polylmer blend de-mixing and morphology ......polymer blend de-mixing and morphology development...
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POLYlMER BLEND DE-MIXING AND MORPHOLOGY DEVELOPMENT DURING TUBE FLOW
Askar Karami
A thesis submitted in confonnity with the requirements for the degree of Doctor of Philosophy
Graduate Department of Chemical Engineering and Applied Chemistry University of Toronto
O Copyright by Askar Karami, 1999
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POLYMER BLEND DE-MIXING AND MORPHOLOGY DEVELOPMENT
DURING TUBE FLOW
PhD. Thesis, 1999
Askar Karami
Department of Chernical Engineering and Applied Chemistry
University of Toronto
Abstract
This work is an investigation of morphology and de-mixing of polymer blends during
melt flow through a tube. Morphology is the relative size, shape and location of each
distinguishable phase present in a polymer blend. De-mixing is the shear-induced
migration of different types of polymers away from each other during the flow. Being
able to tailor de-mixing during extrusion can potentially result in a new family of plastics
waste recycling processes with mixed waste entering an extruder and separate streams of
different polymer types leaving it. Also, control of morphology development can lead to
the formation of layered structures without the need for two or more extruders and co-
extrusion. These ideas formed the basis for a U.S. Patent. However, obtaining an
understanding of the phenornena is critical to improving separations to a practical level.
This thesis is directed at elucidating morphology development and de-mixing of polymer
blends in the most simple process design: melt flow through a tube. The work had four
objectives. The first was to design a process that would enable elucidation of both
morphology and de-mixing along the tube. This was done by attaching a long segmented
tube to a static mixer which in tum was attached to the end of a single screw extruder. At
the conclusion of a run, the tube was quenched and disassembled to provide the needed
samples. The second objective was to develop analytical methods to measure polymer
composition in the samples. A mid-infrared spectrometer technique and a method based
on the use of a differential scanning calonmeter were developed. The third objective
was to use the above accomplishments to elucidate morphology developrnent and
polymer migration. Shear-induced migration was quantitatively shown in various
polyethylene-polypropylene, polypropylene-nylon6 and polyethylene-nylon6 blends.
The theoretical rate of viscous energy dissipation per unit length of the tube was used to
show that the observed shear-induced migration was in accordance with the principle of
energy minimization. The ratio of the viscosity of the dispersed phased to that of the
continuous phase greatly influenced the morphology of polypropylene-nylon6 and
polyethylene-nylon6 blends: a droplet-dispersed phase structure occurred at a high
viscosity ratio whereas a multi-layer structure resuited at viscosity ratios near unity.
Shear-induced deformation and coalescence contributed to formation of the multi-layer
structure. Finally, the fourth objective was to investigate the effect of morphology
development on viscosity measurement by capillary rheometry. The extruder-tube
process was used as the rheometer. Morphology had a large impact on the value of the
measured viscosity and viscosity-composition data were shown to be not readily fit by
two mixing rule models: Lees' model and a sheathtore model. Greatly improved results
were obtained by introducing a "shear-induced interlayer slip factor" into the sheath-core
model.
ACKNOWLEDGEMENTS
I would like to thank my supervisor, Professor S. T. Balke for al1 of his advice, guidance,
and encouragement throughout this work. I would also like to thank Professors Cluett,
Kortschot and Park of my reading cornmittee for providing valuable comments and usehl
guidance.
1 wish to thank the Manufacturing Research Corporation of Ontario and the Ontario
Centre for Matenal s Researc h (now Materials and Manufacturing Ontario) for t heir
support of this work.
Al1 people in the PPMAG group have been helpful, especially Carrie, Chris, Dennis,
Jimmy, Joseph, Keivan, Lianne, Mark, Ognian, Paul, Ramin, Runi, Sak, Sina and Dr.
Saed.
My family and my wife have supported me unfailingly and 1 thank them for the love,
understanding and encouragement that they have always provided. Finally, without the
love and guidance From my deceased parents, 1 would not have the opportunity to obtain
a higher education.
TABLE OF CONTENTS
ABSTRACT ACKNOWLEDGEMENTS TABLE OF CONTENTS LIST OF FIGURES LIST OF TABLES NOMENCLATURE
INTRODUCTION
THEORY
2.1 Phase morphology development in immiscible polymer blends
2.2 De-mixing in immiscible polymer blends
2.3 Rheological property of polymers
2.3.1 Viscosity 2.3.2 Capillary flow 2.3.3 Flow behavior of immiscible polymer blends
2.4 Analytical methods
2.4.1 Fourier Transform Infiared Spectroscopy (FTIR) 2.4.2 Differential Scanning Calorimetry @SC)
3.1 Materials 3.2 Processing equipment 3.3 Anal ytical methods 3.4 Morphology 3.5 Melt viscosity
RESULTS A N D DISCUSSION
4.1 Process development
4.1.1 Design requirements 4.1.2 Shear-induced migration and phase morphology
development in extruder 4.1.3 Tube design
ii iv v vii xi xii
1
3
3
2 2
2 1
21 21 24
30
30 32
36
36 37 41 45 46
47
47
47 48
55
4.1 .4 Extruder-Rheometer 4.1.5 Summary
Analytical methods development
4.2.1 Fourier Transform InFrared Spectroscopy (FTIR) 4.2.2 Differential Scanning Calonmetry @SC) 4.2.3 Summary
Shear-induced migration and phase morphology
4.3.1 S hem-induced migration 4.3.2 Phase morphology development 4.3.3 Summary
Flow behavior of immiscible polymer blends
4.4.1 Systeml 4.4.2 System3 4.4.3 Evaluation of experimental data 4.4.4 Summary
5.0 CONCLUSIONS
6.0 RECOMMENDATIONS
7.0 REFERENCES
8.0 APPENDICES
LIST OF FIGURES
Figure 2.1: Possible morphological structures for an immiscible blend A) core-sheath
B) multi-layers C) dispersed phase D) mix of dispersed pliasc and multi-layers.
Pi y re 2.2: Schernatic representation of encapsulation in simultaneous side-by-side flow
in a coextrusion; the less viscous component encapsulates the more viscous component.
Figure 2.3: A schematic of sinusoidally disturbed thread.
Figure 2.4: Coalesccnce of two Newtonian liquid droplets.
Figure 2.5: Showing the Poiseuille flow in the tube for a Newtonian fiuid.
Figure 2.6: Showing a particle in tubular shear flow.
Figure 2.7: A schcmatic of sheath-core morphology for cornparison: Case 1) high viscosity
component is located at low shear rate region (core), Case II) high viscosi ty cornponent is
located at fiigh shear ratc region (shell) CaseIII) a monocomponent alone is occupied
the wholc tube.
Figure 2.8: Ei/Em versus h' as a function of volume fraction of component 1 (fiom 9,=0.2
to +,=0.7 wvith 0.1% interval) for a=l , log-log scalc.
Figure 2.9: EDEin versus A' as a function of volume haction of componcnt 1 (from 9,=0.2
to +,=0.7 with O. 1% interval) for a= 1, log-log scale.
Figure 2.10: versus A* as a hinction of volumc fraction of high Mscosity component
(from $, =0,2 to $i =O.î with O. 1% interval), log-log scale.
Figure 2.11: Showing orientation in thc flow direction for a b imy blend.
Figure 2.12: Sheath-corc configuration in the flow dircction for a binary blend in a tube.
Figure 3.1: A schemtic of Process 1.
Figure 3.2: A schematic of Prucess 2.
Figure 3.3: A schernatic of Process 3.
Figure 3.4: A schematic of the tube along with the positions for taken the samples.
Figure 4.1: Morphology development of 30170 PA4/HDPE2 perpendicular CO the flow direction 49
at 1 cm downsîream from the scrcw tip; Temp.=230 OC, <,=38 sec-', A= 10.6 A: shell, B: core.
Figure 4.2: Morphology development of 30170 PA-6/HDPE2 dong the transition piece 52
parallel to the flow direction; Tcmp.=230 OC, Y,= 38 sec", A= 10.6, A: shell B: core,
a) Ls3.5 cm b) L=4.8 cm.
Figure 4.3: Morphology development of 30170 PA6/HDPE blends after siatic mixer paralle1 54
to the flow direction; yoa,= 3 1 sec", Temp.=230°C, A: shell B: core (a) 30170 PAdMDPEl
h= 0.64 (b) 30170 PA4HDPE2 h = 10.21.
Figure 4.4: A schematic of a segmented tube. 55
Figure 4.5: A schernatic of the segmented tube for mounting the pressure transducer. 57
vii
Figure 4.6: Pressure readings versus time for i-IDPE1 and PP3 at two temperatures
viii
a) low screw speed (5 r.p.m) b) high screw speed (33 r.p.m.).
Fipre 4.7: End correction determination for HDPE 1 a) dinerent flow rate ( Q = ~ o - ~ m3/sec)
b) upper and lower 95% conlidence intervals for Q = 0 . 1 5 ~ 1 0 ~ (m3/sec).
Figurc 4.8: Comparison of capillary rheometer and extnider-rheometer viscosity rneasufements
for PP3 at Temp.=230°C.
Figure 4.9: ETIR spectra of HDPEFP blends: number of s a n s 60, resolution 4cm-l . Figurc 4.10: Net absorbance at charactcristic peaks vcrsus composition a) characteristic peaks
of HDPE (719.32cm", 730.89cm") b) characteristic p a k s of PP (971.3cm-', 1166.7crn'~).
Figure 4.1 1: Calibration curvcs for HDPEPP blends using SLR a) characteristic pcak of
HDPE 7 19.32cm-' b) characteristic peak of PP 1 166.7cme'.
Figure 4.12: Rcsidual plots of calibration samples and prediction samples from SLR modcls
a) characteristic peak of PP (1 166.7 cm'') b) characteristic pak of HDPE (719.32 cm").
Fi y re 4.13: Plot of the fit with 95% confidencc intervals for HDPEPP blcnds using FTïR
a) nct absorbance at characteristic peak of PP (1 166.7 cm-') b) net absorbancc at charactcristic
peak of HDPE (7 19.32 cm").
Figure 4.14: DSC spcctra of pure polymers (HDPE, PP and PA-6): heating ratc=1O0C/min.
Figure 4.15: Effcct of thcrmal history on thc DSC spectra of 6 O h O HDPE/PP:
heating ntc= IO°C/min.
Figure 4.16: Effect of the grade of HDPE on thc DSC s p c c t m : hcating ratc= 10°C/min.
Figure 4.17: Effect of thc grade of PP on the DSC spectrum: heating nte=lO0Umin.
Figure 4.18: DSC spectm of HDPE/PP blends: heating rate=1O0Umin.
Figurc 4.19: DSC spectra of HDPEPA-6 blends: healing rate= 10°C/min.
Figure 1.20: DSC spectra of PPiPA-6 blends: heating rate=lOOUmin.
Figure 4.21: Area under melting peak as a function of composition for polymer blends
a) HDPEPP blends b) HDPEPA-6 blends c) PP/PA-6 blcnds.
Fi y re 4.22: Plots of residuai for HDPEFP blcnds using DSC: a) SLR model (ana under
melting peak of PP b) SLR model (area under melting peak of HDPE) c) MLR mode1
(arcas under melting peak of PP and HDPE).
Figure 4.23: Residual plot of HDPEiPA-6 blends from SLR model.
Figure 4.24: Rcsidual plot of PP/PA-6 blends from SLR model.
Figure 4.25: Plot of the fit with 95% confidence intervals for HDPEPP blends using DSC
a) area undcr the melting peak of PP b) a m under the melting peak of HDPE.
Figure 4.26: Plot of the fit with 95% confidence intervals for HDPElPA-6 blends using DSC.
Figurc 4.27: Plot of the fit with 95% confidence intervals for PP/PA-6 blends using DSC.
Figure 4.28: Shear viscosity as a fiinction of shear rate of pure HDPE and pure PP for
Systeml at different temperatures: (a) 195OC (b) 2 10°C (c) 230°C.
Figure 4.29: Shear viscosity as a fiinction of sliear rate of pure PP and pure PA-6 for
Systcm2 at different tcmperatures: (a) 230°C (b) 260°C.
Figure 4.30: Shear viscosity as a function of s h w rate of pure HDPE and pure PA4 for
Systcm3 at differcnt tcmperatures: (a) 230°C (b) 260°C.
Figure 4.31: Radiai composition profile of PP in 30170 PP2lHDPE 1 blcnd across
the cross-section of the sample; T= 1 9S°C, y'.,=22 sec-', L/D= 13 7.
Figure 4.32: Shear rate distribution in the tube as a hc t ion of radius for
30/70 PP2/HDPE 1 : T= 195°C.
Figure 4.33: Showing, thc effcct of the lengtii of the tube on de-rnixing for
25/75 PPZ/HDPE 1 blcnds; T=2 10°C, f,=32 sec".
Figure 4.34: PA4/PPI 30170 blend morphology dcvelopment at 0.5 cm aRcr static mixer
pcrpcndicular to Lhc flow direction, f,,=30 sec", Temp.=230°C A: shcll B: corc.
Figure 4.35: Effcct of viscosity ratio on thc morphology of 30170 PA4IPP blcnds;
romp= 39sec-' , Tcmp.=230°C. L/D=94, A: shell B: core (a): PA6/PP2 h=5.89
(b): PAdPP3 h=2.73 (c): PA4/PP1 h= 1.4 1.
Figure 4.36: PA4PP 1 30170 blcnd morphology dcvclopmcnt along thc length of thc tube
parailel to the flow direction; f ~ 3 0 scc", Temp.=230°C , h=1.21 A: shcll El: corc
a) L/D=44 b) L/D=69 c) L/D= 13 2.
Figure 4.37: PA-6PP2 30170 blcnd morphology dcvclopmcnt dong the lcngth of the tube
parallel to the flow direction, Y,,= 9 scc", Temp.=23O0C, h4.04 A: shell B: corc
(a) L/D=56 (b) !,ID= 106 (c) LD= L56.
Figure 4.38: OM showing thc effect of tubc Iength for 30 wt% P A 4 in 70 wvC4 HDPE1,
paralle1 to the flow direction, yO,= 3 1 sec", Tcmp.=230°C, h= 0.64 A: shell B: core
(a) LJD-31 @) L/D=56 (c) L/D=144.
Figure 4.39: Effect of viscosity ratio on the morphology of 30170 PA4/HDPE blends
paralle1 to flow direction, <,=36sec-'. Temp.=230°C, L/D=125, A: shell B: core
(a): PA4/HDPE2 h=10.5 (b): PA61HDPEl h=0.8.
Figure 4.40: Effect of composition on the morphology development of PA-GMDPE 1 blends;
flow direction, y",= 30 sa", Temp.=230°C, LlD=l5O, A: sliell B: core a) LOI90 b) 15/85 c)20180.
Figure 4.41: Effect of temperature on morphology of 30170 PA4/HDPEI blends; 127
L/D=1 IO, screw speed=60 r.p.m, A: shell B: core (a) 230°C (b): 260°C.
Figure 4.42: Melt Mscosity versus composition as a function of wall shear stress for 131
HDPE lPP3 blends a) 195°C b) 260°C.
Figure 4.43: SEM photograph of 20180 PP3fHDPE 1 paralle1 to the flow direction, 133
wall shear stress=0.08 MPa, A: shell B: con a) 19S°C b) 260°C.
Figure 4.44: Melt Mscosity versus composition for HDPEiPA-6 blends, T=230°C. 134
Figure 4.45: OM photograph of HDPE2DA-6 blends in the flow direction, 137
A: shcll B:core, sliear stress=0.01 MPa, Temp.=23O0C a) 80120 b) 60140 c) 50/50.
F i y r e 4.46: OM photopph of HDPE lPA-6 blends in the flow direction, 141
A: shell Bxore, shear strcss=0.06 MPa, Temp.=230°C a)80/20 b)70/30 c)60140 d)50/50.
Figure 4.47: A comparison betwcen thc expcrirncntal data and the modeIs for HDPE l/PP3 b1ends 144
at two tcmperatures; 195°C and 260°C a) 0.02 MPa b) 0.05 MPa c) 0.08 MPa.
Figure 4.48: A comparison betwecn experimentai data and the models for HDPEPA4 blends,
T=230°C a) HDPE2PA-6 blends b) HDPElPA-6 blends.
Fi y re 4.49: Cornparison of sheath-core model and modifred sheath-core mode1 with
cxperirnental data for HDPE 1m-6 blcnds, T=230°C.
Figure 1.1: A schematic of conccntric laycrs in a tubular flow.
Figure 1.2: A schcmatic of shcath-corc morphology for comparison: Case 1) high viscosity
cornponcnt is located at low shcar n tc region (core), Case II) high viscosity component is
locatcd at Iiigh shcar rate rcgion (shell) CaseIII) a monocomponent donc is occupicd
the whoie tube.
Fi y re m. 1: End correction detemination for PP 1 a) dürercnt flow mtc (Q== lo4 m3/scc)
b) upper and lowcr 95% confidence intcrvals for Q=0.16x10~ (m3/scc).
Fi y rc m.2: End correction determination for PP3 a) different flow rate (Q= 1od m3/sec)
b) uppcr and lower 95% confidence intervals for Q=O.l5x 10" (rn3/sec).
Figure IV.1: Arca undcr melting pcaks as a function of composition for HDPElPPS blcnds.
Fi y re IV.2: Arca undcr mclting pcak versus composi tion a) HDPE2PA-6 blcnds
b) PP2PA-6 blcnds.
Figure N.3: Plots of residual for HDPElPP2 blends using DSC a) SLR mode1
( m a undcr mclting pcak of PP) b) SLR modcl (area under mclting pcak of HDPE)
c) hdLR model ( m a undcr mclting peak of PP and HDPE).
