated that there may be some
rdationship between the cell nucleating ability of the additive and its lubricating effect.
Further study needs to be conducted to c1arifL this issue.
4.4.2.4 Effect of blowing agent on the critical shear stress
The effect of the dissolved blowing agent on the criticai shear stress of polypropylene
resins is shown in Figure 4.22. It was observed that the critical shear stress increased
significantly as the amount of dissolved butane increased: for both linear and branched
polypropylene resins, the criticai shear stress increased by 0.025 MPa as the butane content
increased from O to 20 wt%. These results rnean that the surface rnelt fracture is significantly
suppressed by the dissolved butane. The decreased criticai shear stress of Branched PL by
branching was alrnost recovered by the dissolved 20 wt% butane. On the other hrind, the
critical shear stress of Linear Pl was also significantly increased by the dissolved butane.
4.4.3 Conclusions
Experirnental studies were camed out to investigate the effects of branching,
processing temperature, foaming additives, and blowing agent on the critical shear stresses of
linear and branched polypropylene resins. An experimental setup was designed to elucidate
the surface rnelt fracture behaviors of polypropylene melts and polypropylendbutnne
solutions under various conditions. Efforts were made to accurately control the processing
temperature, to well disperse the foaming additives in the polypropylene melt, and to fom a
single-phase polypropyleneibutane solution. A CCD camera wris installed at the die exit to
precisely monitor and analyze the onset of surface melt fracture of the extrudate at the early
stage of foam processing. From the expenrnents conducted in this study the following
conclusions cm be drawn:
1. An on-line technique for detecting the onset of surface rnelt fracture for extmded foam
has been developed by visualization of extrudate using a CCD camera.
2. The long-chain branching of polypropylene rnaterials significantly decreased the critical
shear stress of the resins.
3. The critical shear stress was aimost insensitive to the die temperature; however, the die
temperature significantly affected the criticai shear rate at the onset of surface mett
fracture.
4. The foaming additives of talc and GMS increased the critical shear stress of branched
polypropylene materiats. However, they did not affect the critical shear stress of Iinear
polypropylene rnaterials much.
5. The dissolved butane significantly increased the critical shear stresses of linear and
branched polypropylene resins.
4.5 Summary
Based on the results shown in this chapter, the effects of processing conditions and
foarning additives on the amount of swelling of polypropylene melts, the c~stallization
point, and the onset of melt fracture were identified. The specific volume was highly
dependent on the processing temperatures and pressures, the degree of branching, and the
dissolved butane. The crystallization temperature of polypropylene materials was a sensitive
function of branching, foaming additives, dissolved b1owing agent, and the cooling rates. The
onset of melt fracture was affected by the degree of bnnching, the addition of foaming
additives, and the dissolved butane. These parameters are the bais for development of the
strategies for the production of low-density, fine-celled polypropylene foams, and for
identifying the fundamental mechanisms that governs the volume expansion ratio of
polypropylene foams.
Table 4.1. Materials properties of Iinear and bcanched polypropylene materials
- Material
Branched P 1
Branched PZ
Branched P3
Linear P 1
Linear P2 I
k[FR
(gllornin)
2.3
4.8
3.1
I I
2.8 I
LCB degree
(n/lûûûc)
0.2 1
0.17
O. 13
O
O I
MW
(Wmol)
418
416
327
360
500 I
Mn
(KgImol)
3 1.7
48
46.5
70
97 I
Figure 4-1: Schematic of the tandem extrusion Iine that provided the extruded samples for measuring the PVT propenies of polypropylene/butane solutions
Figure 4.2: Schematic of the degassing oven 1. High Resolution Balance 2. Heater 3. Below Balance Weighing 4. Sample Holder 5. Circulating Fan
8 5 I
1 2 3 4 5 6 7 8 9 t O
the (mh)
Figure 4.3: Calibntion of the degassing oven
Butane Content (%)
(a) Specific volume vs butane content for Iinear polypropylene
12 14 16 18 20 22 24 26 28 Pressure (MPa)
(b) Specific volume vs pressure for linear polypropylene
Figure 4.4: Changes of the PVT data of Iinear propylene
Temperature (C)
(c) Specific volume vs temperature for linear polypropylene
Figure 4.4: Changes of the PVT data of linear propylene (cont'd)
1.20 C 1
O 2 4 6 8 10 12 14 16 Butane Content (%)
(a) Specific volume vs butane content for branched polypropylene
J L.
12 14 16 18 20 22 24 26 28 Pressure (MPa)
(6) Specific volume vs pressure for branched poIypropylene
Figure 4.5: Changes OF the PVT data of branched polypropylene
1.20 . 170 190 210
Temperature (C)
(c) Specific voIume vs temperature for branched poiypropylene
Figure 4.5: Changes of the PVT data of branched polypropylene (cont'd)
I 1
4.6 Design of Cooling Capability for High-Pressure DSC Cell
Crystallization hmperature (C)
Figure 4.7. DSC thennograms of Iinear and branched polypropylene resins
+Eranched P l (onst) +ûranched Pl (peak) +Linear P l (onal) +Linear P l (peak)
100 ! 1
Concentration of Talc (%)
Figure 4.8 (a) Effect of talc on the crystailization behaviors of linear and branched polypropylene resins
0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0
Concentration of Talc (%)
Figure 4.8 (b) Effect of talc on the degrees of crystdlinity of linear and branched polypropylene resins
-6-Branched P l (onsht) . -a-Branched P l (ptak)
+ Linear P l (peak) , 100 .