Figure IV.& Rcsidual plot from SLR Mode1 a)HûPE2/PA-G b1ends b)PP2/PA-6 blends.
LIST OF TABLES
Table 3.1: Blend systems.
Table 3.2: Pertinent polymer properties.
Table 3.3: Opcrating conditions of Proccss 1 and Process 2.
Table 3.4: Opcrating conditions for studying viscosity-composition behavior of HDPE 1/PP3
blends using Proccss 3.
Table 3.5: Opcrating conditions for studying viscosity-composition behavior of HDPEfiA-6
blends using Proccss 3.
Table 3.6: Operating condition: hi studying sliear-induced migration and phase morphology
devclopmcnt using Process 3.
Table 4.1: Composition analysis of 30170 PA-G/HDPE blends after cxtrudcr using DSC,
T=230 O C .
Table 4.2: n, k valucs from extruder-rheorncter, ~ernp.=220'~.
Table 4.3: List of thc characteristic peaks sclccted for HDPE and PP.
Table 4.4: Corrclation coefficients for the characteristic peaks of PP.
Table 4.5: Correlation cocfficients for the charactcristic pcaks of HDPE.
Table 4.6: The measurement of goodness of fit of calibration curvcs for HDPEPP blends
using FTiR mcthod.
Table 4.7: DSC expcnmcntal m s .
Table 4.8: The mcasurement of goodness of fit of calibration cuves using DSC method.
Table 4.9: Prediction error of calibration curvcs using DSC mcthod.
Table 4.10: Composition analysis of 30/70 PP/HDPE blends using FTIR at T=195 OC
and L/D= 1 3 7.
Tabk 4.11: Composition anaiysis of 30170 PA-61HDPE2 blends using DSC at T=230 OC
and L/D= 144.
Table 4.12: Composition analysis of 3O/ïO PA-6/PP2 blcnds using DSC at T=230 O C
and L/D= 13 1.
Table 4.13: Composition analysis of 30170 PA-6IHDPEl blends using DSC at T=230 O C
and L/D= 144.
Table 4.14: Composition analysis of 30170 PA-6fPPl blends using DSC at T=230 OC
and L/D= 134.
Table 4.15: EI/EII for the blend systems used in ihis work, +,=0.7.
Table W.1: Calibration results for immisciblc polymer blends using DSC.
A A A A0 a a a (VI Abs. A m ANOVA B b b b C
C.I. Ca cac CHDPE CLS Corn. CPP D D D d DSC E E Er EII E~ri EVA EVOH e e FTTR GPC H HDPE IL S IR k k
absorbance area under the melting curve in a DSC spectmm distortion amplitude initial distortion amplitude droplet radius the ratio nAlnB absorptivity at v absorbance area under the melting peak per unit mass analysis of variances minor principal axis sample thickness radial position of interface in a sheath-core configuration the ratio k & ~ concentration confidence interval s Capillary number critical Capillary number concentration of HDPE classical least square composition concentration of PP dispersed phase average particle size tube diameter (capillary diameter) diameter droplet diameter differential scanning calorimetry activation energy total rate of viscous energy dissipation per unit length of tube total rate of viscous energy dissipation per unit length of case 1 total rate of viscous energy dissipation per unit length of case II total rate of viscous energy dissipation per unit length of case III ethylene vinyl acetate copolymer ethylene vinyl alcohol copolymer magnitude of the negative intercept residual Fourier transform infrared gel permeation chromatography height high density polyethylene inverse least square infiared power law' constant interfacial tension
xiii
m MF1 MLR
ni OM P PA-6 PCR PID PLS PP PS PVC Q QI 9 q 9 R R Ro r r R~ RMSEC RMSEP S.D. SEC SEM SEP Si SLR T t fb TA Temp. T,
power law' constant of component i ce11 constant tube length (capillary length) major principal axis sample weight number of samples in prediction set melt flow index multiple linear regression power law' constant number of samples number of samples in calibration set power law' constant of component i optical microscopy pressure polyamide-6 (nylon6) principal component regression proportional, integral, derivative partial least square polypropylene polystyrene polyvinyl chloride volumetric flow rate total volume flow rate heating rate number of observations growth rate parameter tube radius (capillary radius) gas constant initial radius of the thread correlation coefficient radial position from the center of the tube (capillary) multiple correiation coefficient squared root mean square error of calibration root mean square error of prediction standard deviation standard error of calibration scanning electron microscopy standard error of prediction the ratio of the radius of the i-lth layer to the ith layer simple linear regression temperature time breakup time thermal analysis temperature g lass transit ion temperature
xiv
glass transition temperature of component i melting point velocity volume volume of component i total volume volume velocity in 2-direction velocity in z-direction of component 1 velocity in z-direction of component 2 mass fiaction mass fraction of component i weight independent variable average of x's dependent variable mean experimental value predicted vale from the f i t experi mental value viscosity ratio (q 1 1 ~ 3 )
interlayer-slip factor characteristic slip factor viscosity viscosity of the blend viscosity of di spersed phase viscosity of component i viscosity of matrix distortion wavelength viscosity ratio (qd/qm)
viscosity ratio (q ,/Q) shear stress in the 2-direction wall shear stress wavenumber degree of fieedom volume fraction volume fraction of component i viscosity viscosity of component i shear rate apparent wall shear rate shear rate of the blend shear rate of component i shear rate at the wall of the tube (capillary) Gibbs fiee energy Gibbs fiee energy of mixing
mf heat of fùsion mm, enthalpy of mixing Pi viscosity of component i AI' pressure drop across the capillary length AP pressure drop d o n g the length of the tube A L entropy of mixing
1.0 Introduction
The product resulting fiom the melt processing of polymer blends can be very different
depending upon both the materials involved and the processing conditions. Morphology
development and de-mixing are the two primary contributors to this difference and are
the subject of this thesis. Morphology is the relative size, shape and location of each
distinguishable phase present. Normally the product of a two wmponent blend will
exhibit two clearly different solid phases. For exarnple, one phase may appear as
spherical particles distributed in a continuous matrix of the other. However, as we shall
see in this work, often each phase will consist of both of the blend's polymer
components. That is not to Say that the product will be homogeneous in composition.
Different types of polymers can separate away fiorn each other during melt flow of the
blend. This separation is termed "de-rnixing".
Polymer morphology and homogeneity of composition of the product both can strongly
influence the resulting mechanical properties. Most often manufacturers require a
uniformly mixed polymer blend which is uniform in composition. Finely dispersed
droplets of one phase in another and the absence of de-mixing are the objective.
However, there are exceptions. For example, multi-layer films are a polymer blend
morphology that is oflen highly desired for improved barrier properties in food
packaging. Co-extrusion is used to obtain such morphologies. The migration (i.e. de-
mixing) of lubricants, slip agents and mold release agents to the surface of the processed
polymer is also highly desirable.
The increasing need to recycle polymer blends and the possibility of more easily forming
usefùl blend morphologies are the motivations for this work. Our recent patent [1]
proposes that control of polymer morphology and de-mixing can allow sorting of
different types of plastics in the melt and can provide desirable morphologies with only a
single extruder. The use of a long tube attached to the extruder was proposed as an
important means of attaining these objectives for a practical process. It soon became
evident that this was not easily accomplished because of the lack of useful knowledge of
this area in the published literature. The reason for this situation also became apparent: a
process design that permitted sampling along the length of the tube along with advanced
analytical techniques were required. Furthermore, it was realized that the results of the
investigation would also be useful to rheologists who were attempting to measure the
viscosity of polymer blends by measuring pressure drop across a capillary containing a
flowing polymer.
Thus, this work has four objectives:
i. to develop a process permitting de-mixing and morphology development during flow
of immiscible polymer blends in a tube to be investigated;
ii. to develop analytical methods to monitor polymer composition variation in the
polymer blend product;
iii. to elucidate the effect of polymer type and processing variables on morphology and
de-mixing ;
iv. to investigate the effect of morphology on viscosity-composition behavior as
measured using a capillary rheometer.
2.0 Theory
2.1 Phase morphology development in immiscible polymer blends
A polymer blend is a mixture of at least two polymers which have been mixed together
while melted. Such blends may be miscible or immiscible. Miscible polymer blends are
homogeneous at the molecular level. Gibbs fiee energy o f mixing, AG, is negative and is
defined by:
AG, = AH, - TAS, (2.1)
where AH,,, is enthalpy of mixing and ASm is entropy of mixing. For miscibility, in
addition to a negative value o f AGm, the following inequality must hold [2]:
where +i is the volume fraction of component i . In immiscible polymer blends, Gibbs
fiee energy o f mixing is positive. The entropy term for long chain macromolecules is
usually low because o f the restricted number of possible molecular confi yra t ions and the
entropic contribution cannot overcome the positive enthalpy of mixing. Thus, rnost
polymer pairs are immiscible and their melt blending usually leads to the formation of
systems in which the minor component is dispersed in the matrix phase. Ail the blends
used in this thesis are thermodynamically immiscible and were mechanically mixed in a
single screw extruder.
As mentioned in the Introduction section, the focus o f this thesis is upon morphology and
de-mixing o f polymer blends as they flow through a tube. This section will examine
morphology development. Then polymer de-mixing will be examined.
As stated above, morphology is the relative size, shape and location of each
distinguishable phase present. Aithough the dispersion of sphencal particles in a
continuous matrix is probably the most common morphology, there are many others.
Some of the possible structures for immiscible polymer blends are shown in Figure 2.1.
Many properties and subsequent uses of a blend depend on the morphology. For example,
while fibrillar reinforcement is desirable for monofilaments to obtain high uniaxil
strength, platelets oriented normal to the flow direction are required for enhancement of
barrier properties.
Blend morphology can be affected by the way in which the two components flow during
forming. Based on the degree of phase separation, there are two main types of multiphase
fiow [3]. The first one is stratified multiphase tlow, in which two polymers form
continuous phases separated fiom each other by continuous boundaries. The second is
dispersed multiphase flow, in which one polymer exists as discrete phase dispersed in the
polymer forming the continuous phase. These two types of multiphase flow are
discussed in turn in the following sections.
Figure 2.1: Possible morphological structures for an immiscible blend A) core-sheath B)
multi-layers C) dispersed phase D) mix of dispersed phase and multi-layers.
Stratified flow may be sequential or simultaneous side-by-side flow. In sequential flow,
two volumes of polymer melt are injected into a mould. An example of this is the
sandwich injection moulding process [4]. In simultaneous side-by-side flow, the two melt
phases are CO-extruded.
Several authors studied the simultaneous side-by-side extrusion of two polymer melts [5-
1 O]. It has been experimentally established that the less viscous melt always tends to
move into the region of high shear rate thus displacing the more viscous melt. If the flow
path is sufficiently long, this phenornenon results in a sheath-core configuration with the
less viscous component encapsulating the more viscous component, as illustrated in
Fi y r e 2.2. The effect is ofien attributed to the principal of energy minimization: the
multi-layer system is configuring so as to minimize its energy and pressure drop for fixed
flow rate [9].
Lower viscosity 4
b Eigher viscosi t y
Figure 2.2: Schematic representation of encapsulation in simultaneous side-by-side flow
in a coextrusion; the less viscous component encapsulates the more viscous component.
White et. al. [5] studied the effect of normal stress differences on the interface shape.
They showed that the less viscous component encapsulated the more viscous component
regardless of the relative values of the first and second normal stress differences between
the two components. They concluded that the factor dominating the interface shape is the
viscosity mismatch. This was also confirmed by other researchers [6,8,1 O].
Everage [8] studied the effect of the viscosity ratio and the length of the flow path on the
encapsulation. He showed that as the viscosity ratio increases, there is an increase in
encapsulation. He also showed that encapsulation is a very slow process and a very long
die was required for complete encapsulation to take place in a very viscous bicornponent
polymeric system.
Everage [7] explained the encapsulation observation by showing that the encapsulation
was preferred when two Newtonian (or power law) fluids flowed side by side through a
cylindrical tube or a slit because the energy of the system was minimized.
In contrast to stratified flow, dispersed multi-phase fiow refers to the usual random melt
mixing of different polymers to form blends. Morphology is then mainly determined by
drop deformation, drop breakup and drop coalescence [Il-421.
Our knowledge of drop deformation and drop breakup in polymer blends builds upon
the knowledge base established for Newtonian fluids. Thus, work in Newtonian fluids
will be examined first followed by the work in polymers.
The deformation of drops in the Newtonian systems is based on the interaction of
viscous force and interfacial force. The ratio of these two forces is defined as a
dimensionless Capillary nurnber, C,:
k = interfacial tension
q = viscosity of rnatrix
a = droplet radius
f = shear rate
Taylor [13,14] analyzed the deformation of a droplet in simple shear flow. By balancing
viscous force with interfacial force, he developed an expression for the drop deformation:
where  is the ratio of the viscosity of the dispersed phase to the matrix (or continuous)
phase (qdq,,), L is the length of the major principal axis and B is the length of the minor
principal axis of the drop. q, and q d are the viscosity of the matrix and the viscosity of
the dispersed phase respectively. Equation (2.4) shows that the larger the capillary
number, the less the droplet deforms during simple shear flow field. Taylor also proved
that no drop breakup occurred when h>25 Equation (2.4) is only valid for small
deforrnations.
Several authors have extended the observations of Taylor and generally confirmed that
maximum drop size is a function of viscosity ratio and Capillary nurnber [ l 5-1 71. Karam
et. al. [IS] and Tavgac [17] have studied Newtonian systems over a wide range of
viscosity ratio. Their results showed that there is an upper and lower limit of viscosity
ratio beyond which no droplet breakup occurs. At very high h, the shear stress exerted by
the continuous phase cannot overrome the interfacial force. However, at very low h the
drop was deformed but did not break.
The deformation of the droplet becomes unstable and will burst, when the interfacial
force can no longer balance the viscous force. The parameter descnbing the critical
condition of breakup is the critical Capillary number. C,c. When Ca< C,c, the deformed
droplet will break [18].
There is another mechanism for breakup due to capillary instability of long cylindrical
column which is shown in Figure 2.3 [18]. The kinetics of the process were theoretically
studied by Rayleigh [19] and Tomotika [20]. Tomotika [20] showed that a sinusodial
distortion grows exponentially with time and can be used as a description of degree of
instability:
where & is the distortion at t=O and q is the growth rate parameter of a sinusoidal
distortion:
where k is the interfacial tension, q, is the matrix viscosity, A is the distortion
wavelength, h is the viscosity ratio, Q(A, l ) is tabulated function and & is the initial
radius of the thread. The thread breakup take places when A=&0.81&. The time for
breakup (tb) can be expressed as:
Therefore, two important conditions rnust be provided for breakup of droplets; C,<Cac is
the requirement for breakup, but to achieve it the tirne t>tb, must be provided [18].
Figure 2.3: A schematic of sinusoidally disturbed thread [18].
In contrast to the Newtonian case, there are relatively few publications on droplet
behavior in non-Newtonian systems [12,22,27-291. Based on arguments similar to
Taylor's theory, it would be expected that deformation and breakup of drops in non-
Newtonian systems could be determined by the Capillary number and the viscosity ratio.
Vanoene [29] argued that melt elasticity also has an effect on the deformation of the
drops. The elasticity of the droplet is expected to reduce the deformation and sustain
higher cntical shear rate before rupture than purely viscous drops. The elasticity of the
matrix is expected to increase the deformation and decrease the critical shear rate before
rupture.
The effect of viscosity ratio, h, on drop size has been studied [22,23,27]. In general, a
high viscosity ratio results in coarse morphology, while viscosity ratio close to unity
results in fine morphology. Wu [22] studied the phase morphology of mbber/polyarnide
blends in twin screw extruder. He showed that the dispersed rubber phase can breakup
even at D4. He also found that Capillary number relate to the viscosity ratio according to
the following equation.
D = dispersed phase average particle size
)C = viscosity ratio
= average shear rate
Tlm = viscosity of matnx
k = interfacial tension
Minus sign (-): when h<l
Plus sign (+): when Dl
Regarding Equation 2.8, the deformation and breakup of drops during mixing are
enhanced by small interfacial tension, high matrix viscosity and large shear rate. For the
same processing conditions the minimum size will obtained when the viscosity ratio is
close to unity.
As in the case of drop defonnation and drop breakup, the main body of knowledge
available for coalescence is with Newtonian fluids. When the droplet concentration is
increased the coalescence must be taken into account and its effect on final morphology
needs to be considered. Coalescence occurs not only in the flow but also in quiescent
systems. Several authors [30-321 have studied coalescence in Newtonian systems. In the
case of quiescent coalescence, it has been reported [3 11 that the contact time required for
drop coalescence is proportional to drop diameter, the matrix viscosity and the density
difference between the drop and the matrix. Flow induced coalescence of two Newtonian
drops is pictured as shown in Figure 2.4 [30].
Collision Film Draining
Film Rupture
Figure 2.4: Coalescence of two Newtonian liquid droplets.
First, the two droplets corne close to each other or aggregate, remaining as a pair of drops
rotatinç in flow field. The film between the two drops drains until the surfaces of the
drops contact each other. Then, film rupture followed by coalescence occurs.
Most coalescence research in polymer blends has been done for either solvent cast
blending 1331 or melt blending under quiescent conditions [34-361. Fortelney et. al. [34]
found that the amount of coalescence in melt blending, under no mechanically induced
flow, decreased significantly if the viscosity of the matrix was above a critical value and
the volume fraction of the dispersed phase was below a cntical value. Also, some studies
have shown that there are large coalescence effects during annealing of blends [33,36].
Flow induced coalescence of polymeric systems is similar to Newtonian systems.