0.0 0.2 0.4 0.6 0.8 1 .O 1.2
Concentration of GMS (%)
Figure 4.9 (a) Effect of GMS on the crystallization behaviors of linear and branched polypropylene resins
Concentration of GMS (%)
Figure 4.9 (b). Effect of GMS on the degrees of crystdlinity of Iinear and branched poiypropylene resins
Cooling Rate (Clmin)
Figure 4.10. Effect olcooling rate on the crystdlization betiaviors of linear and branched polypropylene resins
i oo J I
O 1 2 3 4 5 6
Pressure (M Pa)
Figure 4.1 1. Effect of pressure on the crystallization behavior of Linear polypropylene
Pressure (MPa)
Figure 4-12. Effect of pressure on the crystallization behavior of Branched poiypropylene
--
n n w O w
1 - -e-Branched P l (onat) +Branched Pl (peak) :
I i -e- Linear P l (onmt) 1 +Linear P l (peak)
100 4 O 1 2 3 4 5 6
Pressure (MPa)
Figure 4.13. Effect of hydraulic pressure on the crystallization behaviors of linear and branched potypropylene resins
Figure 53 (a): the expansion ratio versus the butane content for branched polypropylene materials using butane as biowing agent
Figure 5.3 (b): the biowing agent efficiency versus the rnei t temperature for branched polypropylene materials using butane as biowing agent
Figure 5.3 (c): the ce11 density versus the meIt temperature for branched polypropylene materiais using burane as blowing agent
Figure 5.3 (d): the ce11 density versus butane contents for branched polypropylene materids
Figure 5.3 (el: the die pressure versus the meIt temperature for branched polypropylene materials using butane as blowing agent
- t 4 50 C
g 40 W
30
20
10
O
100 120 140 160 110
Tempantun (C)
Figure 5.4 (a): the expansion ratio versus the melt temperature for linear polypropylene materials using butane as blowing agent
100 120 11P 1W 180
Ternpmtun (C)
Figure 5.4 (b): the blowing agent efficiency versus the melt temperature for linear polypropyIene materials using butane as blowing agent
1W 120 110 160 180
Tmnptalun (C)
Figure 5.4 (c): the ceii density versus the meIt temperature for linear polypropyiene materials using butane as blowing agent
O 5 10 16 N ZS
Butina &nmi(%l
Figure 5.4 (d): the cell density venus butane contents for Iinear polypropylene materials
Figure 5.4 (e): the die pressure versus the meh temperature for linear polypropylene materials using butane as blowing agent
= ~~r-~7- -tlO%COI - - - +i% COI
2 44% COI
-. 100 120 140 160 110
T i m p i n i u n (C)
Figure 5.5 (a) the expansion ratio versus the melt temperature for branched polypropylene materials using CO, as blowing agent
Figure 5.5 (b) the blowing agent efficiency versus the meIt temperature For branched polypropylene materials using CO, - as bIowing agent
Figure 5.5 (c) the cell density versus the melt temperature for branched polypropylene materials using CO, - as blowing agent
Figure 5.5 (d) the ce11 density versus CO2 contents for branched polypropylene materials
Figure 5.5 (e) the die pressure versus the melt temperature for branched polypropylene materials using CO, as biowing agent
Temperabin (C)
Figure 5.6 (a) the expansion ratio versus the melt temperature for branched polypropylene materiais using a blend of butane and CO, as blowing agent -
Figure 5.6 (b) the blowing agent efficiency versus the melt temperature for branched po1ypropy:ene materials usinp a biend of butane and CO, as blowing agent
Figure 5.6 (c) the cell density versus the melt temperature for branched polypropylene materials using a blend of butane and CO, - as blowing agent
lf*0 O 20 40 60 80 100
RihUvr CO2 Contint h W n (?&)
Figure 5.6 (d) the cell density versus CO2 relative amount for branched polypropylene materials using a blend of butane and CO, as biowing agent -
Figure 5.6 (e) the die pressure versus the meIt temperature for branched potypropylene materials using a blend of butane and CO, as bIowing agent
t O d X Talc
10 -1.6% Talc - c Z . 4 % Talc
"- 60 à 1 : 50 - I
g 40
30
20
10
O 100 120 140 160 110
Temperature (C)
Figure 5.7 (a) the expansion ratio versus the melt temperature for branched polypropylene materials using butane as blowing agent for various talc contents
Figure 5.7 (b) the blowing agent efficiency versus the melt temperature for branched polypropylene materials using butane as blowing agent for various talc contents
100 120 140 160 110
Timprnhirr (C)
Figure 5.7 (c) the cell density versus the melt temperature for branched polypropylene materials using butane as blowing agent for various talc contents
1.E4 ! l.E+OQ 4
O M 1 1 . 5 2 2 5
Talc Conion! ( X )
Figure 5.7 (d) the ceil density versus the melt temperature for branched polypropylene materials using butane as blowing agent for various talc contents
Ttmprnlurr (C)
Figure 5.7 (e) the die pressure versus the melt temperature for branched polypropylene materiais using butane as blowing agent for various talc contents
Drawdown velocity in mmb
Figure 5.8 the melt strength and melt extensibility of linear and branched polypropylene materials blend
Figure 5.9 (a) the expansion ratio versus the melt temperature for linear and branched polypropylene materiais biend using butane as blowing agent
Figure 5.9 (b) the blowing agent efficiency versus the meIt temperature For linear and branched polypropylene matecïds blend using butane as blowing agent
Figure 5.9 (c) the ce11 density versus the melt temperature for linear and branched polypropylene materials blend using butane as blowing agent
Figure 5.9 (d) the ce11 density versus linear polypropylene content for Iinear and branched poIypropylene materials bIend using butane as blowing agent
Figure 5.9 (e) the die pressure versus the rnelt temperature for linear and branched polypropylene materiais blend using butane as blowing agent
Figure 5.10 (a) the expansion ratio versus the melt temperature for re-extruded Brancheci PI polypropylene materials using butane as blowing agent
Figure S. 10 (b) the blowing agent eficiency versus the melt temperature for re-extruded Brmched P l polypropylene materials using butane as blowing agent
Figure 5.10 (c) the ceIl density versus the melt temperature for re-extruded Branched Pl polypropylene materials using butane as blowing agent
Figure 5.10 (d) the ce11 density versus the butane contents for re-extmded Branched P 1 polypropylene matends
Figure S. 10 (e) the die pressure versus the melt temperature for re-extruded Branched P 1 poIypropylene mateciais using butane as blowing agent
100 120 140 160
Ternpwrlurr (C)
Figure 5.1 1 (a) the expansion ratio versus the melt temperature for re-extruded Branched P2 polypropylene materials using butane as blowing agent
Figure 5.11 (b) the bIowing agent efficiency versus the melt temperature for R-extruded Branched P2 polypropylene materiaIs using butane as blowing agent
i oa 120 140 160 180 t r m p r i u i n (C)
Figure 5.1 I (c) the cell density versus the melt temperature for ce-extruded Branched P2 polypropylene materials using butane as blowing agent
Figure 5.1 1 (d) the ce11 density versus butane contents for re-extruded Branched P2 polypropylene materials using butane ris blowing agent
O J 100 120 110 160 110
Trmptnmn (C)
Figure 5.1 1 (e) the die pressure versus the melt temperature for re-extruded Branched P2 polypropylene mater& using butane as blowing agent
- . 100 120 140 160 110
Ternperiluie (C)
Figure 5.12 (ri) the expansion ratio versus the melt temperature for branched polypropyicne materials using butane as blowing agent for Die B
- ,
100 120 140 160 180
Timp.ritur~ (C)
Figure 5.12 (b) the blowing agent eficiency versus the melt temperature for branched polypropylene materials using butane as blowing agent for Die B
Figure 5.12 (c) the cell density versus the melt temperature for branched polypropyiene materials using butane as bIowing agent for Die B
Figure 5.12 (d) the ce11 density versus butane contents for branched polypropylene materials for Die B
40 T
Figure 5-12 (e) the die pressure versus the rnelt temperature for branched polypropylene materials using butane as blowing agent for Die B
Figure 5.13 (a) the expansion ratio versus the melt temperature for branched polypropylene materials using CO, - as blowing agent for Die B
Figure 5-13 (b) the blowing agent efficiency versus the melt temperature for branched polypropylene materials using CO, as blowing agent for Die B
Figure 5.13 (c) the cell density versus the melt temperature for branched polypropylene matends using CO, - as blowing agent for Die B
Figure 5.13 (d) the cell density versus CO2 contents for branched polypropylene materials for Die B
Figure 5-13 (e) the die pressure versus the melt temperature for branched polypropylene materials using CO, - as blowing agent for Die B
Chapter 6
Fundamental Mechanisms of Volume Expansion Behavior of Polypropylene Foam Filaments
6.1 Introduction
In the previous chapter, the effects of processing and materials parameters on the
foaming of polypropylene materials were thoroughly investigated, The results show that
the volume expansion ratio was govemed by the processing and materials parameters. It
was shown that the processing temperature, the blowing agent arnount and type, the
nucleating agent amount and type, the materials branching, and the die geometry are
affecting the finai volume expansion ratio of polypropylene foains. in order to improve
the expandability of polypropylene foams, we need to have a clear understanding of the
fundamental mechanisms that govern the volume expansion ratio. Despite the recent
snidies that addressed the production of low density, fine-celled polypropylene foams, no
research has been conducted to investigate the mechanisms that govern the expandability
of polypropylene foams.
In this chapter, a qualitative modeling for the volume expansion behavior was
described. initially, the fundamental mechanisms goveming the volume expansion of
polypropylene foams were identified and the volume expansion phenomena were
described based on our expetimental observation. The system setup used for monitoring
the expansion phenomena of extmded foam is then described. The procedure for
monitoring the expansion mechanisms and the image analysis will be elucidated.