However, the frequency of occurrence may be less [35]: elastic recoil may cause
polymeric drops to separate dunng the collision step and the high viscosity of the matrix
may cause a long draining time before the film ruptures.
Many experimental studies have shown that the final size of the droplet is usually lager
than predicted by Taylor's theory for polymer systems because of coalescence effects and
the deviation increases with increasing the concentration of the dispersed phase
[23,24,3 5,3 7,401. The coalescence can be accelerated by the same factors that facilitate
the breakup of the droplets such as higher shear rate and lower viscosity of the dispersed
phase. Tokita [39] suggested that equilibrium droplet diameter in blending is a result of
continuous breakup and coalescence of the dispersed phase. Willis et. al. [40] have shown
that during the mixing process of immiscible polymer blends the dispersed phase
experiences a combination of both particle breakup and coalescence. They showed
quantitatively that the particle size increases with the concentration of the minor phase.
This effect is due to the greater probability of phase contacts and particle-particle
interactions. Coalescence results in a coarse morphology. Favis et. al. [2 1,40,42] showed
that coalescence can be significantly suppressed i n the presence of interfacial modifier
(compatibilizer) which reduce the interfacial tension. The phase size decreases with
adding interfacial modifier and also the dispersed droplets become physically separated
from each other during a given collision, by the interphase region, as a result the
coalescence decreases.
Ostwald ripening is another mechanism responsible for coarsening of dispersion [18].
The small droplets have higher interfacial energy than the larger ones. This difference
gives rise to the difision of molecules from smaller drops to the larger ones. Also, the
total number of the minor phase droplets in the system decreases with time.
2.2 De-mixing in immiscible polyrner blends
De-mixing is the migration of polymer components during melt blending under flow to
increase heterogeneity. Other terms have also been used in the literature [43-451 to
describe de-mixing: shear-induced migration, shear-induced segregation and field-
induced phase fiactionaiion for example. The term "shear-induced migration" is
particularly descriptive of observations here and will be frequently used in this thesis.
De-mixing occurs mostly under a non-homogeneous flow field. Couette and tubular
flows represent non-homogeneous fields and are common in most polymer operations.
Since this work deals with multi-phase polymer flow systems in a tube, this section
examines de-mixing during tube flow. Also, elastic effects, although possibly
contributing to de-mixing in immiscible polymer blends, are considered to be much less
important than viscous effects in the conditions examined here. This topic is the subject
of recent Ph.D. thesis work here [46].
During flow in a tube, the shear stress and shear rate are not homogeneous: they Vary
with tube radius. The high shear rate region is located near the wall of the tube while the
low shear rate region is located at the center of the tube. Figure 2.5 shows the shear rate
distribution and velocity distribution for laminar flow of a Newtonian fluid in a tube.
Vclocity distribution x
Figure 2.5: Showing the Poiseuille flow in the tube for a Newtonian fluid [75].
When a single polymer is present (either as a melt or in a solution of srnaIl molecules) it
has been suggested that the presence of a stress field (such as in tubular flow) in a
polymer system perturbs the average chain dimensions, elongating the chains in the flow
direction. The system tends to counterbalance the situation, a thermodynamic force wants
to restore the unperturbed dimensions of the macromolecules. This driving force will
cause the chains to diffise fiom regions of high stress to those of low stress [44,47,48].
For example in a polyrner solution flowing through a capillary, the solution near the wall
would, according to this argument, become depleted of polymer, thus reducing its
viscosity. The depleted polymer solution near the wall would lubricate flow of the bulk
fluid in the capillary, and the macroscopic consequence of the depletion layer would be
an apparent slip velocity. The existence of the slip velocity has been experimentally
observed [49-5 21.
Shear-induced migration in polymer solutions has been reported [53-6 11. For example,
Tirrell 153,541 showed that the composition inhornogeneity in polymer solutions
increased with flow rate and the viscosity ratio ( ~ p l F d ~ s o i w n t ) . They concluded that
conditions for observing the detectable extent of migration are achievable if a large
residence time is provided in long tubes and the length scale over which diffusion must
take place is small.
The most characteristic result of molecular migration in flow of a single polydisperse
polymer should be the creation of a molecular weight gradient across the flow cross-
section [62-651. Schreiber et. ai. [62] extnided polyethylene through a long capillary, and
found more of the low molecular weight in the outer section compared to the whole
extrudate. They attributed this to molecular fiactionation induced by flow.
Migration of low molecular weight additives, such as lubncants and slip agents in
polymers during processing, has been reported [66-691. An example of such a
phenomenon is the flow of plasticized PVC where, the additive rnigrates to the wall of
the tube, forming a lubricant layer 1681.
It has been experimentally shown that the interface shape between two polymers changes
during side-by-side flow of two polymer phases in CO-extrusion, when two polymers have
a different viscosity [5- 1 O].
There are only a few publications in the literature dealing with shear-induced migration
during dispersed two-phase flow of immiscible blends in a capillary flow [43,70-731.
White et. al. [70] studied dispersed two-phase flow of HDPEPS blends in a capillary
instrument @=1.37mm, maximum L/D=50). They showed that there was a greater
quantity of the lower viscosity component (PS) in the surface region of the extrudate than
in the core region by using Gel Permeation Chromatography (GPC) analysis. The effect
was small in terms of composition variation across the extrudate.
Utracki et. al. [43] showed the existence of shear-induced migration during melt blending
of HDPERA-6 blends in a capillary rheometer. They studied the morphology of the
extrudate in the center and the surface. They showed qualitatively that at lower
temperature (1 50°C) the concentration of PA-6, the higher viscosity component, was
higher in the center than at the surface, while at higher temperature (250°C) the
concentration of PA-6, the lower viscosity component, was lower in the center than at the
surface. Lohfink et. al. 1721 also showed qualitatively that shear-induced migration took
place for PP/EVOH blends in capillary extrusion process.
Maclean [74] used an energy approach to compare two cases of bicomponent flow. One
of the two cases is a sheath-core configuration and the other is a single interface
configuration. He concluded that the sheath-core configuration has a lower energy than
the single interface configuration. So, the sheath-core configuration is the energetically
preferred configuration.
Another mechanism for shear-induced migration is based on the forces applied on the
particles [44,75]. Figure 2.6 illustrates a particle in tubular shear flow. There is a shear
stress difference across the particle because of the velocity gradient in the tube. As a
result, there will be a driving force to rnove the particle to the center of the tube.
Figure 2.6: Showing a particle in tubular shear 80w [44].
However, as evident kom the above cited published literature, the most common
explanation for the migration of the low viscosity polymer cornponent in a blend towards
the high shear rate region of the flow is that the system attempts to minimize it's energy
Predicting Polymer De-mixing
Flow of a polymer blend in a tube is visualized as taking place in a series of radial
concentnc layers. Shear rates are considered sufficiently low that isothermal conditions
exist. Initially, at the entrance to the tube, the viscosity of each layer is identical.
However, the presence of the velocity profile results in a shear rate gradient and a
consequent gradient in the rate of viscous dissipation.
The tendency of polymer viscosity to Vary with shear rate is most often accommodated
by using the constitutive equation well known as the "power law":
where r = shear stress, = shear rate, k and n are power law's constants. Then viscosity
The derivation of shear rate profiles for concentric flow of two power law fluids in a tube
has been previously published [IO]. Appendix 1 extends the analysis to any number, i, of
concentnc layers and derives the following expression for the rate of viscous energy
dissipation per unit length of the tube in the ith layer where power law constants of the
fluid in that layer are ki and ni:
where
Si = T ; - ~ / T ~
The total rate of viscous energy dissipation per unit length of tube in al1 m layers is:
For a Newtonian fluid, k i ~ i , the Newtonian viscosity in the ith layer and ni=l.
Therefore:
Equation 2.14 can also be written for each component if each alone was in the tube as:
The rate of viscous energy dissipation (E) is equal to the rate of the work done on the
fluid by viscous force. E is a measure of the rate that viscous energy is being converted
into thermal energy and it is an irreversible conversion [76]. A fluid flow such as tubular
flow tends to minimize its rate of energy dissipation for a fixed flow rate [77]. To
illustrate this for tube flow we consider the three cases s h o w in Figure 2.7. In case 1, the
high viscosity component (component 1) is located in the core (low shear rate region), in
case II, the high viscosity cornponent is located in the shell (high shear rate region) and in
case III, as a reference, the whole tube is occupied by cornponent 3 (qs). It is assumed al1
components are Newtonian fluids. The effect of interfacial tension is not considered and a
constant total flow rate is also assumed in each case.
The equations derived in Appendix 1 expressing the ratio of total rate of viscous
dissipation in case 1 to that of case III and case II to that of case III are:
Combining Equations 2.17 and 2.18:
Where $i and cb2 arc the volume fraction of component 1 (the high viscosity component)
and component 2 (the low viscosity cornponent) respectively. X* is the ratio of the
viscosity of the high viscosity component to that of the low viscosity component
( L * = ~ '1). a is the ratio of the viscosity of the high viscosity component to that of
component 3 (a=ql/qi). Equation 2.19 shows that EI/E~I is independent of the reference
viscosity used for case III.
Figure 2.8 to Figure 2.10 show variation of the rate of viscous dissipation with h* at
different values of the volume Fraction of component 1 (41). It can be seen (particularly in
Figure 2.10) that for each value of viscosity ratio (h*), case I has a lower rate of viscous
dissipation energy than that of case II. Thus, case 1, where the lower viscosity component
is in the high shear region near tube wall, is the energetically preferred confi yration. A
flowing blend with mismatched viscosity in the tube has a tendency to minimize its total
rate of viscous energy dissipation for a fixed flow rate. Thus, to accomplish this the lower
viscosity component migrates to the wall.
Casc 1 Case KI
Case II (1): High viscosity component Csisc ITI (2): Low viscosity component (3): Componçnt 3 R: Tube radius rll and rtz: Radial position of interface for case 1 and casc iI
Figure 2.7: A schcrnatic of sheath-core morphology for comparison: Case 1) high viscosity component is located at low shear rate region (core), Casc II) high viscosity component is located at hiçh shear rate region (shell) CaselII) a monocomponent alone is occupied the whole tube.
Figure 2.8: Er/Eiir versus h* as a fbnction of volume fraction of component 1 (from 4 ,=0.2 to 41~=0.7 with 0.1% interval) for a=l, log-log scale.
Figure 2.9: Eii/Eiii versus A' as a function of volume fraction of component 1 (fiorn 4,=0.2 to &=0.7 with 0.1% interval) for a=l, log-log scale.
Figure 2.10: Ei/Eii versus h* as a fùnction of volume fraction of high viscosity component (fiorn 41~=0.2 to Oi=O. 7 with 0.1 % interval), log-log scale.
2.3 Rheological property of polymers
2.3.1 Viscosity
As was mentioned in the previous section, the power law is fiequently used to describe
the viscosity of polymers. For a Newtonian fluid, viscosity is not a function of shear rate
but it is a function of temperature. For a Non-Newtonian fluid, such as most polymers,
viscosity is a function of both shear rate and temperature [78,79]. Viscosity (at fixed
shear rate for shear thinning fluid) decreases with increase of temperature. At low shear
rates, and at temperatures much higher than the glass transition temperature of the
polymer, the relationship between the viscosity of polymer melts and temperature can be
represented by an Arrhenius relationship 1801 :
where A and E are constants, the later being the activation energy for 80w in J/mol, R is
the gas constant in J/mol K and T is the absolute temperature in degrees Kelvin.
2.3.2 Capillary flow
A capillary rheometer is the most popular type of apparatus used to measure polymer
melt viscosity because it is very similar to common polymer processing flows in extruder
dies and pipes [81,82]. The capillary rheorneter uses pressure-driven flow and the flow is
generated by pushing the fluid from a reservoir. In this work, a single screw extruder is
used as a capillary rheometer by extruding into a long tube. Pressure drop and flow rate
through the tube are used to determine the viscosity.
In deriving the viscosity relation the important assumptions are as follows:
1. Flow is fùlly developed, steady state and isothermal.
2. There is no velocity in the r and 0 directions.
3. There is no slip at the walls and the velocity of the fluid at the wall of the tube is
zero.
4. Fluid flow is time-independent.
5. The melt is incompressible.
With these assumptions, the shear stress distribution is as follows:
where r, is the shear stress in the z-direction, t, is the wall shear stress, r is the radial
position fiom the capillary center, AP is the pressure drop across the capillary length L
and R is capillary radius.
The shear rate can be calculated from the volumetric flow rate, Q:
In order to determine the true shear rate at the wall for non-Newtonian fluids, the
Rabinowitsch correction must applied. For the case of a power law fluid behavior, it can
be shown that the true shear rate at the wall is given by :
It is clear that the quantity ( ~ Q I ~ R ~ ) , which is equal to the wall shear rate in the case of a
Newtonian fluid, no longer has this significance when the fluid is non-Newtonian. It is,
however, often referred as the "apparent wall shear rate". The symbol Y., is used for this
quantity.
Combining Equations 2.22 and 2.24 and considering the definition of the viscosity, it can
be shown that
g[Y] against ~ o g [ s ] Values of n and k can be obtained by regressing LO -
End effects
In capillary rheometers, the driving pressure in the reservoir is used to calculate the wall
shear stress values. However, because the barre1 diameter is much larger than the
capillary diameter, the flow of the polymer meit reaches a fully-developed state only after
traveling some distance from the entrance of the capillary [8 11. As the melt flow has
velocity components in directions other than that in the length direction of the capillary,
the measured pressure drop is larger than the true pressure drcp for fully-developed flow
(entrance effect).
The Bagley end correction is used to account for end effects in capillary rheometers.
Bagley [83] showed that a plot of AP vs. (L/D) at constant shear rate is linear and that
extrapolation of such a linear plot to the abscissa (Al?=O) gives the length over which the
flow is not fùlly developed. The shear stress can then be corrected for end effects
using :
where e is the magnitude of the negative intercept.
2.3.3 Flow behavior of immiscible polymer blends
The melt rheology of immiscible polymer blends has been studied extensively [84-901. In
order to predict the rheological properties of the blend from known rheological properties
of the components, different blending laws have been proposed [9 1-94]. It has been
shown that the melt flow curves generally lie between those of the pure homopolymer
components [95-991. However, it also observed that the flow curves of the blends may
ofien be lower or higher than both the pure components [100,101]. The melt viscosity -
composition behavior of immiscible blends at constant shear stress has been studied
[85,97,99,102,103]. Linear, negative deviation and positive deviation behavior was
observed depending on the viscosity ratio and type of flow. This diverse rheological
behavior may be reflecting an interaction between the flow conditions, the local viscosity
ratio of the components and the morphology as well as the state of de-mixing of the
blend. Assuming de-mixing negligible and focusing upon morphology as the cause of the
different observed variation of viscosity with composition, two mixing rule models are
examined: Lees' model and a sheath-core model.
1) Lees' model
Lees [104] assumed that the dispersed phase is oriented in the flow direction. This is
shown schematically in Figure 2.11. He also assumed that there is a constant shear stress
everywhere in the flow and there is no slip between the layers.
Figure 2.11: Showing orientation in the flow direction for a binary blend.
The average shear rate in a volume element based on the volume fractions is:
The stress is the same in al1 layers so that
where 4, and 4 1 ~ are the volume fraction of component 1 and component 2 respectively.
Lees' model predicts a rnonotonic variation of the reciprocal viscosity with volume
Fraction.
2) S heath-core model
For this model a sheath-core structure is assumed for an immiscible polymer blend: the
high viscosity fluid is located at the center of the tube (low shear rate area) and the low
viscosity fluid is located at the wall of the tube (high shear rate area), which is show in
Figure 2.12. The fluids are assumed Newtonian and incompressible. Flow is assumed
fully developed, steady state and isothermal. It is also assumed that there is no slip at the
wall of the tube and at the interfaces.
(1): High viscous fluid (2): Law viscouu fluid R: Tube mdius b: Radial position of intcrfacc
Figure 2.12: Sheath-core configuration in the flow direction for a binary blend in a tube.
The axial mornentum equation is [IO]
Combining two equations and applying the boundary conditions give the velocity profiles
The volumetric flow rates are found fiom Equations 2.33 and 2.34:
where
The total volume flow rate, c., then is
Q, = QI + Q2
From the macroscopic point of view, it is known for a fluid of the same density of
viscosity q~ that
By setting Qe = Qt, we can solve for the effective coefficient of viscosity
Combining Equations 2.41 and 2.42
It can be seen from Equation 2.43 that the sheath-core model predicts a stronger
dependence upon volume fraction of components in the blend than does Lees' model
(Equation 2.30).
Slip at the component interîace
As mentioned earlier, viscosity-composition behavior of some immiscible polymer
blends at a constant shear stress showed a negative deviation from the mixing rules.
Neither Lees' mode1 (Equation 2.3 0) or the sheath-core model (Equation 2.43) predicts
this type of behavior. Slip between the two immiscible polymeric phases may be the
cause. The consequence of slip is a higher flow rate at constant shear stress and hence a
lower viscosity. Lin [105, 91,921 assumed shear-induced interlayer slip between two
unlike adjacent compositions with concentric layers morphology and he introduced
interlayer slip factor, P:
For a given temperature and shear rate:
Where rw is the wall shear stress and piz is the characteristic slip factor, a constant for a
given system. PI2 has the same units as shear stress and was shown by Lin [105] to be
related to the frictional force between the interfaces of the two unlike adjacent polymers
Lin's theory predicts the negative deviation fiom mixing rule, independent of viscosity
ratio, caused by the interlayer slip. When P is equal to unity the mixing rule (Equation
2.30) results.
In this work, the interlayer slip factor is used to modify sheath-core mode]:
The shear stress at the interface is used as the characteristic slip factor. Using Equations
2.21, 2.22 and 2.42:
Where R is the tube radius, b is the radial position of the interface and $i is the volume
fraction of high viscosity component (component 1).
2.4 Analytical methods
2.4.1 Fourier Transform Infrared Spectroscopy (FTIR)
FTIR spectroscopy has become an independent tool in characterizing both organic and
inorganic compounds. This section begins with a brief description of spectroscopy
fundamentals and then qualitative and quantitative analysis of FTIR spectra for
polymers will be discussed.
Spectroscopy is the measurement and interpretation of electromagnetic radiation
absorbed or emitted when atoms, ions or molecules change energy level.
Electromagnetic radiation, which is energy transmitted through space in the form of
waves, can be classified according to its wavelength and fiequency. The frequency of
infiared radiation is usually defined to be between 10 cm" to 12900 cm-' [106]. This
range can be divided into far infrared, mid-infrared and near infrared. The mid infiared
region (the 400 -4000cm" range) is often used for characterization of polymers. Infiared
spectroscopy is based on the absorption of radiation in the infrared fiequency range due
to the molecular vibrations of the functional groups contained in the polyrner chain.
Infiared spectroscopy has long been recognized as a poweriùl tool for polymer
characterization [107]. Prior to FTIR, infiared spectroscopy was carried out using a
dispersive instrument utilizing prisms or gratings to geometrically disperse the infiared
radiation.
The increased speed, higher signal-to-noise ratio [108], more precise fiequency [log]
and higher infiared beam throughput [110] of FTIR relative to the dispersion infiared
has led to a substantially greater number of applications of infiared in polymer research
especially in quantitative analysis. For example, the time advantage of the FTIR has
been particularly important for the study of polymerizations [111], degradation
processes [ 1 121 and other time-dependent processes of polymers [ 1 1 3,1141.
Qualitative analysis
Absorption bands corresponding to certain atomic groups appear in certain spectral
fiequencies characteristic of the given atornic group. The frequencies producing these
bands are called characteristic band and group frequencies. Group fkequencies depend
on the nature of the atoms, the nature of the chemical bonds and the structural
characteristics of the molecule.
The position of the characteristic peak fiequency in a spectrum is indicative of the type
of the chemical group present. The IR spectra of some polymers have been studied in
detail and the absortion bands assigned to some fundamental factor group vibrations
[Ils]. For example, Varsanyi [Il61 found that the very strong absorbtion band at
approximately 700 cm-' in a polystyrene spectrum is due to the benzene ring in the
polymer.
The group frequencies and distinctive patterns in the "fingerprint" region of the
spectrum can be used for qualitative identification of polymers. Identification of
polymeric materials by modern cornputer-assisted instruments is now ofien readily
accomplished: libraries of standard spectra stored in the computer rnemory and used in
conjunction with search routines have greatly simplified the process.
Quantitative analysis
The quantitative analysis of polymers in transmission studies is based on the Beer-
Lambert law
A(v) = a(v)bc (2.48)
where v is the wavenumber, A is the absorbance at v , a is the absorptivity at v, b is the
sample thickness and c is the concentration of the cornponent of interest. For a
multicomponent mixture, the total absorbance at v is given by:
where subscript i refers to the ie component of the mixture and it is assumed that there is
no interaction between the i components.
Deviations from Beer's law have been observed in some cases. Instrumental eflects,
such as stray radiation, insuscient resolution and non-uniform samples and chemical
effects, such as dissociation and molecular interaction, can cause deviations. For modem
FTIR spectrometers, the most important source of error apart from the chemical efTects
is resolution [l 17,1181.
Anderson et.al. [119] showed that the non-linear behaviour is worst at higher absorbance
and narrower bands. This finding means that absorption bands with an absorbance of
less than 0.7 are best used for quantitative analysis.
In this work the characteristic peak and the range of concentration for HDPE/PP blends
were chosen in such a way that there was no deviation from Beer's law and Equation
2.48 can be used (see Section 4.2). The absorbance at a characteristic peak of PP and a
characteristic peak of HDPE is correlated to the polymer composition. Simple linear
regression (SLR) was applied to find the calibration curves and various statistical
measures associated with SLR (see Appendix II) provided a measure of the goodness of
fit of the straight line to the data and precision of predictions. Other least square methods
which have been used by spectroscopists [120-1241 include classical least square (CLS),
inverse least square (ILS), principal component regression (PCR) and partial least square
(PLS). These approaches were ~ o t necessary in this work.
2.4.2 Differential Scanning Calorimetry @SC)
DSC is a thermal analysis technique. Thermal analysis (TA) is the tem applied to a
group of methods and techniques in which a physical property of a substance is
measured as a hnction of temperature while the substance is subjected to a controlled
temperature progam [125]. DSC is defined as a technique in which the differences in
energy inputs into a sample and a reference are measured as a fùnction of temperature or
time, while they are subjected to a controlled temperature program. The reference is a
thermally inert substance, which has no phase changes in the temperature range of the
experiment .
A signal provided by a DSC instrument is a direct measure of the excess heat absorbed
or released by the sample. A DSC spectrum is a plot of heat flow rate as a function of
temperature. When a polymer is heated, many processes may occur in the polymer.
Those pertinent to thermal analysis include glass transition, crystallization, melting,
degradation, crosslinking and oxidation. DSC is one of the easiest of modem analytical
techniques to use: minimum sample preparation is required and quantitative information
can be obtained fiom only a few milligrams of material. Some of the applications of the
DSC for polymers, especially blends, will be discussed in the following sections.
Glass transition temperature
DSC is used to measure g las transition temperature (Ta of polymers. The DSC used
with a step change in the heat flow detects the Tg. At the Tg of polymers there is a heat
capacity change due to the change of mobility of polymer chain segments.
One of most widely used methods to check the miscibility of polymer blends is based on
the measurement of the Tg. The detection of a single transition temperature whose value
falls somewhere between the Tg's of the component polyrners is an indication of a
miscible system [126-1281. Conversely, a non- miscible system will show two Tg's
corresponding to the T,'s of the individual components [129]. For completely
immiscible polymer blends, the Tg's of the two phases in the blend should be equal to
the Tg's of the respective pure polymers 11301. A shift in the T, of the polymer in the
blend with respect to the pure polymers may indicate some degree of miscibility [13 11.
Attempts have been made to relate the T, of a miscible blend to blend composition
[ 132,1331. A few miscible blends such as PVC in EVA have been found to exhibit Tg
versus composition dependencies that cm be predicted by the following equation:
where Wi and Wz represent the mass fractions of the components and Tg, Tgl and Tg2 are
the Tg's of the blend, component 1 and component 2 respectively. However, most
miscible blends do not obey the above equation and other equations have been proposed
[134-1361.
Melting temperature and crystallinity
Low molecular weight materials have a single valued sharp melting point. This is not the
case in polymers. Melting takes place over a wide temperature range and the melting
point (Tm) is ideally taken as a temperature at which the last trace of crystallinity
disappears. This is the temperature at which the largest and most perfect crystals are
melting. The most popular thermal technique for determining Tm is DSC, which also
gives the enthalpy or heat of fusion [137]. The peak value of the melting endotherm is
frequently assigned as Tm.
The area under the melting curve (A) in a DSC spectrum is proportional to the heat of
fusion (AHf) according to the following relationship:
K is a ce11 constant, m is a sample weight and q is a heating rate.
With the proper calibration of the instrument cell, DSC then can be used to determine
polyrner heats of fusion. It must be noted that the value of AHf, detennined by DSC,
depends on the amount of crystallinity present in the sample. In the case of completely
immiscible blends, the melting point is independent of the composition but the heat of
fusion is proportional to the composition [138].
The polymer blends that were used in this work are immiscible. Each polymer has a
distinct melting peak and there is no overlap between the peaks for each combination of
two polymers. The idea of using DSC to correlate the area under the melting peak to
polymer composition to obtain the calibration curves for quantitative anaiysis
(composition) will be discussed in detail in Section 4.2. Chemical structure, molecular
weight, chain branching and thermal history can affect the crystallization and the
melting point [ 139,1401. For example, the molecular weight influences the degree of
crystallinity that can be achieved. As molecular weight increases, the degree of
crystallinity decreases until the limiting value of about 25-30% is reached. This
reduction in crystallinity may be attributed to entanglements that restrict the
crystallization process [14 11.
Different grades of PP and HDPE with different molecular weight were used in this
work. A different calibration curve (a plot of area under the melting peak versus blend
composition) was needed for each pair of polymers due to different degrees of
crystallinity. Linear regession was used to provide quantitative results (Appendix II).
3.0 Experimental
3.1 Materials
Various grades of high density polyethylenes (HDPE) and polypropylenes (PP) and one
grade of nylon6 (PA-6) were used. Al1 materials were provided in pellet form. Three
blend systems were employed in this work which is s h o w in Table 3.1.
Table 3.1: Blend systems.
System Component 1 Component 2
1
Pertinent properties of the materials are summarized in Table 3.2.
(PP/PA-6) 3
(HDPE/PA-6)
Table 3.2: Pertinent polymer properties.
HDPE 1 PP 1, PP2, PP3
PA-6 HDPE 1, HDPEZ
Name
HDPE1
HDPES
Manufacturer
Fina Oil & Chernicals
I
PP2
PP3
PA-6
Trade Name
7208 L
JE 6100 PP 1
0.952 Dow Chernical
Montel Polyolefins
Amoco Chernical Company
Shell Chernicals
Alliedsignal Plastics
Density ( g l d
0.959
IP-40
0.904
0.77
Amco7934
SM 6100
8209F
MeIt Density (glcm3)
0.77
40 1
MF1 Dl238
(gllomin) 0.6
0.73
0.908
0.905
1.135
2.1
0.73
0.73
0.92
35
11
4.5
3.2 Processing equipment
Process 1
Figure 3.1 shows a schematic picture of Process 1. It consists of % " single screw
extruder (C. W Brabender, Mode1 1629) equipped with a PL-2000 computet-ized
plasticorder and a metering screw with an LID ratio of %/l and a transition piece
attached to the end of the extruder. The length of the transition piece is 10 cm. The
beginning of the transition piece has a diarneter of 18 mm, but the diarneter reduces
suddenly (90 deg.) to 10 mm at the length of 5 cm. Process 1 was used to study shear-
induced migration and phase morphology development of HDPEBA-6 blends during
flow in the extruder alone. Since PA-6 absorbs water, a11 materials were dried at 90°C in
the oven under vacuum for 48 hours. Pellets were dry blended in a plastic container and
then transferred to the hopper. The hopper was covered to prevent PA-6 from absorbing
water. The level of rnaterial in the hopper was maintained at about ?A of the heiçht of the
hopper in al1 the mns. These procedures were done for al1 experiments in this work. The
extruder was tumed off, when the systern reached steady state. The transition piece was
cooled by cold water and samples were taken from the transition piece for analysis. The
operating conditions are shown in Table 3.3.
Hopper
Transition Piece
Figure 3.1: A schematic of Process 1.
Table 3.3: Operating conditions of Process 1 and Process 2.
Polymer 1
Polymcr n
HDPEI
Corn. of Polymer 1
(wt%)
Tcmp. Scrcw f mess
': Run that is perfomed iwice
Process 2
Process 2 is the same as Process 1 , except that a static mixer (Kenics type) and a breaker
plate were attached to the end of extruder. A schematic pichire of Process 2 is shown in
Figure 3.2. The operating conditions are shown in Table 3.3. The purpose of using
Process 2 was to study the phase morphology development of immiscible polymer blends
entering the tube.
Hopper
Plate
Static Muer
Figure 3.2: A scliematic of Process 2.
Process 3
A schematic picture of Process 3 is shown in Figure 3.3. It is similar to Process 2, except
that a long tube was added to the end of process 2 and two pressure transducers (Gefran
Mode1 SP830ISPSM/H) were installed to monitor the pressure drop along the tube
(Section 4.1). Temperature control is achieved along the length of the tube using P D
control from an extemal multiloop teniperature controller.
Process 3 was used as a capillary rheometer to measure melt viscosity of pure polymers
(HDPE, PP) and to study viscosity-composition behavior of HDPEiPP blends and
HDPE/PA-6 blends. Screw speed, pressure (P2, P3), pressure transducer positions, die
temperature, melt temperature, and mass flow were recorded for each run. When the
system reached steady state, five extrudate samples were taken over measured time
intervals and averaged to determine the mass flow rate. Different shear rates were
obtained by cnmging the screw speed.
In the case of measuring melt viscosity of pure homopolymers, extrusion temperatures
were 155OC, 195*C, and 195°C in the extruder barre1 zones and 195°C in the transition
piece. The static mixer and the breaker plate were not installed between the extruder and
the tube. Temperature along the length of the tube was 2 2 0 ' ~ and L/D was varied.
The operating conditions for studying viscosity-composition behavior is shown in Table
3.4 and Table 3.5.
Process 3 also used to study shear-induced migration and phase morphology development
during the flow of immiscible polymer blends in the tube for three immiscible polymer
blend systems. Table 3.6 shows the operating conditions for this case. A schematic
picture of the tube along with the positions from which samples were taken for analysis is
shown in Figure 3.4.
The cooling procedure used depended on the method of studying morphology. To study
morphology development along the length of the tube, the transition piece after the static
mixer and the whole tube were cooled by cold water. Othenvise, only sections 1 1,12,13
were cooled (see Figure 3.4). It was estimated that under these conditions the centre of 8
mm tube was cooled to l OO°C in 10 S. Thus, it was considered unlikely that cooling rate
was a significant source of error. In addition, the extrudate samples after the tube, for al1
the runs, were cooled by cold water when the system reached steady state. The samples
were stored in a freezer to prevent morphological changes with time.
3.3 Analytical methods
Quantitative Fourier Transform Infrared Spectroscopy (FTIR)
FTIR spectra were obtained using a Matson Galaxy 6020 FTIR spectrometer. The system
operated in transition mode with 60 scans at 4 cm-' resolution fiom 400 to 4000 cm". The
FTlR spectrometer was interfaced with a computer equipped with a FIRST Utility
Mattlib software package used to collect and process the data. Specteral peaks were
analyzed using the Winfirst software package. The FTIR was used to obtain calibration
curves to determine polymer composition for HDPE/PP blends. These calibration curves
were used to quantifi shear-induced migration.
Quantitative Differential Scanning Calorirneter (DSC)
A Dupont, mode1 9 1 OC DSC was used to obtain polymer composition. Samples were
weighed to approximately 5-10 mg and placed in a 6.5mm D X 1 .Omm H aluminurn pan.
The lid of the pan was crirnped on with a press (Dupont Encapsulation, Mode1 000733).
The DSC interfaced with a computer and a Real-time Plot, Utility Program. Each sample
was heated to either 200°C or 250 O C (depending upon the blend system) to destroy its
thermal history at a heating rate of 10 "C/min. It was held at that temperature for 5 min,
slowly cooled to room temperature and then heated again at the same heating rate to
obtain the second DSC spectrum. The DSC spectra were analyzed using the
accompanying General Analysis, Utility progam.
Table 3.4: Operating conditions for studying viscosity-composition behavior of HDPE 1 /PP3 blends using Process 3.
Table 3.5: Operating conditions for studying viscosity-composition behavior of
Shcar Stress W P ~ )
0.02
0.05
0.08
KDPEIPA-6 blends using Process 3.
HDPEl (W%)
0,20,40,60,80,100
0,20,40,60, 80,100
0,20,40,60,80,100
Tcmp. (OC)
Polymer Type
HDPE (wt%)
Shcar Strcss (Mpa)
HDPEl 0.06
Screw Specd (r.p*m)
LID
Table 3.6: Operating conditions for studying shear-induced migration and phase morphology development using Process 3.
System Temp. ("C)
I-IDPE/PP 195,2 10,230 blends
PP/PA-6 blends
Composition Screw (wtO/o)
LID
25-162
25-162
25-162 HDPEPA-6 blends
Figure 3.4: A schematic of the tube along with the positions for taken the samples.
230,260 90/10,851151,80/20,70/30,50/50 5,15,30,60
Sample preparation
A total of 19 samples of polymer blends were prepared using virgin PP3 and HDPEI, in
which the concentrations of PP ranged fiom O percent PP to 100 percent PP in 5 percent
increments. A CO-rotating intermeshing twin-screw extruder (C.W. Brabender Instrument,
Inc.) with a L/D ratio of 24:l was used for melt blending of the polymers. A constant
screw speed of 5 rpm and a temperature dong the extruder of 1 9 5 ' ~ were used for al1
mns. The extruded polymer strand was pelletized and then reformed into films by
compression rnolding ( Specac Constant Thickness Film Maker). The absorbance spectra
were recorded for 19 samples of known composition with 3 replicates by using FTIR.
The blend samples for DSC analysis were used in two fons : extrudate blend samples,
mechanically mixed, directly obtained fiom the twin screw extruder and pelletized blend
samples, not mixed, directly prepared from the pure polymer pellet. The extrudate
samples were used only for HDPE VPP3 blends. Samples for quantifying shear-induced
migration were prepared fiom core region and shell region of extrudates. The area of the
core and the area of the shell were estimated based on the total area of the extmdate and
the feed composition:
A, = w,A, (3.1)
As = (1 - w,)A, (3-2)
where A,, At and A, are core, total and shell cross sectional areas respectively. wi is the
weight fraction of component 1 (high viscosity component). It was assumed that the
components had the same density.
3.4 Morphology
Staining technique
To enhance optical contrast between the dispersed phase (PA-6) and the matrix phase (PP
or HDPE), polished sarnples were stained with a O. 1 wt % dye solution of Brilliant blue.