Consequently, the effects of processing parameters on the volume expansion behavior
were depicted based on the extrudate images. A theoreticai mode1 for calculating the
voIume expansion ratio, ce11 size, and ceIl wail thickness was developed as a hnction of
extrudate diameter. The mode1 was used to describe the volume expansion behavior at
various distances from the die using the observed CCD images of extruded foams.
6.2 Qualitative Modeling of Volume Expansion Behavior
6.2.1 Fundamental Mechanisms Governing Volume Expansion of Polypropylene
Foams
A careful analysis of extended experirnental results obtained at various processing
conditions indicates chat the final volume expansion ratio of the extruded polypropylcne
foarns blown with butane is governed either by Ioss of blowing agent through the foam
skin or by crystailization of the porymer matrix. In general, upon exiting the die, the
foaming extrudate exhibits one of the shapes depicted in Figure 6.1, depending on the
temperature of the extrudate. At higher temperanires, rhe cross section of the extrudate
expands suddenly, Le., it has a higher initiai angle 0, and this iingie decreases as the
processing temperature decreases. Below an optimum processing temperature, the
foaming of extrudate is inhibited and the initial angie is substantially reduced. The
mechanisms governing the volume expansion of polypropylene foams are explained in
the following sections.
62.1.1 Gas los
The gas ioss phenornena which occur during f o m processing cm be correlated
with the melt temperature. The diffusivity of bIowing agents at elevated temperatures is
very high, therefore, if the processing temperature is too high, the gas cm easily escape
from the extruded foam because of its higher diffusivity at eIevated tempennires. In
addition, as the ceII expansion increases, the cell w d l thickness decreases and the
resulting rate of gas diffusion between cells increases. Consequently, the rate of gas
escape from the foam to the environment increases. Gas escape through the thin cell
wails decreases the amount of gas availabie for the growth of cells resulting in lowered
expansion. Moreover, if the cells do not freeze quickly enough, they tend to sht-hk due to
Ioss gas through foam skin causing overail f o m contraction. This mechanism is
schematicdly shown in Figure 6.2.
This phenornenon of the gas escape at hi& temperatures and the resultant foam
contraction during the ceIl growth stage at elevated temperatures was observed during
experimentation, It was observed that the extruded foam initially expanded as the
nucleated cells grew very fast and then eventually contracted. The maximum expansion
occurred very close to the die exit with the diameter of the initially expanded foam being
larger than that of the final foam extrudate. This indicates that the volume expansion ratio
of the initially expanded state was considerably high, and therefore, the shapes of
expanded cells at this state were not spherical but polyhedral with thin cell walls (Figure
6.2).
6.2.1.2 Crystallization
The crystallization behavior of semisrystalline materials is another critical factor
that affects the maximum expansion ratio in plastic foam processing. For semicrystalline
poiymers, the polymer melt solidifies at the moment of crystallization during cooling.
Therefore, in the foam processing of polypropylene, the foam structure "freezes" at the
crystallization temperature during the foaming process. If the crystallization occurs in the
primitive stage of foaming, i.e., before the dissolved blowing agent fully diffused out of
the plastic matrix and in the nucIeated cells, then the foam cannot be fully expanded.
Therefore, in order to achieve the maximum volume expansion ratio from the
polypropylene foarn, the crystallization (or solidification) should not occur before al1 the
dissolved gas diffuses out into the nucleated cetls. Upon exiting the die, the temperature
of melt decreases due to external cooling outside the die and the cooling effect due to
isentropic expansion of the blowing gases. Thus, the processing temperature at the die
determines the time after which the polymer melt solidifies. Therefore, in order to give
enough time for the gas to diffuse into the polymer matrix. the processing temperature
should be high enough. U should be noted that if the processing temperature is too close
to the crystallization temperature, the poIyrner melt would be solidified too quickly
before the foam is expanded fully, as shown in the initial section of Figure 6.2.
On the other hand, if the tempenture is too high, then the solidification time
might be too long and the gas that has difised out of the plastic melt to the nucleated
cells might escape out of the foam as discussed above. This indicates that there is an
optimum processing temperature for achieving maximum expansion as shown in the
middle section of Figure 6.2. if the melt temperature (Le., the processing tempenture) is
too high, then the maximum volume expansion ratio is governed by gas loss and the
volume expansion ratio will increase as the processing temperature decreases. if the
processing temperature is too Iow, then the volume expansion ratio is govemed by the
soIidification (Le., the crystailization) of polypropylene, and the volume expansion ratio
will increase as the temperature increases.
In addition to the effects of the processing parameters on the crystailization, the
foaming additives and materials parameters can also contribute to changes in
crystailization temperatures, The effect of these parameters on the volume expansion
ratio is shown in Figure 6.3.
6.2.2 Visualization of Expansion Behavior Using a CCD Carnera
6.2.2.1 System setup
For verifying the mode1 described above, photographie images of the extmdate
were taken and analyzed for each processing and material parameter. The system setup
consisted of the tandem extrusion system already described in Chapter 3. In addition, an
on-line progressive scan imaging system was mounted at the die exit, to capture images
of the extrudate coming out of the die. using a frame gnbber and image processing
software, The progressive scan imaging system consists of a CCD camera (CV MlO),
which has a very high shutter speed up to 1180000 second, with a magnifjing lens
(Navitar), a frarne grabber (PC vision), and image processing softwrire (Sherlock).