The morphology of the extrudate frorn each of the pure polyrners was evaluated using the
staining technique to differentiate between the morphology of the pure homopolymers.
This staining technique allowed us to use optical microscopy for the morphological study
of HDPEPA-6 and PP/PA-6 blends. To enhance contrast between PP phase and HDPE
phase in HDPE/PP blends, several different staining techniques were tried. However,
none of the staining techniques were successfùl. Therefore, optical microscopy was not
used to study blend morphology for HDPE/PP blends. In this case the morphology was
obtained using scanning electron microscopy (SEM) on the fiactured surfaces.
Optical Microscopy (OM)
A Nikon-AFX-DX optical microscopy was used to study morphology of HDPE/PAd and
PPPA-6 blends. Samples were cut in two diflerent directions: parallel to the flow
direction and perpendicular to the flow direction. They were then polished to a smooth
surface. The polished samples were stained with a 0.1 wt % dye solution of Brilliant blue.
In the following figures, which show the morphology of the blends in this work, the PA-6
phase appears white in a black PP or HDPE matrix. The tube radius is 4mm. It was not
always possible to photograph the sample beginning near the wall or beginning near the
centre. In such cases, at the side of the photograph, the starting radial distance is shown
for the centre and the ending radial distance for the wall.
Scanning Electron Microscopy (SEM)
A Hitachi S-520 SEM was used to study the morphology of HDPE/PP blends on
fractured surfaces. The fractured surfaces were prepared in liquid nitrogen and vacuum
coated with a thin layer of gold.
3.5 Melt viscosity
Capillary rheometer
A Rosand Mode1 RH741 capillary rheometer using a lmm diameter capillary with a
length of 16 mm was used to measure melt viscosity of pure polymers. Measurements
were conducted at McMaster University in Hamilton, Ontario, Canada. The melt
viscosities were measured at different shear rates at different temperatures ranging fiom
195 'C to 260 'c.
4.0 Results and discussion
4.1 Process development
4.1.1 Design requirements
The design requirements for a process to allow the study of shear-induced migration and
phase rnorphology development during flow of immiscible polymer blends in a long
capillary flow were as follows:
1. A single screw extruder will provide the input to the capillary. That is the process
will be an add-on to an extmder.
. . 11. The design will be a simple open tube. Since the residence time in the tube has a
major effect on the shear-induced migration and morphology, the tube will be
designed in such a way that its length can be changed by adding or removing a
length of tubing. The design will also allow water quenching and sampling from
different sections of the tube for off-line analysis.
.**
111. The design will allow the determination of the effect of the tube on shear-induced
migration and morphology separately from the extruder.
iv. The design will allow the control of temperature as well as the use of temperature
progarnming along the length of the tube.
v. The design will allow the measurement of melt viscosity by using the extruder-
tube combination as a capillary rheometer.
4.1.2 Shear-induced migration and phase morphology development in extruder
In a single screw extruder, shear rate is not homogeneous and the highest shear rate area
is close to the screw surface while the lowest shear rate is close to the barrel. The flowing
immiscible polymer blends can become nonhomogeneous (i.e. shear-induced migration
can occur) when the blend components have a different viscosity.
Figure 4.1 presents the morphology development of 30/70 PA-6RIDPE2 perpendicular to
the flow direction at downstream fiom the screw tip. The white color represents the
dispersed phase (PA-6) and the black color the matrix (HDPE2). The viscosity ratio of
the dispersed phase to the matrix (h) is 10.6 at the experimental conditions.
The figure shows that the average size of the dispersed phase in the core region is smaller
than the shell region. This is may be attributed to higher shear rate close to the core. One
can also notice differences in the concentration of PA-6 at both regions because of shear-
induced migration. Table 4.1 shows composition analysis along with standard deviation
(S.D.) of PA-6/HDPE blends after the extruder. There is no significant composition
difference between the core region and the shell region at a viscosity ratio close to one,
regardless of the shear rate in the system. However, at high shear rate and high viscosity
ratio, there is a significant composition difference between the core and the shell: there is
more of the PA-6, the high viscosity component, in the shell region.
43.3 mm)
Figure 4.1 : Morphology development of 3 O/7O PA-6/HDPE2 perpendicular to the flow direction at 1 cm downstream fiom the screw tip; Temp.=230 OC, fn,,=38 sec*', h=10.6 A: shell, B: core.
Table 4.1: Composition analysis of 3O/7O PA-6/HDPE blends after extruder using DSC,
Feed Composition
(wtOh)
Apparent Shear Rate
(sec-')
-
Viscosity Ratio A
~ P A - ~ ~ I F I D P E
0.62
Composition in Core f S.D.
(wtOh)
0.47
Composition in Shell + S.D.
(wtO/o)
P A 4 29.7 i 0.94 HDPE 70.3 i 0 . 9 4
8
The shear rate distribution in the transition piece after the screw tip is opposite to that of
the extruder. The higher shear rate locates close to the wall of the transition piece. Figure
4.2 shows morphology development along the transition piece parallel to the flow
direction for 30170 PA-6/HDPE2. The average size of the dispersed phase (PA-6)
decreases at the shell and increases at the core along the transition piece.
PA-6 3 1 . 1 f 1.25 HDPE 68.9 f 1.25
PA-6 30.3 f 1.17 HDPE 69.7 i 1.17
10.6
A static mixer (Kenics type) was installed afier the extruder. The purpose was to obtain a
P A 4 32.2 f 0.93 HDPE 67.8 f 0.93
P A 4 28.2 k 0.55 HDPE 71.8 IO.55
homogeneous structure as possible previous to entenng the tube. Figure 4.3 shows the
morphology of HDPEPA-6 blends after the static mixer. The dispersed phase (PA-6) is
distributed in the matrix uniformly. There is no significant difference in polymer
composition between regions. It means that there is no indication of shear-induced
migration immediately after the static mixer.
P A 4 29.1 k 0.87 HDPE 70.9 f 0.87
P A 4 23.7 i: 0.82 HDPE 76.3 I0 .82
P A 4 32.9 I 1.50 HDPE 67.1 f 1.50
Wall (r=8mm)
Core (d .5mrn)
Figure 4.2a
Wall (r=
Core (r=0.6mm)
Figure 4.2: Morphology development of 30/70 PA4RIDPE2 dong the transition piece parallel to the flow direction; Temp.=230 38 sec-', A= 10.6, A: shell B: core, a) L=3.5 cm b) L4.8 cm.
Wall
Core
B
Figure 4.3a
Wall
Core
Figure 4.3: Morphology development of 30170 PA-6RIDPE blends d e r static mixer pamllel to the flow direction; y',= 3 i sec", Temp.=230°C, A: shell B: core (a) 30/70 PA-6MDPEl h= 0.64 (b) 30/70 PA-6MDPE2 h = 10.2 1.
4.1.3 Tube design
The designed tube is composed of individual pieces which screw together to make up the
entire length. Each piece is in effect a small capillary die 10 cm length with inner
diameter of 0.8 cm and outside diameter of 4 cm. Figure 4.4 shows a schematic picture of
one piece. All, or some of the pieces, can be used such that L/D ratio can be varied.
Therefore, this design allows taking samples from each segment to study shear-induced
migration and phase morphology development along the length of the tube.
Temperature control is achieved along the length of the tube using PID control fiom an
extemal mult iloop temperature controller. There are twelve temperature zones along the
length of the tube. Type J thermocouples contact the outer die wall beneath twelve band
heaters spaced uniformly along the die length.
< > 1 Ocm
Figure 4.4: A schematic of a segmented tube.
4.1.4 Extruder-Rheometer
In a capillary rheometer, rheological characterization of polymer melts is performed off-
line and separate to the processing stage. It is difficult to accurately measure the pressure
in the capillary: the pressure is usually measured in a reservoir with a large diameter
preceding the capillary. In order to overcome the above mentioned problerns, it was
decided to use a single screw extruder as an extruder-rheometer to measure the viscosity
of polymer melts. The extruder is operated analogously to a capillary rheometer by
extmding into a capillary shaped die of extremely long land length.
Two pressure transducers (Gefran Mode1 SP8301SP5MRI) were used to monitor the
pressure drop in-line across the tube length. An inherent transducer uncertainty off 0.5%
of the transducer span was specified by the transducer manufacture. Pressure drop values
were used to generate flow curves.
Each pressure transducer was mounted into one of the tube pieces which is shown in
Figure 4.5. The position of one of the pressure transducers was fixed while the other can
be moved along the length of the tube by rearranging the tube segments.
Accurate measurement of the pressure drop across the die length is critical to the
successful obtaining of rheological properties of polymers [ 142,1431. Tadmor [144] has
classified pressure fluctuations occumng in an extruder by their Frequency of occurrence.
Sources of fluctuation in order of decreasing frequency are: screw rotation, solid breakup,
cycling of temperature controllers and irregular feeding to the extruder.
Placing the transducer within the capillary die away from the screw tip minimized
pressure fluctuations due to screw rotation. The use of PID temperature control
eliminated any temperature induced pressure fluctuations. Maintaining a constant level of
resin in the feed hopper prevented irregular feeding. Figure 4.6 shows the pressure versus
time for HDPEl and PP3. Mierent transducer error, variability in raw polymer matenal,
and solid bed breakup are likely causes of the fluctuations in pressure readings. The mean
pressure was used over the steady-state range of values to minimize errors caused by
pressure fluctuations.
-<- Transducer mount
< 1 Ocrn
Figure 4.5: A schematic of the segrnented tube for mounting the pressure transducer.
End effects
End effects are typically significant in capillary rheometer experiments due to large
entrance effects arising fiom converging melt flow fiom the barre1 to the die of the
rheometer. The Bagley end correction is used to account for end effects in the capillary
rheometer. End corrections were determined in the extruder-rheometer for each polymer
at each shear rate by performing linear regression on pressure vs. L/D data. Bagley plots
and 95% confidence intervals were plotted on the same axes for each shear rate. An
example of such a plot is shown for HDPEl in Figure 4.7. Bagley plots for al1 polymers
can be found in Appendix III. The origin (zero end correction) falls within the 95%
confidence intervals for ail polyrners. We can conclude that the end efTects are negligible
in the extruder-rheometer and there is no need to correct wall shear stress and Equation
2.22 can be used directly. This is an advantage over the capillary rheometer.
HDPW T cmp.=260 OC
4 #-
O 50 100 150 200 260 300 360
lime (sec)
Figure 4.6: Pressure readings versus time for HDPEl and PP3 at two temperatures a) low screw speed (5 r.p.m) b) high screw speed (33 r.p.m.).
Figure 4.7: End correction determination for HDPEl a) different flow rate (Q=IO" m3/sec) b) upper and lower 95% confidence intervals for Q=O. 15x10~ (m3/sec) [145].
Flow curves
The extruder-rheometer was used to measure melt viscosity of HDPE and PP. To
calculate the melt viscosity, it is necessary to know wall shear stress and wall shear rate.
The wall shear stress can be determined from the pressure drop values across the tube,
while the wall shear rate can be determined fiom the tube radius and the volumetric flow
rate. The shear rate range (3sec" to 50sec-') investigated was restncted by the torque
limitations on the extruder drive, and the maximum screw speed attainable by the motor
(90rpm). The polymers exhibited power law fluid behavior as evidenced by the linearity
of the log-log plots of wall shear stress zgainst apparent shear rate. Table 4.2 shows n, k
values along with standard deviation (S.D) at 220°C by using the extruder-rheorneter
[I45],
To evaluate the accuracy of the extruder-rheometer measurement, a capillary rheometer
was used to measure the viscosity of the polymers. The maximum shear rate attainable in
the capillary rheometer is much higher than in the extruder-rheometer experiment. The
only data collected from the capillary rheometer at low shear rate were used to compare
the two rheometers. Figure 4.8 shows the viscosity determined from the capillary
rheometer along with those determined from the extruder-rheometer for PP3. There is
good agreement between them.
Table 4.2: n, k values fiom extruder-rheometer, ~ e r n ~ . = 2 2 0 ~ ~ .
O 10 20 30 40 50 60
Shear Rate (sec'')
Polymet
HDPE 1
PP 1
PP3
Figure 4.8: Cornparison of capillary rheometer and extruder-rheometer viscosity measurements for PP3 at Temp.=230°C.
l 1 I I I I
n
0.666
0.608
0.703
S.D.
0.008
0.03 1
0.023
k
MPa.(sec)"
0.0098
0.007 I
0.0017
S.D.
O. O003
0.0003
0.0002
4.1.5 Summary
This section showed the development of a process based on the flow of the extruder
output through a long segmented tube to allow the study of shear-induced migration and
phase morphology development during melt blending of immiscible polymer blends.
The extruder provides the input for the tube. The shear-induced migration took place in
the extruder because of non-homogenous flow when there was viscosity difference
between polymer components at high shear rate. There was more of the dispersed phase
(PA-6), the high viscosity component, at shell region. The shear-induced migration was
quantified by using DSC. Morphological studies also showed the average size of the
dispersed phase was larger at the shell region. However, as polymer entered to a
transition piece afier the extruder the average size of the dispersed phase (PA-6)
decreased at the shell region and increased at the core because the high and low shear rate
regions in the transition piece are switched in location compared to the extruder barre1
situation. A static mixer was installed afler the extruder to obtain a uniform dispersed
morphology before entering the tube. This allows the effect of the tube on shear-induced
migration and phase morphology development to be examined separately from the
extruder.
A single screw extruder was used as an extruder-rheometer to measure polymer melt
viscosity by extruding into a capillary shaped die of extremely long land length. Pressure
lost along the capillary, volumetric flow rate and capillary dimensions were measured.
Bagley plots (end effects) were constructed for al1 polymers and showed that no end
correction was necessary for the extruder-rheometer data. This is an advantage over the
capillary rheometer. Power law fluid behavior was found for polymers and n, k values
were determined. In order to verie the accuracy of the extruder-rheorneter viscosity
measurements, a capillary rheometer was used for comparative viscosity measurements.
Good agreement between the capillary rheometer and extruder-rheometer was obtained.
4.2 Analytical methods development
The main objective of this section is to develop analytical methods to monitor de-mixing
in immiscible polymer blends by measuring polymer composition. Two methods were
used in this work to obtain polymer composition: Fourier Transform Infrared
spectroscopy (FTIR) and Differential Scanning Calorimetry @SC). These two methods
are each described in turn in the following sections under three headings: method
description, fit of the calibration curve and prediction capabilities of the method.
4.2.1 Fourier Transform Infrared Spectroscopy (FTIR)
Method description
Two characteristics of IR spectra are particularly important: (a) position of the
characteristic peak (indicating the type of chemical group) and (b) intensity of the peak
(indicating the concentration of the chemical group).
Many of the more important data processing algorithms used in FTIR spectrometry rely
on the intensities of spectral bands being linearly proportional to the concentration of
each component in the sample. For absorption spectrometry, the law relating band
intensity to the concentration is Beer's law.
Calibration curves for individual polymer components were obtained from analysis of the
extruded polymer blends of different compositions while assuming that the presence of
one polymer did not affect the absorbance of the other.
Figure 4.9 illustrates the raw FTIR spectra of HDPE/PP blends in the range of 500 cm"
to 1400 cm-'.
The criteria for selecting the characteristic peaks were that the peaks must be distinct and
non- overlapping. Table 4.3 shows the chosen characteristic peaks for PP and HDPE.
The absorption bands at characteristic peaks of PP in Table 4.3 are a result of methyl
(CH3) group vibrational modes [146]. Strong methyl rocking vibrations take place at
97 1.3 cm" and a methyl wagging mode take place at 1 166.7 cm-'. Absorption band at
7 19.32 cm" and 730.89 cm" are a result of CH2 rocking.
The plots of the net absorbance (the height of the peak) at the above wave numbers
(Table 4.3) as a fùnction of the concentration are shown in Figure 4.10. The figure shows
the average of three replicates.
The figure indicates that at low level concentration of either PP or HDPE (< SON), there
is not only a linear relationship between the net absorbance and the concentration but also
a significant change in the net absorbance with concentration at the al1 wave numbers.
However, there is only a very small change in the net absorbance with the concentration
at high concentration levels (>50%) especially in the range of 50% to 80%.
The change in net absorbance with increasing concentration must be suficient to measure
the concentration of the blend precisely. For best sensitivity, it was decided to make two
calibration curves, one using the characteristic peak of PP for low concentration of PP
(<50%) and the other using the characteristic peak of HDPE for low concentration of
HDPE (GO%).
Fit of the Calibration Curve
Linear regression was used to obtain calibration curves. For this type of application the
main consideration was the "goodness of fit" of the equation to the data and the error in
the predicted values obtained when the equation is used. The "goodness of fit" measures
used were: the multiple correlation coefficient squared, the correlation coefficient, the
standard error of calibration, and the plot of residuals (Appendix II). The prediction
ability of the calibration curves was justified using the 95% confidence interval of the
predicted values about the fitted line and the standard error of prediction (Appendix II).
Correlation analysis was used to find the best wave number for simple linear regression
(SLR) analysis. The correlation coefficients for the absorbance values and concentration
of PP (< 50 wt%) and HDPE (< 50 wt%) are summarized in Table 4.4 and Table 4.5.
Regarding Table 4.4 and Table 4.5 the net absorbances at 1 166.7 cm-' and 7 19.32 cm"
have the highest correlation coefficient with the concentration of PP and HDPE
respectively. Therefore, they were chosen to obtain the calibration curves using Simple
Linear Regression (SLR).
Fi y r e 4.1 1 shows the calibration curves and Table 4.6 shows the rneasure of goodness of
fit. The figure shows a linear relationship between the net absorbance at the characteristic
peaks (1 166.7 cm-', 719.32 cm") and the concentration (PP, HDPE). The resulting
equations are as follows:
where, C is the concentration and A is the net absorbance.