6.2.2.2 Experirnental Procedure
The foamed extmdate was monitored while changing the processing temperatures.
From the captured images the angle of initial swelling of foamed polymer and the
extrudate diameter as a function of the distance from the die exit were measuted, The
images of the foamed extrudate were analyzed and the effects of gas ioss and
crystailization were extracted. Some sample images are shown in Figure 6-4. Sample
results for branched polypropylene materid with 15% butane concentration ai hree
different processing temperatures are elucidated in the following sections.
6.2.2.3 Effect of Processing Temperature on the Initial Expansion Rate and Final
Diameter of the Extudate
Figure 6.5 shows the effect of processing temperatures on the initial volume
expansion behavior characterized by the initiai angle 0 of the extrudate. The figure shows
that at the higher temperatures, the initial expansion rate, and consequently 8, was quite
significant. However, since gas Ioss was accelerated at this high temperature, the final
diameter of the extnided foam was smdl. As the temperature decreased. the initial
expansion rate (8) was decreased and the final diameter of the extnided foams was
increased due to the decreased gas loss amount. There existed an optimum temperature
to achieve the maximum diameter of the extrudate. When the temperature was further
decreased, and because of the earIy crystallization, the foamed extrudate was frozen
before the extrudate was fully expanded. As a consequence, the final diameter d was
small. The results are summarized in Figure 6.6.
Figure 6.7 shows the changes occurring in the extrudate diameter, as the distance
from the die exit increased, At the optimum die temperature of 130°C, the extrudate
diameter reached its maximum value of 5 mm at a distance of 5 mm away from the die
exit, At this point the foam extrudate reached the crystallization temperature and retained
its shape due to solidification of the polymer. At the lower temperature of 120°C, the
polymer reached the crystallization tempenture earlier, about 3 mm away from the die
exit and freezed at this point. Thus, at the same gas content of 15%, the maximum
diameter achieved was about 4.6 mm instead of 5 mm which would result in sacrifice of
the expansion ratio. in the meanwhiIe, at a higher temperature of 200°C, the effect of gas
loss was more pronounced. First of dl, the gas at a higher temperature caused a sharper
initia1 expansion ratio and the rate of diameter growth was greater. Simultaneously, the
gas loss due to the increased diffusivity of the gas resuited in a lowered maximum
diameter in the extrudate, which was about 3.6 mm and was observed about 2.5 mm away
from the die exit. As the temperature of the extrudate was stiIl above the crystdlization
temperature of the polymer, it continued to experience hrther gas loss and the diameter
continued to decrease beyond the range of the photographic equipment. This decrease in
diameter could attnbuted either to contraction due to cooling or due to additional gas loss
by diffusion through the foam skin.
6.3 Theoretical Mode1 for Calculating the Expansion Ratio, Cell Size, and Cell Wall
Thickness from the Observed Foam Profile
6.3.1 Development of a Theoretical Mode1
A theoretical mode1 was developed for describing the expansion ratio in terms of
distance from the die exit at each processing temperature. The rnodel takes into
consideration gris loss at high temperatures and solidification due to the early
crystallization at low temperatures.
In order to determine the instantaneous diameter of the extrudate upon leaving the
die exit, a first order mode1 based on the conservation of mass was considered. The
continuity equation rnay be written in the following form:
riz = p Av =constam. (6.1)
where m is the rnass flow rate, p is density, v is speed, of a polymerlgas solution
expenencing flow through the channel having cross-sectional area A (see Figure 6.8). By
definingx-axis as the flow direction, Eq. 6.1 at any given point becomes
Px Axvx = mpr * (6.2)
where m,, is the combined flow rate of both polymer melt and gas. Since
Eq. 6.2 becomes
where @, is the volume expansion ratio at the distance x, ppx = p, = the polymerfgas
solution density, p, is the polymer melt density, and d, is the extrudate diameter.
Assumingv, = v p = constant at any point in the cross section at the distance x, the
continuity equation for polyrner melt is:
p p Apvp = f i p , (6.5)
where m, is the flow rate of polymer melt. and Ap is the polymer area which is a
hnction of the ce11 shape. The ce11 shape depends on the processing condition as follow:
Case 1: when x is very small, the ceIls are not fully grown and are generally of
spherical shape, separated from each other and uniformly distributed (Figure 6.8a).
Case 2: when x is at the maximum expansion point (Figure 6.8b). The cells are
hlly-grown and are polyhedral in cross-section with thin polymer walls separating them.
As a first approximation for this model, it is assurned that the cells are cubic in shape.
Case 3: when x is large, there are two possible cases. First, the polymer reaches
the crystallization tempenture before x became large, and freezes in the state that it was
in corresponding to either case 1 or case 2. Second, the temperature was above the
crystallization limit. In this case the gas will diffuse out through the thin ce11 walls and
outer skin and the cells will shrink. in this situation, the approximation of spherical cells
can again be applied and we can use the model described
1s t part of Figure 6.8a.
Modell
If the ce11 is represented by ri sphere of diameter D (Cases
in Case 1. This is shown in the
1 and 3, then
(6.6)
whiIe n is the number of cells per cross-section (see Figure 6.8). Substitution of Ap into
Eq. 6.5 yields the following:
The volume expansion ratio, 4,. is defined as
where V, and V, are the occupied volume of gas and polymer melt in the foarn.
respectively. Since the ceil density is defined to be the number of cells per unit volume of
polymer:
Combination of Eq.'s 6.9 and 6.10 results in:
Using Eq. 6. i 1, Eq. 6.8 is then rransfotmed into:
Since the flow is assumed unidirectionai. and v, = v , =v,, Eq. 6.1 1 may be substitured
into Eq. 6.4:
when $, is too Iarge, equation 6.14 becomes
If the ceIl is represented by a L-sized cube (Case 2), then
Eq, 6.6 could be rewntten as:
and therefore
For a cubical cell, Eq. 6. IO becomes:
Keeping the definition introduced in Eq. 6.9, an equation for L may be derived:
Combined with Eq. 6.1 8, this equation yields:
Then combination with Eq. 6.4 produces the following formula:
when 4,is too large, equation 6.23 becomes
Figures 6.9 (a) and 6.9 (b) show the calculated values of the volume expansion
ratio using Models 1 and 2, respectively, for a fixed ceIl density (N) of 8.7 x106 cells/cm3
while varying the number of cells per cross sectional area (n) from 200 to 800. It can be
seen that the expansion ratios predicted by both models are close to each other. However,
Model 1 would show yet better prediction in the low expansion region in which the cells
have not met each other and therefore maintain a spherical shape. On the other hand,
Model 3 would show better prediction in the high expansion region where the ce11 shape
can not be spherical but rather polyhedral due to the contact of the cells. The predicted
average ceIl sizes based on both rnodek are also close to each other within 25% (see
Figure 6.10). Since Model 1 predicts a negative cell wall thickness for the large
expansion region (see Figure 6.1 1), Model 2 seems to be a better model overall to
describe the expansion ratio, cell size, and cell wall thickness. Therefore, Model 2 has
been selected to estimate the expansion ratio, ce11 size, and cell wall thickness of the
extmded polypropylene foams from the observed extrudate profile as shown in the next
section. However, it should be mentioned that this model is based on the observed outer
shripe of the extrudate, the finally observed number of cells in the cross-section, and the
ce11 density of obtained foam. in order to predict the shape of the extrudate as a function
of time, a set of equations that describe the dynamics of the polymer-gas system
including the flow of polymer melt, diffusion of gas, and heat transfer needs to be solved
simultaneously- Further research is required for this work.
6.3.2 Determination of Expansion Ratio, Ce11 Size and Cell Wall Thickness of
Extruded Foams from the Observed Profiles
Figures 6.12, 6.13, and 6.14 show the calculated volume expansion ratio, ceIl size
and celt wall thickness of the extmded foarn, respectively, as a function of the distance
from the die exit at the melt temperature of 120°C. 130°C, and 200°C. Table 6.1 shows
the number of cells per cross section area and the ce11 density for each temperature.
Figure 6.12 shows that for the lower temperatures of 120°C and 130°C, the expansion
ratio initiaily increased rather rapidly with increasing extrudate diameter. This trend
continued up to 60-fold expansion. At the point of maximum diameter, rit a distance of 5
mm away from the die exit, the maximum volume expansion predicted was 90-fold, for
the optimum processing temperature of 130°C. On the other hand, at the processing
temperature of 120°C, the polymer reached the crystallization ternperature before
attaining the maximum expansion, and consequently the achieved maximum expansion
ratio was only about 78-fold. At the higher temperature of 30O0C, the plot shows that the
volume expansion ratio was very low. The excessive gas loss, due to increased gas
diffusion at the higher temperature resulted in only a very low maximum volume
expansion ratio of 13-fold. On the other hand, it is worth of noting that the ceil sizes of
120°C and L3O0C were very close to each other (Figure 6-13), whereas the cell wall
thickness were quite different (Figure 6.14). This demonstrates chat the lower temperature
prohibited the ceII walls from being seretched and thinned during expansion becnuse of
the eariy crystailization,
6.4 Summary and Conclusions
The fundamental mechanisms governing the volume expansion behavior of
polypropylene foams were determined based on the experimentai results shown in
Chapter 5. Tt turned out that either gas loss or polymer crystallization governs the
expansion behavior of polypropylene foams. A progressive image scanning setup was
configured to capture the images of the foamed extrudace coming out of the die on a PC.
The images captured were anaIyzed and the data were used to verify the proposed
mechanisms, A theoreticai mode1 h a been proposed, which relates the instantaneous
expansion ratio, ce11 size, and cell wai1 thickness to the exwdate diameter as the distance
from the die exit varies. The developed mode1 was effectively used to describe the
variation of the cell morphology in the extruded foarn from captured images. However,
the mode1 used the experimentaily observed parameters such as the ceil density, the
number of cells per extrudate cross-sectiond area, and the final volume expansion ratio.
Further studies are required to describe the gas Ioss phenomena over tirne.