Table 4.6 indicates that in the case of Equation 4.1 and 4.2 R ~ , the multiple correlation
coefficient squared, values are 0.984 and 0.990 respectively which implies that 98.4%
and 99% of total variation in Y can be explained by fitted equations.
RMSEC is the standard deviation of the residuals due to differences between the actual
values and the predicted values for samples within the range of calibration data. RMSEC,
shows a good fit if its value is low with respect to composition.
600 600 700 800 900 1000 f i00 1200 1300 1400
Wavenurnber (cm")
Figure 4.9: F"TlR spectra of HDPElPP blends: number of scans 60, resolution 4cm".
Table 4.3: List of the characteristic peaks selected for HDPE and PP.
Polymer Characteristic Peaks
(cm-')
HDPE
HDPE (Wh)
O. 9
0.8 -
0.7 -
y 0.6 -
5 0.5 - 5 2 0.4 - Ci Q)
0.3 -
0+2
Figure 4.10: Net absorbance at characteristic peaks versus composition a) ch;rractenstic peaks of HDPE (71 9.32cm-', 730.89cm-') b) characteristic peaks of PP (97 1.3cm-', 1 166.7cm-').
I o. 1
o. 0 1 1 1 1 I 1 1 I
10 20 3 0 40 50 60 70 80 90 100
* 719.32 (cm")
0 . . a * * a i )
r n . . =
rn 730.89 (cm-') 0
4 . e m m rn
0.6 -
0.5 -
3 5 0.4 -
3 s 0.3 - Y Q Z
0.2 -
0.1- A
II
a
971.3 (cm-') m
m B
I . m
A A A
A A * A A A C m
A 1 166.7 (cm-') 0 m . A . A A
A A
0.0 T I I 1 I I 1 1 I t
10 20 30 40 50 60 70 80 90 100
PP (wt'h)
Table 4.4: Correlation coefficients for the characteristic peaks of PP.
PP
(wt O/@)
Abs. at
1166.7
(cm-')
Abs. at 1166.7
(cm-' )
Abs. at 971.3
(c mm')
Abs. at
971.3
(cm-')
0.992
0.975
Table 4.5: Correlation coefficients for the characteristic peaks of HDPE.
HDPE
(wtO/o)
HDPE (wt%) l Abs. at 1 Abs. at
Abs. at 719.32
(cm-') - Abs. at 730.89
(cm-' )
719.32
(cm*')
0.995
0.986
730.89
(cm-')
Figure 4.1 1 : Cahbration curves for HDPERP blends using SLR a) characteristic peak of HDPE 7 19.32crn-' b) characteristic peak of PP 1 166.7cm-' .
The RMSEC is low (<1.5 wt%) for both models, indicating that the chosen models fit the
calibration data very well.
Table 4.6: The measurement of goodness of fit of calibration curves for HDPElPP blends
The plot of the residuals for the calibration samples and prediction samples is show in
Figure 4.12. The magnitude of the scatter in the data and the presence of any non-random
trends can be seen in the plot of residuals. No trend is evident in this figure and the
magnitude of the residuals ranges fiom -1.5 wt% to 2 wt%.
using FTIR method.
Prediction Capabilities
Model
SLR
SLR
RMSEP indicates the standard deviation for the residuais due to differences between the
actual value and the predicted value for samples outside of the calibration set using a
specific calibration equation. RMSEP and the plot of 95% confidence intervals were used
to check the prediction capability of the FTIR method. Figure 4.13 shows a plot of the
95% confidence intervals calculated using Equation (11-8). The vertical distance between
the two curved lines above and below the regression line is the 95% interval for the
predicted mean value of composition at each value of absorbance. The interval is a
minimum at the mean value of the absorbance data used. The value of RMSEP for
Equation 4.1 and Equation 4.2 are 1.025 and 1 .O98 wt% respectively which is low for
both calibration curves.
RMSEC
(wt%)
1.449
1.154
Characteristic
Peak
(cm-')
1166.7
7 19.32
R~
0.984
0.990
~ a l l b r o t i o n Samplrs
A Validetlon Srmplrs
O 5 10 15 20 2 5 30 3 5 40 4 5 5 0
Estimated (wtX PP)
A Validatlon Samplrs
1 O 15 20 25 30 35 4 0 4 5 50 Estimated (wt% HDPE)
(b)
Figure 4.12: Residual plots of calibrabon samples and prediction samples fiom SLR models a) characteristic peak of PP ( 1 166.7 cm-') b) characteristic peak of HDPE (7 19.32 cm").
Figure 4.13: Plot of the fit with 95% confidence intervals for HDPEPP blends using FTIR a) net absorbance at chancteristic peak of PP (1 166.7 cm") b) net absorbance at characteristic peak of HDPE (719.32 c d ) .
4.2.2 Differential Scanning Calorimetry @SC)
Method description
The DSC method is mostly used in polymer systems as a quantitative tool to measure
heat of transit ion and reaction, activation energies, and rates of crystal lization, reaction,
and melting. Here it was decided to attempt to use DSC to quantify polymer composition
for immiscible polymer blends. The main assumptions were:
1) The area under the melting curve in a DSC spectrum per unit mass (&) is
proportional to the heat of fusion (AHf) according to the following relationship:
where K is the ce11 constant and q is heating rate.
Heating rate influences shape and position of the peaks. The heating rate was constant
(lO°Clmin) for al1 the runs. The value of AHG determined by DSC, depends on the
amount of crystallinity present in the sample. For imrniscible blends, the melting point is
independent of the composition but the heat of fusion is proportional to the composition.
2) There is no overlap between the DSC melting peak of the pure components. Figure
4.14 shows the DSC spectrum of the pure polymers. It can be seen that each polymer has
a distinct peak and the melting points of HDPE, PP and PA-6 are 132.2OC, 164.8OC and
222. 1°C respectively. There is no overlap between the peaks for each combination of two
polymers.
3) For an immiscible polymer blend, it can be assumed that (in DSC terms) one polymer
doesn't recognize the existence of the other and vice versa. The spectrum of the blend
would reflect the simple addition of the individual components. Therefore, the area under
the melting peak, for the immiscible blends, is a function of polymer composition.
The area under the melting peak, which is a finction of degree of crystallinity, is used to
measure the polymer composition. The degree of crystallinity can be affected by the
thermal history of the sample and the polymer grade. The first step in developing a DSC
method to measure polymer blend composition was to study the effects of these two
factors on the DSC spectrum.
Thermal history
As mentioned above, the degree of crystallinity and the area under the melting peak are
affected by the thermal history of the sample. Figure 4.15 shows the DSC spectrum for
two samples of 6 O h O HDPE/PP blend with different thermal histories. One is an
extrudate sample, mechanically melt mixed, directly obtained from the twin screw
extruder and the other is the dry blended sample, not melt mixed, directly obtained from
the pure polymer pellets. The DSC measurement was camed out in two cycles: the
sample as received (first cycle) and after being cooled slowly in the DSC ce11 to destroy
the thermal history (second cycle).
The values of the area under the melting peak are also shown in Figure 4.15. The area
was calculated manually identieing two points that define the base line at the selected
range of temperatures.
Fi y r e 4.15 shows that there is a significant difference in the area under the melting peak
of HDPE and PP between the extrudate sample and the pellet sample in the first cycle. In
the second cycle, afier destroying the thermal history, the area under the melting peak is
almost the same for both samples. This result can be used as one of the advantages of
using DSC to measure polymer composition in terms of calibration samples preparation.
There is then no need to use a twin screw extruder or any other melt mixing tools to
prepare the calibration samples. We can make the calibration samples directly from the
60 80 100 120 140 160 180 200 220 240 260
Temperature ( O C )
Figure 4.14: DSC spectra of pure polymers O P E , PP and PA-6): heating rate= 10 OC/min.
5 - Fint Cycle Extnidatc
Temperature ( O C )
Figure 4.15: Effect of thermal history on the DSC spectra of 6OMO HDPEIPP: heating rate= 1 O°C/rnin.
pure polymer pellets so long as the second cycle of the DSC spectrurn is used to measure
the area under the melting peak.
Polymer grade
As mentioned before, different grades of HDPE and PP were used in this work. Each
grade of polymer has a different degree of crystallinity that influences the area under the
melting peak. Figure 4.16 shows the DSC spectrum for two grades of HDPE. The effect
of PP grade on the DSC spectnim is shown in Figure 4.17.
The figures indicate that the intensity of the DSC spectra is quite diferent for each grade
of polymer, regardless of whether or not its thermal history is destroyed. HDPE1 has a
higher melting point than HDPEZ because of its higher molecular weight and the effect of
molecular weight is show in Figure 4.16. The melting points of HDPEl and HDPE2 are
132.2OC and 130.1°C respectively, so the DSC spectrum of HDPEl shifts to the higher
temperature.
The grade of polymer has a strong effect on the DSC calibration curves. This result
means that a separate calibration curve is required for each pair of polymers. This may be
considered as one of disadvantages of the DSC method to quantify polymer composition.
Table 4.7 shows the DSC experiments which were done to obtain the calibration curves.
In the following section, the DSC results of HDPEl/PP3, HDPE 1PA-6 and PP l/PA-6
will be used to illustrate the interpretation method. DSC results for the other blend
compositions are shown in Appendix IV.
Fi y r e 4.18 to Figure 4.20 show the DSC spectra for the three different immiscible
polymer blends.
5
Fimi Cycle Second Cycle
6 0 80 1 O0 1 20 140 160 160 200
Temperature (OC)
Figure 4.16: Effect of the grade of HDPE on the DSC spectrum: heating rate= 1 O°C/min.
60 80 1 O0 120 140 160 180 200
Tem perature foc)
Figure 4.1 7: Effect of the grade of PP on the DSC spectrum: heating rate= 1 O°C/min.
Table 4.7: DSC experimental mns.
System
HDPEffP blends
HDPE/PA-6 blends
60 80 100 120 140 160 180 200
Temperature ( O C )
Combination
HDPE 1/PP2, HDPE 1PP3
HDPE 1 /PA-6, HDPE2/PA-6
PP/PA-6 blends
Figure 4.18: DSC spectra of HDPE/PP blends: heating rate=lO°C/min.
PP 1 /PA-6, PP2lPA-6
Increasing % HDPE
60 80 1 00 120 1 40 160 1 80 2 O0 220 240
Tarn pcrature ( O C )
Figure 4.19: DSC spectra of HDPEPA-6 blends: heating rate= 1 O°C/min.
60 80 100 120 1 40 1 60 1 BO 200 220 240
Temperature (OC)
Figure 4.20: DSC spectra of PP/PA-6 blends: heating rate= 10°C/min.
Figure 4.21 shows the area under the melting point of each component as a function of
composition. The figure shows that the area under the either PP melting peak or HDPE
melting peak is proportional to the concentration in the blends regardless of the blend
system. This illustrates that there is a linear relationship between the area and the
concentration at the whole range of the concentration in the blends. However, the
reiationship between the area under the melting peak of PA-6 and the concentration
appeared irregular and therefore more uncertain. Thus, the area under the melting peak of
either HDPE or PP, depending on the blend system, were used to determine composition.
(Figure 4.21a)
Figure 4.21: Area under melting peak as a fùnction of composition for polymer blends a) HDPERP blends b) HDPERA-6 blends c) f PPA-6 blends.
Fit of calibration curves
SLR and MLR (with two independent variables) were used to detenine the calibration
curves for the blends. It should be mentioned that MLR was used only for HDPEPP
blends. The equations for the calibration curves are as follows:
IEIDPEPP blends
HDPERA-6 blends
PPPA-6 blends
where, C and A are the concentration and the area under the melting peak respectively.
Table 4.8 shows the measures of goodness of fit of the calibration curves. The SLR
model shows multiple correlation coefficient squared (R2) values ranging fiom 99.4% to
99.9% for the blends. Thus large percentage of total variation in composition can be
explained by the fined equations. Regarding the calculated R', goodness of fit in the
MLR model is better than the SLR model for HDPEPP blends.
RMSEC is low in respect to the composition for three different blends, indicating that the
SLR models fit the calibration samples well. The value of the RMSEC for the MLR is
lower than the SLR model.
Figure 4.22 to Figure 4.24 show the plots of residuals for three different blend systems.
The residuals of both calibration and prediction are randomly distributed along the zero
l i ne.
Prediction capabilities
Figure 4.25 to Figure 4.27 show a plot of the 95% confidence intervals calculated using
Equation (11-8). A set of four samples, randomly chosen from the original samples, was
used as validation samples to check prediction ability of the calibration curves using
RMSEP. Table 4.9 shows the values of RMSEP for three different blend systems. The
values of RMSEP show that the prediction ability of SLR is quite good for al1 of the
blends. The prediction error of MLR is superior to that of SLR for HDPE/PP blends
Table 4.9: Prediction error of calibration curves using DSC method.
System
EDPE/PP
Mode1
1 blends
RMSEP
Eq.4.4 SLR
1
(wt%)
1.102
Eq.4.5 SLR 1
Eq.4.6 MLR 1 0.941
1 HDPEIPA-6
1 -299
blends
PP/PA-6
Eq.4.7 SLR 2.062
Eq.4.8 SLR 1.447
Table 4.8: The measurement of goodness of fit of calibration curves using DSC method.
System Model
HDPEIPP 1 Eq.4.4 SLR 1 0.997 1 1.64 I blends 1 Eq.4.5 SLR 1 0.994 1 2.271 1
O 10 20 30 40 50 60 70 80 90 100 110
Estirnated (wt% PP)
HDPEIPA-6
blends
PPtPA-6
blends
Figure 4.22a
Eq.4.6 MLR
Eq.4.7 SLR
Eq.4.8 SLR
0.999
0.998
0.996
0.9 10
1.406
I .529
+CaHbration Samplrs
Valldatlon Samplrs
O I O 20 30 40 50 60 70 BO 90 100 110
Estimated (wtXHDPE)
1 m Valldatlon barnplis l
O 10 20 30 40 50 60 70 80 90 100 110
E stimatcd (wt%HDPE)
(cl
Figure 4.22: Plots of residual for HDPEPP blends using DSC: a) SLR model ( a m under melting peak of PP b) SLR model (am under melting peak of HDPE) c) MLR model ( areas under melting peak of PP and KûPE).
O 10 20 30 4 0 50 60 70 80 90 100 t10
Estimated (wt% HDPE)
Figure 4.23: Residual plot of HDPE/PA-6 blends fiom SLR model.
6 Calibntion Sarnples 6 +
E Validation Sarnples ,il , , , ,
Figure 4.24: Residual plot of PPPA-6 blends fiom SLR model.
Figurc 4.25: Plot of the fit with 95% confidence intervais for HDPE/PP blends usbg DSC a) a r a under the melting peak of PP b) area under the melting peak of HDPE.
Figure 4.26: Plot of the fit with 95% confidence intervals for HDPE/PA-6 blends using DSC.
Figure 4.27: Plot of the fit with 95% confidence intervals for PPPA-6 blends using DSC.
4.2.3 Summary
This section showed the development of off-line analytical methods to measure polymer
composition in an immiscible polymer blend. Two methods were used: Fourier
Transform Infrared Spectroswpy (FTIR) and Differential Scanning Calorimetry @SC).
In FTIR spectroscopy, the net absorbante at 1166.7 cm-' (characteristic peak of PP) and
71 9.32 cm" (characteristic peak of HDPE) were correlated with polymer composition. In
DSC, the area under the melting peak was correlated with polymer composition.
Preparation of the calibration samples in the DSC method is much easier than in the FTIR
method. In the DSC method, there is no need to prepare the calibration samples, by using
a twin screw extruder, so long as the thermal history of the sample is destroyed and the
pellet samples, directly prepared from the pure polymer pellet, can be used for the
calibration samples. However, in the FTIR method, the preparation of the calibration
samples involved melt mixing using a twin screw extruder, and making film using
compression molding.
One calibration curve must be made for each pair of different polymers in the DSC
method, because the area under the melting peak is a function of the degree of the
crystallinity of the polymer sample and each grade of the polymer has a different
tendency to crystallize.
Deviations fiom Beer's law were observed at high concentrations of either PP or HDPE
(>50 wt%) in FTIR spectroscopy. Therefore, two calibration curves, using SLR, were
made for HDPE/PP blends. One used the characteristic peak of PP (1 166.7 cm-') for low
concentration of PP (550 wt%) and the other used the characteristic peak of HDPE
(719.32 cm") for low concentration of HDPE (a50 wt%). There was a linear relationship
between the area under the rnelting peak and the composition in the DSC method over the
whole range of the composition.
Finally, well-known statistical measures such as, R*, RMSEC, RMSEP, plots of residuals
and plots of 95% confidence intervals were used to evaluate the calibration curves. SLR
models fit the calibration data ffom FTIR and DSC very well. MLR was used only for
DSC analysis of HDPEPP blends.
4.3 Shear-induced migration and phase morphology
The main objective of this section is to study shear-induced migration and phase
morphology development during the flow of immiscible polymer blends in a long tube
for three different blends. This was accomplished by using the segmented tube discussed
in section 4.1.3 and analyzing the polymer in each segment (See section 3.3).
Concentrations of each polymer present as well as the resulting structures were of
interest.
Figure 4.28 to Figure 4.30 show shear viscosity of homopolymers as a function of shear
rate for three different blend systems. In Systeml (HDPEIPP blend), HDPEl exhibits a
higher viscosity than three grades of polypropylene over the whole shear rate range and
PP1 has higher viscosity than PP2 and PP3. At higher shear rates the viscosities of PP2
and PP3 are close to each other. However, PP3 has higher viscosity than PP2 in the range
of the processing conditions in this work (shear rate400 sec") at al1 temperatures.
In System2 (PP/PA-6 blends), the viscosity of the minor phase (PA-6) is higher than PP2
and PP3 over the whole shear rate range. In System3 (HûPE/PA-6 blends), at low shear
rates (<200 sec"), HDPEl has a higher viscosity than PA-6 and HDPE2 has a lower
viscosity than PA-6 at whole range of shear rate. PA-6 melt viscosity appears almost
independent of shear rate at low shear rates which means that melt viscosity behavior of
PA-6 shows Newtonian behavior at low shear rates.
Ur
a .- II)
8 P L n Y r V)
1 O 1 0 0 1000
Shear Rate (sec")
(Figure 4.28a)
I 1 - - -, 1 O 100 1000 1OOOO
Shear Rate (sec-')
Shear Rate (sec")
Figure 4.28: Shear viscosity as a function of shear rate of pure HDPE and pure PP for Systeml at different temperatures: (a) 195°C (b) 210°C (c) 230°C.
S hear Rate (sec-')
(a)
100 1000
Shcar Rate (sec-')
(b)
Figure 1.29: Shear viscosity as a fiinction of shear rate of pure PP and pure P A 4 for System2 at difFerent ternpcratures: (a) 230°C @) 260°C.
100 1000
S hcar Rate (sec")
(a)
10 100 1000 10000
Shear Rate (sec")
(b)
Figure 4.30: Shear viscosity as a function of shear rate of pure HDPE and pure PA-6 for Systern3 at differeni temperatures: (a) 230°C (b) 260°C.
4.3.1 Shear-induced migration
Table 4.10 to Table 4.14 show composition analysis along with standard deviation (S.D)
in the core region and shell region for three different blend systems. Apparent shear rate
was calculated using Equation 2.25 and viscosity ratio (X) which is the ratio of the
viscosity of the dispersed phase to the matnx phase, was estimated fiom Figure 4.28 to
Figure 4.30.
It can be seen that there is no significant difference in the composition of the core region
and the shell region at al1 shear rates when the viscosity ratio is close to one. There is also
no significant composition difference between the core and the shell at low shear rates in
spite of large viscosity differences between the components of the blend. However, there
is a significant composition difference between the core and the shell at high shear rates
when the viscosity ratio is far fiom one; there is a greater fraction of the high viscosity
component in the core region although the separation is low.
Table 4.15 shows Ei/Eir fiom Equation 2.19 for the blend systems used in this work.
There is no composition variation in the core and the shell when Er/ErI is close to one.
The values of Ei/EII for HDPEIPP 1, HDPEZPA-6 and PA-6PP1 are 0.8, 0.85 and 0.95
respectiveiy and composition measurements show no indication of shear-induced
migration for these blend systems. However, there is composition variation in core and
shell when Er/EiI is significantly less than one. The values of Er/& for HDPE I/PPZ, PA-
6MDPE2 and PA-6/PP2 are 0.37, 0.32 and 0.42 respectively and for these blend systems
polymer de-mixing was observed. Composition measurements showed that there was
more low viscosity cornponent in the shell than in the core. We can therefore conclude
that the shear-induced migration results are in accordance with the principle of energy
minimization.
Table 4.10: Composition analysis of 30/70 PPIHDPE blends using FTIR at T=195 OC and L/D=137.
Viscosity Ratio 1
~ P P ~ ~ H D P E
Composition in Core k S.D.
(wtO/o)
Feed Composition
(wt%)
Apparent Shear Rate
(sec-')
Composition in SheN 3. S.D
(wtOh)
PPl 30 HDPEl 70
PP 29.5 k 1.07 HDPE 70.5 f 1.07
Pf 28.4 f 0.92 HDPE 71.6f 0.92
PP 31f1.13 HDPE 69f 1.13
PP 29 I 1.03 HDPE 7 1 * 1.03
PP 1 30 HDPEl 70
PP2 30 HDPEl 70
3.1
PP2 30 HDPEl 70
PP 26 f 0.72 HDPE 74 10.72
0.11
PP 36f 0.88 HDPE 64 10.88
Table 4.11: Composition analysis of 30170 PA-6/HDPE2 blends using DSC at T=230 OC
PP 28.3 f 0.93 HDPE 71.7 f 0.93
and L/D= 144.
PP 30.3 f 0.74 HDPE 69.7 10.74
Feed Composition
(wtOh)
Apparent Shear Rate
(sec-')
Viscosity Ratio h
q~~dtl HDPE
Composition in Core S.D.
(wtOh)
Composition in Shell k S.D
(wt%)
PA4 28.8 f 1.10 HDPE 71.2 f 1.10
PA4 27.4 f 1.23 HDPE 72.6 f 1.23
PA4 30.8 f 1.50 HDPE 69.2 f 1.50
PA-6 28.2 I 1.3 1 HDPE 71.8f 1.31
PA4 35.9 f 0.68 HDPE 64.1 1 0.68
PA4 26.8 1t 1.30 HDPE 73.2 f 1.30
Table 4.12: Composition analysis of 30170 PA-6/PP2 blends using DSC at T=230 OC and L/D= 131.
Viscosity Ratio h
Feed Composition
(wt0/0)
Composition in Core + S.D.
(WtYD)
Apparent Shear Rate
(sec")
Composition in Shell k S.D
(Wt%)
Table 4.13: Composition analysis of 3O/ ïO PA-6/HDPE1 blends using DSC at T=230 O C
and L/û=144.
1 Apparent Shear Composition in Shell k S.D
(wt%)
Feed Viscosity Ratio Composition in Core 4 S.D.
(wt%) Composition
(wtO/o) Rate
(sec-')
PA-6 30 HDPEl 70
P A 6 30.9 f 1.14 HDPE 69.1 1t 1.14
P A 6 3 1.2 rt 1.32 HDPE 68.8 f 1.32
PA-6 30 HDPEl 70
P A 4 29.6 I0.88 HDPE 70.4 zk 0.88
P A 4 31.1 I 1 . 2 7 HDPE 68.9 I 1.27
PA-6 30 HDPEl 70
P A 4 31.3 I 1.85 HDPE 68.7 1 1.85
Table 4.14: Composition analysis of 3ORO PA-6/PPl blends using DSC at T=230 OC and
Feed Apparent Shear Composition Rate
(wt%) (sec-')
Viscosity Ratio Composition in h Core + S.D.
~ P A ~ ~ P P (wt%)
Table 4.15: Ei/Eir for the blend systems used in this work, 41=0.7.
Blend System Temperature
HDPE IPP2
Composition in Shell k S.D
(wtO/o)
A radial, cross-sectional, concentration profile was measured across 8mm diameter of
30/70 PPZRIDPEl blend extrudate. The purpose was to determine the concentration
profile of the PP2 across the cross-section of the sample. Figure 4.3 1 shows the
composition of PP2 as a function of radius. The figure is scaled fiom O mm to 4 mm,
corresponding to the radial ends of the cross-section with O mm at the center. Error bars
correspond to the standard deviation of each point. The figure shows that PP2, the low
viscosity dispersed phase, migrates to the wall of the tube. It also shows that the change
in the concentration of PP2 with radius is higher close to the wall of the tube than at the
center of the tube.
O 0.5 1 1 .S 2 2.5 3 3.6 4
Riidius (mm)
35 -
30 - .
25 + n
f 20
Figure 4.31: Radial composition profile of PP in 30/70 PPZ/HDPEI blend across the cross-section of the sample; T= 195"C, fa,=22 sec", L/D=137.
b a
1s -
10 -
6 -
+- Radial Composition
- Feed Composition
It was s h o w that al1 of the polymers used in Systeml exhibited power law fluid behavior
as evidenced by the linear log-log plots of wall shear stress versus apparent shear rate
(Section 4.1). Therefore, the power law behavior was assumed in order to calculate the
shear rate distribution in the tube.
The shear rate distnbution in the tube for 30/70 PPZIHDPEI blend at the experimental
conditions is shown in Figure 4.32. The power law's constants (n, k) of the matrix
(HDPEI) were used to calculate the shear rate distribution. It can be seen that the shear
rate as a fùnction of radius is not linear because of non--Newtonian behavior. The change
in shear rate with the radius close to the center of the tube is much lower than close to the
wall of the tube. As a result, the shear-induced migration close to the center of the tube is
lower than close to the wall of tube. This is in agreement with the concentration gradient
of PP2 being higher near the wall.
O 0.6 1 1.6 2 2 6 3 3.6 4
Radius (mm)
Figure 4.32: Shear rate distnbution in the tube as a function of radius for 30170
PP2/HDPEl: T=195OC.
In order to understand and explain shear-induced migration during melt blending in the
tube, it is important to study de-mixing along the length of the tube. For this purpose a set
of experiments was carried out for PPZ/HDPEI blends. Figure 4.33 shows the
composition of PP2, the low viscosity component, dong the length of the tube in the core
region and the shell region for 25/75 PP2/HDPEl blend. Error bars correspond to the
standard deviation of each point. The figure suggests several important points. Firstly, it
indicates that shear-induced migration increases along the length of the tube. This may be
attributed to the higher residence time in the long tube that allowing significant shear-
induced migration. Secondly, the average composition separation per unit length of tube
(one cm.) is only 0.11 wt%. This emphasizes that the shear-induced migration is a very
slow process during melt blending of the immiscible polymer blends in the tube under the
condit ions used.
- Fetd -c core * Shdl
Figure 4.33: Showing, the effect of the length of the tube on de-mixing for 25/75 1 PPZ/HDPEI blends; T=2 10°C, &=32 sec- .
4.3.2 Phase morphology development
Morphological studies of immiscible polymer blends in a capillary die have been reported
in the literature [20, 43, 1471 and the extrudate samples afier the die mostly were used to
evaluate morphology development. However, in order to understand and explain final
morphology of an extmded product it is important to study morphology development
along the length of the die. The process design that pem~itted sampling along the length
of the tube along with the staining technique in this work allowed study of the
morphology development along the tube. In addition, using a static mixer allowed the
evaluation of the effect of the tube on the phase morphology developrnent separately
fiom the extruder. The staining technique can be used for System2 and System3 then
these systems were used to study phase rnorphology development in the tube.
System2 (PP/PA-6 blends)
Morphology development before entering the tube
Figure 4.34 presents rnorphology of 3OIîO PA-6BPl perpendicular to the direction of the
flow at 0.5 cm downstream from the static mixer. Figure 4.34 shows the morphology of
the melt entering the tube. The two phases can be distinguished in the blend. The white
color represents the dispersed phase (PA-6) and black color represents the matrix (PPI).
It shows qualitatively that the dispersed phase is distributed in the matrix uniformly.
Figure 4.34 also indicates a coarse morphology for this blend. One reason for this coarse
morphology is a large intefiacial tension of the PPPA-6 blends [20]. It can be concluded
that a dispersed phase structure is entering the tube and there is no indication of the
fibrillation of the dispersed phase (PA-6) in both regions at the experimental conditions.
Wall
Core (r= 1.3 mm)
Figure 4.34: PA-6PP1 30/70 blend morphology development at 0.5 cm afier static mixer perpendicular to the flow direction, fap=30 sec", Temp.=230°C A: shell B: core.
Morphology development in the tube
A set of experiments with three grades of PP was conducted. The order of the viscosities
of the four polymers at the experimental conditions were PA-6>PP l>PP3>PP2
Figure 4.35 shows the morphology of 30/70 PA-6PP blends with three grades of PP at
L/D=94 in the flow direction, The viscosity ratios of the dispersed phase (PA-6) to the
matrix phase (PP) (h) are 5.89, 2.73 and 1.41 at the experimental condition. It can be seen
that at high viscosity ratio (5.89) only a dispersed phase structure develops in both
regions, but at low viscosity ratio (1.41), a multi-layer structure develops in the shell
region. [Note: in these and subsequent photographs showing non spherical dispersed
phase the morphology was lamellae not fibers.]
Figure 4.36 shows the rnorphology development along the length of the tube in the core
and the shell regions for 30/70 PA-6PPl blend. The viscosity ratio is 1.21 (h) at the
experimental conditions. Comparing al1 the locations in Figure 4.36, it can be seen that a
multi-layer structure is developed gradually along the length of the tube in the shell
region. However, the morphology development along the length of the tube in the core
region does not change drastically. One explanation for this morphology development is
that ir the shell region, deformation of the dispersed phase is facilitated because of higher
shear stresses.
Examination of the shell region (Picture A) fiom Figure 4.36(c) reveals that the size of
deformed dispersed phase, especially at higher Lm, is larger than it was before entering
the tube. This may show that the deformation is associated with the shear-induced
coalescence. This association of shear-induced deformation with shear-induced
coalescence is also reported by Utracki [43].
Fiyre 4.37 shows the morphology development as a fùnction of the length of the tube for
30/70 PA-6/PP2. The viscosity ratio for this system is about 6.04 at the experimental
conditions. It can be seen that the length of the tube has no marked effect on the
morphology development in terms of developing a multi-layer structure and there is no
indication of shear-induced deformation. This means that the viscous force cannot
overcome the interfacial force and no deformation then takes place under experimental
conditions for this blend. Figure 4.37 also reveals that the size of the dispersed phase
(PA-6) increases along the length of the tube in both regions: shear-induced coalescence
increases along the length of the tube.
B
Figure 4.35 (a)
Wall
Core
R
Wall
Core (d.8mm)
R Figurce 4.35 (b)
Core (r= 1.5mm)
Figure 4.35: Effect of viscosity ratio on the morphology of 3O/7O PA-6PP blends; y",= 39sef1, Temp.=230°C, LID=94, A: shell B: core (a): PA-6/PP2 h=5.89 @): PA- 6/PP3 h=2.73 (c): PA-6/PPl h=1.41.
Figure 4.36 (a)
Wall
Core
Core ( d . 9 m m ) B
Figure 4.36 (b)
Wall
Core (r= 1.7mm)
Figure 4.36: PA-6PP 1 30/70 blend morphology developrnent dong the length of the tube parallel to the flow direction; i,=30 sec-', Temp.=230°C , Acl.2 1 A: shell B: core a) Lm44 b) L/D=69 c) L/D=132.
Wall
Core
B
Figure 4.37 (a)
Figure 4.37 (b)
Wall
Core
R Figure 4.37: PA-6RP2 30170 blend morphology development dong the length of the tube parailel to the flow direction, y,= 9 sec-', Temp.=230°C, A4.04 A: shell B: core (a) LD=56 (b) L/D=106 (c) L/D=156.
System3 (HDPEPA-6 blends)
The results fiom Systern2 revealed that viscosity ratio plays an important key role in flow
induced morphology of PPPA-6 blends in the tube, especially in developing the multi-
layer structure along the length of the tube. At low viscosity ratio, the dispersed phase
(PA-6) deformed gradually in the shell region and also shear-induced coalescence
occurred until finally a multi-layer structure developed along the length of the tube.
However, at high viscosity ratio, there was no indication of shear-induced deformation in
both regions and only a dispersed phase stnicture was observed along the length of the
tube.
A set of experiments with 30/70 PA-6/HDPEl was done to evaluate the phase
morphology development in the tube at low values of viscosity ratio. Experimental set up
and processing conditions were the same as for System2.
Figure 4.38 gives the morphology development along the length of the tube parallel to the
flow direction for 30170 PA-6/HDPEl. The viscosity ratio of the dispersed phase (PA-6)
to the matrix phase (HDPEI) is 0.64(h<l) at the expenmental conditions. As mentioned
earlier, the dispersed phase structure entered the tube and that there was no indication of
the fibrillation of the dispersed phase (PA-6) in both regions (See Figure 4.3(a)).
However, flow in the tube tends to change the dispersed phase morphology, especially in
the shell region, as shown in Figure 4.38 a, b, and c. It can be seen that the dispersed
phase deforms gradually along the length of the tube and shear-induced coalescence
occurs among the deformed dispersed phases in the shell region. The size of the PA-6
layers especially at higher L/D (see Figure 4.38(c)) confirmed that shear-induced
coalescence had taken place during flow in the tube.
It seems that there is a difference in the concentration of the dispersed phase (PA-6)
between the core region and the shell region in Figure 4.38(c). However, composition
analysis confirmed that there is no significant composition difference at both regions (see
Table 4.13). Fellahi et. al. [41] studied morphological stability in injection molded
HDPE/PA-6 blends. They showed that there was an apparent absence of the minor phase
(HDPE) near the die wall using SEM but composition analysis revealed the presence of
both components on the surface of the blends and it was very close to the feed
composition. The apparent absence of the dispened phase as observed by SEM was
attributed to the very fine size of the dispersed phase in the skin relative to the core.
Effect of Processing Variables
Several processing variables were evaluated regarding their effect on the morphology of
HDPEIPA-6 blends: viscosity ratio, composition and tube temperature. The main
consideration in these studies was how they affect the morphology development.
Viscosity ratio
The effect of the viscosity ratio on the phase morphology development of 30/70 PA-
6/HDPE blends is show in Fi y r e 4.39. The viscosity ratios for PA-61HDPE 1 and PA-
6RIDPE2 at the experimental conditions are 0.8 and 10.5 respectively. Figure 4.39(a)
shows that there is no indication of the shear-induced deformation at the high viscosity
ratio. As a result, only dispersed phase structure is obtained in both regions. At the low
viscosity ratio (Figure 4.39(b)), shear-induced deformation and shear-induced
codescence take place during flow in the tube at the shell region and, as a result, the
multi layers structure is obtained in the shell region.
Composition
Concentrations of 10 wt%, 15 wt% and 20 wt% of PA-6 in HDPEl matrix were
investigated. The influence of the dispersed phase concentration on the phase
morphology in the flow direction is shown in Figure 4.40. It can be seen that as
concentration increases the multi-layer structure becomes more pronounced [72].