Tabie 6.1 Observed number of cells per cross section at various temperatures
I (Cl 1 cross secrion (n) 1 IN) I Die Temperature Number of celIs per Cd1 Density
Crystallizalion ,,, Gus Loss 4-- >
Figure 6.2 Effect of gas Loss and crystallization on the volume expansion
Crystallization by gas loss I
Temperature
Figure 6.3 Fundamental volume expansion mechanism of polypropylene Foarns
Figure 6.4 images of the foarn extrudate coming out of the die
Figure 6.5 Effect of processing temperature on the initial expansion
Ternparatum (C)
Figure 6.6 Effect of Processing temperature on the initial diameter
0.00 2.00 4.00 6.00 8.00 10.00
Oistanci tom Oia (mm)
Figure 6.7 Extrudate diameter as a function of the distance from the die
die at a
Figure 6.8 Description of ce11 shape models
O ! t
O 1 2 3 4 5 6
Eztnrdate Dkmibr (mm)
6.9 (a) The calculated volume expansion ratio based on Model 1 for ceII density = 8 . 7 ~ IO6 ceIls/cm
0 ! O 1 2 3 4 5 6
Extrudita Dhmimr (mm)
6.9 (b) The caiculated volume expansion ratio based on Model 2 for ce11 density = 8 . 7 ~ 106 cells/cm
O 5 1 O 15 20 25 30 35 40
Expanrlan Rafio
Figure 6.10 Calculated average ceIl size for for cell density = 8.7~ IO6 cells/cm
Figure 6.1 1 Calculated ce11 wail thickness for ce11 density = 8.7~ IO6 ceIls/cm
1°- 0.00 ff . 0.00 200 4.00 6.00 8.00 10.00
Distance from Die (mm)
Figure 6.12 Calculated expansion ratios of extnided foams from the observed profiles
0.00 2.00 4.00 6.00 8.00 10.00
Distance (rom Die (mm)
Figure 6.13 CdcuIated average ce11 size of extmded foams from the observed profiles
0.00 200 4.00 6.00 8.00 10.00
Distanci fmm O k (mm)
Figure 6.14 Calculated ce11 wall thickness ofextnrded foams from the observed profiles
Chapter 7
Summary and Conclusions
7.1 Sumrnary
A continuous extrusion foaming process has been studied for the manufacture of large
volume expansion ratio with a fine-celled structure in extruded polypropylene foams. A
tandem-extrusion foaming system was designed and anaiyzed based on the axiomatic design
approach. Experiments were performed to check the functionaiity of the designed system.
The strategies for promoting ultra low-density polypropylene foarns were developed.
The effects of processing and materials parameters, including those of the processing
temperature, arnount of blowing agent, amount of nucleating agent, blending of blowing
agents, long-chain branching of polymer, blending of linear and branched materids, and die
dimensions on the finai foam properties were investigated in detail.
The fundamental mechanisms for the promotion of high volume expansion ratio of
polypropylene foams were elucidated. A qualitative modeling for the volume expansion
behavior was described at various processing temperatures A theoretical model for
calculating the volume expansion ratio, average ce11 size and ceIl wall thickness was
developed. This model was used to describe the volume expansion behavior at various
distances from the die using the observed CCD images of extnided foams.
Basic studies were also cam'ed out co investigate the effects of dissolved gas on the
PVT data, crystallization kinetics and meIt fracture behaviors of polypropylenelgas soIutions.
The obtained results were usehl in the understanding of the processing technology that
govems the production of low-density, fine-celled polypropylene foams.
7.2 Conclusions
The experimental work presented in this thesis leads to the following conclusions:
1. A tandem extrusion system for the production of low-density, fine-celled polypropylene
foams was designed and constmcted. This system will considerably enhance the
potentid to scale up the existing system into an industrial production system.
3. The basic strategies employed for the promotion of a large volume expansion ratio with
polypropylene materials were fourfold: to use a branched material for preventing ce11
coalescence; to use a long-chah blowing agent with low diffusivity; to lower the meIt
temperature for decreasing gas loss during expansion; and to optimize the processing
conditions in the die for avoiding too quick crystdlization.
3. Low-density, fine-celled polypropylene foams were successfully produced using the
designed system. The branched polypropylene foams have a maximum volume
expansion ratio in the range of 90 times, and a ce11 density of higher than 1o6 ce11s/cm3.
The results show the effectiveness of the fundamental strategies adopted to promote
large expansion.
4. The effect. of dissolved butane on the PVT relationships of linear and branched
propylene materials in a molten state were investigated. The specific volume of the
propylenehutane solution increased significantly with an increase in the percentage of
butane injected in both branched and linear propylene. At al1 experimental conditions
selected, the specific volume was found to be higher for the linear propylene than for the
branched propylene. When butane was dissolved in the propylene matrix, the sensitivity
of the specific volume with respect to pressure increased with the butane content for
both the linear and branched resins, whereas there was no significant changes in the
sensitivity with respect to the temperature.
5. A series of experiments were conducted to investigate the effects of materid branching,
foa ing additives, cooling rate, hydraulic pressure, and dissolved gas on the
crystallization behaviocs of potypropylene resins. Branching and foaming additives in
the polypropylene matrix caused a significant increase in the crystailization temperature.
The crystallization temperature was a sensitive function of the cooling rate and it
decreased as the cooling rate increased. Crystdlization of polypropylene materials was
enhanced as the hydraulic pressure increased. But the dissolved N2 and CO1 lowered the
crystailization temperatures of polypropylene resins.
Experirnental studies were carried out to investigate the effects of branching, processing
temperature, foaming additives, and blowing agent on the critical shear stresses of linear
and branched polypropylene resins. An on-line technique for detecting the onset of
surface melt fracture for extruded foam has been developed by visualization of the
extrudate using a CCD carnera. The long-chain branching of polypropylene matends
significantly decreased the critical shear stress of the resins. The critical shear stress was
insensitive to the die ternperature; however, the die temperature signitlcantly riffected
the critical shear rate at the onset of surface melt fracture. The foaming additives of talc
and GMS increased the critical shear stress of branched polypropylene rnaterials more
than that in linear rnaterials. The dissolved butane significantly increased the critical
shear stress'es of linear and branched polypropylene resins.
The effects of processing parameters such as the temperature, the materials parameters
such as the blowing, nucleating agents, and long chain branching, and the die geometry
on the foam density and ceIl density of polypropylene foarns were investigated.
Despite its high flamrnability, butane is a very effective blowing agent for low-density
polypropylene foams cornpared to COz. A higher ceIl density and lower volume
expansion was obtained by increasing the amount of CO2 when using a blend of CO2
ruid butane as blowing agent.
There exists an optimum temperature for achieving the maximum expansion ratio of
polypropylene foam with butane. This optimum ternperature decreased as the amount of
blowing agent increased because of the plasticizing effects.
IO. The cell dènsity increased proportionally when the amount of taIc increased frorn O to
2.4 5%- However, there was an optimum talc amount for achieving the maximum volume
expansion ratio.
11. The expansion ratio obtained from the linear polypropylene materials was rnuch lower
than that from the bnnched polypropylene materiais because of severe ce11 coalescence,
The expetimental results indicate that branched polypropylene rnatcrials are effective for
lowdensity foam application because of the reduced degree of ceII coalescence.
12. The use of a blend of linear and branched polypropylene materials revealed that the ce11
density in the foamed material increased as the amount of branched polypropylene was
increased. Despite that there was no clear goveming mechanisms observed for the ce11
density of linear and branched polypropylene blends, the ceil coalescence could
obviously be reduced by the addition of branched HMS material, giving the blend higher
melt strength and melt extensibility. On the other hand, the volume expansion increased
when the amount of branched rnaterial in the blend increased in the same way as the
previous phenomena. The optimum Lempenture for producing the maximum volume
expansion ratio decreased as the branched material content was increased. This is
believed to be due to the high viscosity of the linear propylene together with the
extensional characteristics of the added HMS component.
13. The use of re-extruded branched materials resulted in a lower ce11 density and a lower
volume expansion ratio- This was attributed to the breaking of the long-chain branching
of the high melt strength polypropylene, which led to ceIl colilescence in the foam
structure.
14. The maximum volume expansion ratio was decreased using a lower pressure drop rate
die. However, the expansion ratio was higher with a lower pressure drop rate die in the
high temperature range due to the reduced gas loss from the foam extrudate. The ceIl
density was higher in the case of higher pressure drop rate as per previous studies
[17,18].
15. A careful anaiysis of extended experimental results obtained at various processing
conditions indicates that the final volume expansion ntio of the extruded polypropylene
foams blown with butane is governed either by loss of blowing agent or by
crystallization of the polymer matrix. The CCD images were anaiyzed to illustrate both
these mechanisms of gas [oss and crystallization during foarning at various temperatures,
and it was observed that the maximum expansion ratio was achieved when the governing
mechanism was changed from one to the other. Therefore, it is highly recomrnended to
look for this transition point to maximize the expansion ntio for the low-density foam
applications.
Recomrnendations and Future Work
The following suggestions are made for the direction of future research on the
production of low-density, fine-celled polypropylene foam:
1. The extrusion process for the manufacture of low-density, fine-celled polypropylene
foarns filaments developed in this thesis cm provide useful insights to extend the
system to sheet extrusion and injection molding processes. Using the processing
conditions determined here, appropriate design strategies can be developed for other
foam processes,
3. This study was confined to branched and Iinear polypropylene homopolymer. It was
found that each material has particular processing conditions. However, other
thermoplastics materials with different properties could be examined using the
process developed to determine suitable processing conditions for various materials.
3. The die design needs to be modified to improve the surface melt fracture of the
exuudate. A number of parameters, including the Iength to diameter ratio, the die
exit angle, and the die material cm be altered to investigate how they influence the
surface melt fracture.
4. The use of a CCD canera in this study was confined to the image rinalysis of the
filament extrudate to describe the fundamental mechanisms for the volume expansion
ratio of polypropylene foams. However, the study of the ce11 growth mechanisms
based on computational mathematical modeling such as development of a finite
element anaiysis simulation software will be beneficial to determine the growth of
bubbles in any processing conditions.
5. Incorpcirating a completely automated computer control for the tandem extrusion
system for the production of low-density fine-celled plastics foams wiI1 be of a great
value in reducing the error bounds for the measured fundamental pmperties of
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