Increasing concentration of minor phase in the blend system increases the particle size.
The increased particle size enhances particle deformation and assists the formation of a
multi-Iayer structure. This increase in particle size is attributed to an increase in
coalescence [72].
Temperature
Fiyre 4.41 shows the morphology development in the tube for 30/70 PA-6/HDPEl
blends at two temperatures (230°C, 260°C). It can be seen that multilayer structure
develops in the shell region for both temperatures. However, changing the temperature
has a significant effect on the morphology development close to the core region [72]. At
higher temperature, the dispersed phase defoms more easily and this then contributes to
an increase in the shear-induced coalescence in the core region.
Wall
Core (r=O . 6 m )
Figure 4.38 (a)
Figure 4.38 (b) R
Wall
Core (r=O. 4mm)
Figure 4.38: OM showing the effect of tube length for 30 wt% PA-6 in 70 wt% HDPE 1, parallel to theflow direction, yO,= 31 sec-',Temp.=230°C, b=0.64 A: shell B: core a) WD=31 b) L/D=56 c) L/D=144.
Core
B
Figure 4.39 (a)
Wall
Figure 4.39: Effect of viscosity ratio on the morphology of 3O/7O PA-6/HDPE blends parallel to flow direction, fa,,=36sec-', Temp.=230°C, LD=125, A: shell B: core (a): PA-6RIDPE2 h=10.5 (b): PA-6MDPEl A,=0.8.
Wall (r=3.6mm)
B
Figure 4.40 (a)
B
Figure 4.40 @)
Core (d. 5mm)
F=i
Wall
Core
F=-i Figure 4.40: Effect of composition on the morphology development of PA4HDPEI blends; flow direction, y',= 30 sec", Temp.=230°C, L/D=150, A: shell B: core a)10/90 b)l5/85 c)2O/8O.
Wall
Core (~û.7mm)
B
Figure 4.41 (a)
Wall
Core (r- lmm)
Figure 4.41: Effect of temperature on morphology of 30170 PAdMDPE 1 blends parallel to flow direction, L/D=110, screw speed=60 r.p.m, A: shell B: core (a) 230°C (b): 260°C.
4.3.3 Summary
This section deals with shear-induced migration and phase morphology development
during flow of immiscible polymer blends in a long tube.
The results confirmed that shear-induced migration occurred when there were viscosity
differences between the polymer cornponents. The concentration of the low viscosity
component in the shell region was higher than in the core region. Evaluation of shear-
induced migration along the length of the tube showed that this phenornena was a very
slow process and the separation was low in the long tube. However, separation increased
with increasing viscosity ratio, apparent shear rate and tube length. The shear-induced
migration close to the wall of the tube was higher than close to the center of the tube
because of the non-linear shear rate profile resulting ffom non-Newtonian behavior.
An energy approach based on the theoretical rate of viscous energy dissipation per unit
length of the tube was taken to show that shear-induced migration results are in accord
with the principle of energy minimization.
The morphological studies reveal that a coarse dispersed phase structure iç entering the
tube and there is no indication of the fibrillation of the dispersed phase (PA-6) in the core
region and the shell region. The coarse rnorphology is attributed to high interfacial
tension of the blends. However, flow in a long tube tends to change the morphology of
the blends, depending upon viscosity ratio, shear rate, tube length and concentration of
the dispersed phase.
Viscosity ratio plays an important role in the morphological changes in terrns of
obtaining the dispersed phase structure and the multi-layer structure. Morphological
studies along the length of the tube reveal that at high viscosity ratio, the dispersed phase
stmcture develops along the length of the tube. There is no indication of developing a
multi-layer structure at the maximum attainable (shear rate)-(concentration of the
dispersed phase)- (tube length) combination used in Our experiments. However, at low
viscosity ratio (close to unity), a multi-layer structure develops along the length of the
tube to a degree dependent on the concentration of the dispersed phase (PA-6), shear rate
and the length of the tube.
The multi-layer structure is apparently a result of shear-induced deformation and shear-
induced coalescence. These shear effects are observed for the blends. The dispersed
phase deforms in the shell region and also shear-induced coalescence takes place. Then,
the multi-layer structure develops gradually along the length of tube in the shell region.
As concentration of the dispersed phase increases, there is an increase in particle size
because of coalescence effects which enhances the deformation of the dispersed phase.
Therefore, the multi layer structure is then more pronounced.
4.4 Flow behavior of immiscible polymer blends
4.4.1 Systeml (HDPERP blends)
An immiscible polymer blend consists of two phases, dispersed and matrix. The shear
rate at the interface between the two phases is discontinuous because of differences in
their viscosities [10]. However, the shear stress at the interface is continuous and
therefore the shear stress is used in correlating the viscosity-composition data.
Figure 4.42 shows the apparent viscosity of HDPElPP3 blends versus blending ratio as a
function of the wall shear stress at two temperatures. PP3 is a lower viscosity component.
The apparent viscosity decreases with increasing wall shear stress. The apparent viscosity
is a rnonotonic function of composition at both temperatures.
Figure 4.43 shows SEM photomicrographs of extrudate of ZO/8O PP3RIDPEl blends at
two temperatures: 195OC and 260°C. PP3 is the dispersed phase, which is embedded in
HDPE 1 matrix. The white areas represent the PP phase and the dark areas represent the
HDPE phase. It can be seen that the dispersed phase aligns in the flow direction in the
shell region. This is may be attributed to the higher shear rate area close to the wall of the
tube.
4.4.2 System3 (HDPE/PA-6 blends)
Figure 4.44 shows the eEect of blending ratio on melt viscosity of HDPEPA-6 blends
with two grades of HDPE at 230°C. HDPEl has a higher viscosity than HDPEZ. It c m be
seen that the viscosity-composition behavior is strongly depending on the grade of
HDPE. In the case of HDPE2RA-6 blends, the melt viscosity lies between those of the
pure homopolymers. In the case of HDPEIRA-6 blends, the melt viscosity goes through
a minimum a t approximately 50 wt% of KDPEI . These different types of behavior for
HDPE/PA-6 blends may be associated with phase morphology.
O 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1
HDPE (WC%)
O 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1
HDPE (wt%)
Figure 4.42: Melt viscosity verms composition as a function of wall shear stress for HDPEl/PP3 b1ends a) 195°C b) 260°C.
B
Figure 4.43 (a)
Figure 4.43: SEM photograph of 20180 PPîRIDPEl paralle1 to the flow direction, wall shear stress=0.08 MPa, A: shell B: core a) 195°C b) 260°C.
Figure 4.45 and Figure 4.46 shows optical microscopy (OM) photomicrograph of
HDPE/PA-6 blends parallel to the flow direction. In these pictures, the PA-6 phase
appears white and the HDPE phase appears black. A dispersed phase structure developed
for HDPEZPA-6 blends and there is no indication of shear-induced deformation and
developrnent of a multi-layer structure. However, a multi-layer structure developed for
HDPEIPA-6 blends and is more pronounced as the concentration of PA-6 increases.
This can been seen for 50/50 HDPE l/PA-6.
Figure 4.44: Melt viscosity versus composition for HDPEPA-6 blends, T=230°C.
Figure 4.45 (a)
Figure 4.45 (b)
Wall
Core
m Figure 4.45: OM photograpb of HDPE2/PA-6 blends in the flow direction, A: shell B:core, shear stress=O.O 1 MPa, Temp.=230°C a) 80/20 b) 60/40 c) 50150.
Figure 4.46 (a)
Figure 4.46 (b)
Wall
Core (r= 1.2m
Figure 4.46 (c)
Wall
Core (F 1.4mm)
Figure 4.46: OM photograph of HDPEl/PA6 blends in the flow direction, A: shell B:core, shear stress=0.06 MPa, Temp.=230°C a)80/20 b)70/30 c)60/40 d)50/50.
4.4.3 Evaluation of experimental data
The experimental data for viscosity-composition behavior of the blends was compared
with two models: Lees' model with viscosity of the blend, q ~ , expressed by Equation
2.30 and the sheath-core model with q~ expressed by Equation 2.43.
Figure 4.47 and Figure 4.48 show a cornparison between the experimental data and the
models. It can be seen that there is a good agreement between the experimental data and
Lees' model for HDPEl/PP3 blends and the melt viscosity does not follow the sheath-
core model. In the case of HDPE2PA-6 biends, neither of two models predicts the
experimental data. This is in agreement with morphological observation for this blend in
that there was no indication of orientation of the dispersed phase in the flow direction at
the experimental conditions. The multi-layer structure was observed for the HDPEIPA-6
blends but the experimental data does not follow Lees' model and the melt viscosity of
the blend is smaller than those predicted by both models. However, addition of the
slippage factor to the sheath-core model (Equations 2.46) greatly improved the fit.
Figure 4.49 shows this result for the HDPEl/PA-6 blends.
4.4.4 Summary
This section deals with measurement of viscosity of polymer blends. A single screw
extruder-tube combination was used as a capillary rheometer to measure the apparent
melt viscosity of the blends at a constant wall shear stress for HDPEiPP blends and
HDPE/PA-6 blends. The viscosity-composition behavior was compared with two models:
Lees' model and a sheath-core model. The change in melt viscosity of HDPElRP3
blends with composition was monotonie and there was good agreement between the
experimental data and Lees' model. However, the viscosity-composition behavior of
HDPEPA-6 blends was strongly infiuenced by the rnorphology of the blend. The melt
viscosity of the blend lay between those homopolymers when there was a dispersed phase
structure and neither of two models fit the experimental data. The melt viscosity showed
a minimum when there was a multi-layer structure and the experimental data was lower
than those predicted by sheath-core model. This may be associated with the poor
adhesion between the layers and resulting slip. Greater improved agreement was found
by using a shear-induced interlayer slip factor in the sheath-core model.
- Experirnental b t a - - Shelh-Core Mdel
4000 -
3000 -
2000 -
O 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1
BDPE (vol%)
Figure 4.47 (a)
- Lee's Model ExpcrimenW Data SheathZore Model
HDPE (vol%)
(b)
- Lee's Modd Experlmental Opta -- S heathxore Model
Figure 4.47: A cornpanson between the experimental data and the models for HDPE UPP3 blends at two temperatures; 195°C and 260°C a) 0.02 MPa b) 0.05 MPa c) 0.08 MPa.
- - - - -
Sheath-Core Model
Shear StressrO.O1 MPa
HDPE (vol%)
(8)
-+- Experimental Data
....-. Lee's Model
--- SheathCore Model
1 Shear Stres~O.06 MPa
0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 .O
HDPE (vol%)
Figure 4.48: A cornparison between experimental data and the models for HDPE/PA-6 blends, T=230°C a) HDPE2PA-6 blends b) HDPE I/PA-6 blends.
_-______.----------- a . . . .
* * W.. . ..*'... ....
Shear Stress=O.OB MPa
1 1 1 4 4 I I I I
0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 .O
HDPE (vol%)
Figure 4.49: Cornparison of sheath-core mode1 and modified sheath-core mode1 with experimental data for HDPE 1PA-6 blends, T=230°C.
5.0 Conclusions
A process for investigating both shear-induced migration and morphology
development of polymer melt flowing in a long tube was successfÙlly designed and
constmcted. It consists of a long segmented tube attached to a static mixer which is
in turn attached to the end of a single screw extruder. By quenching with water and
breaking apart the tube, samples along its length were obtained for a variety of
processing conditions.
Two analytical methods were shown to be suited to measure composition profiles of
the samples: a Fourier Transform Infrared (FTIR) spectroscopy method and a method
based upon Differential Scanning Calorimetry @SC). Well-known statistical
measures associated with linear regression such as, R', RMSEC, RMSEP, plots of
residuals and plots of 95% confidence intervals were used to evaluate the calibration
curves.
By using the segmented tube process and the developed analytical methods with
various polyethylene-polypropylene, polypropylene-ny(on6 and polyethylene-nylon6
blends, it was quantitatively shown that shear-induced radial migration can occur
during tubo flow.
Calculation of the theoretical rate of viscous energy dissipation per unit length of the
tube used to show that shear-induced migration results are in accordance with
principle of energy minimization.
The morphology of polypropylene-nylon6 and polyethylene-nylon6 blends along the
tube showed viscosity ratio of the dispersed to the continuous phase to be a highly
influential variable: a droplet dispersed phase structure was observed at high
viscosity ratio whereas a multi-layer structure was observed at ratios near unity. The
latter was attributed to a combination of shear-induced deformation and shear-
induced coalescence.
O The effect of morphology development on viscosity-composition measurement by
capillary rheometry was s h o w to be highly significant by using the extrusion-tube
combination as a capillary rheometer. Morphology had a large impact on the
measured viscosity. It was demonstrated that a well known mixing mle (Lees'
model) fits the viscosity versus blend composition data of polyethylene-
polypropylene blends. Data from polyethylene-nylon6 blends with a dispersed phase
morphology and a multi-layer morphology was tit by neither Lees' model or a model
based upon a sheath-core structure. However, better agreement was found by
considering shear-induced interlayer slip factor into the sheath-core model.
6.0 Recommendations
r Numerical solution of the Navier-Stokes equations and energy equations together
should be conducted with the objective of more accurately characterizing shear rate
profiles and rates of viscous dissipation at high shear rates.
A larger variety of viscosity ratios should be examined. One method of rnodifjing the
viscosity ratio of a polymer blend is to add to it polymer of the same type as one of
the components but of a different molecular weight. Another method is to utilize a
chernical reaction with one or more of the cornponents. For example, the injection of
free radical initiator into polyethylene/polypropylene blends during extrusion is well
known to cause a decrease in the viscosity of the polypropylene due to a degradation
of molecular weight and an increase in the viscosity of the polyethylene due to chain
branching and chain extension.
r More studies should be conducted regarding morphology development in the tube.
These studies should be more directed at quantifying shear-induced coalescence and
shear-induced deformat ion.
The effect of morphology on viscosity-composition behavior using the extruder-tube
combination as a capillary rheometer should be examined at diflerent values of wall
shear stress and larger viscosity ratios.
1 .
2.
3.
4.
5.
6.
7.
8.
9.
1 o.
I l .
12.
13.
14.
15.
16.
17.
18.
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H. F. Mark, N. M. Bikales, C. G. Overberger and G. Menges, "Encyclopedia of
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M. J. Folkes, and P. S. Hope, "Polymer Blends and Alloys", Blackie Academic,
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S. S. Young, J. L. White, E. S. Clark and Y. Oyanagi, Polym. Eng. Sci., 20, 798
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8.0 APPENDICES
Appendix 1: Rate of Viscous Energy Dissipation in Tube Flow
For laminar flow in a tube the Navier-Stokes equations simplify to:
where t, is shear stress, p is pressure and r is radius.
d r , 1 a~ +-Te = --
dr r dz
Integrating:
where Ci is an integration constant
For a tube containing different fluids in concentric layers (See Figure 1-l), we can wnte
op Cl* rrr2 = --- +- & 2 r
:. For any layer, i: ap = ---
Te az 2
Assume that each layer of fluid obeys a power law
where n and k are power law's constants.
Then for each layer
The rate of viscous energy dissipation is:
where V is the volume.
Rate of viscous energy dissipation per unit length of tube in each layer can be obtained fiom:
where
Let
The total rate of viscous energy dissipation per unit length of tube in al1 m layers is
obtained:
In the case of one fluid occupying the whole tube
We consider three cases shown in Figure 1.2. In case 1, the high viscosity component
(component 1) is located in the core (low shear rate region), in case Il, the high viscosity
component is located in the shell (high shear rate region) and in case III, as a reference,
the whole tube is occupied by component 3 (q3). It is assumed al1 components are
Newtonian fluids. The effect of interfacial tension is not considered and a constant total
flow rate is also assumed in each case.
Case 1
Set h'=qilrlz (l'ri)
when the whole tube is occupied by component 3:
where a=ql/q3. By setting QI = QiiI, the ratio of pressure drop of Case 1 to Case III can be obtained
Using Equation 2.39:
[(a)' ++(a)')] = $(%),
Combining Equations 1- 19 and 1-2 1 :
ri indicates the position of the interface:
where VI is the volume that occupied by component 1 , Vt is total volume and 01 is the
volume fraction of component 1.
Combining Equations 1-22 and 1-23 :
when the whole tube is occupied by the high viscosity component (a=l):
Case tI
(a)'
Combining Equations 1-27 and 1-28:
r ~ z / R indicates the position of the interface:
where Vz and 42 are the volume that occupied by component 2 and the volume fraction of component 2 respectively.
Combining Equations 1-29 and 1-30:
when the whole tube is occupied by the high viscosity component (a=l):
Combining Equations 1-24 and 1-3 1 or Equations 1-25 and 1-32:
Figure 1.1: A schematic of concentric layers in a tubular flow.
Case I Case m
Case II (1): Hi& viscosity component Case ïII (2): Low viscosity component (3): Cornponent 3 R: Tube radius rll and riz: Radial position of intaface for case 1 and case II
Figure 1.2: A schematic of sheath-core morphology for cornparison: Case i ) high viscosity component is located at low shear rate region (core), Case iI) high viscosity component is located at high shear rate region (shell) Casem) a monocomponent alone is occupied the whole tube.
Appendix II: Statistical analysis [148,149]
Linear Regression
Linear regression is used to fit equations linear in the unknown parameters. SYSTAT was
use for linear regression analysis. Simple Linear Regression (SLR) involves fitting a
straight line:
In the above equation, bo is the intercept of the line with the y-axis and bi is the slope of
the line.
Multiple linear regression @ER) c m be used to fit equations containing more than one
independent or dependent variables. For example, in this work, an equation of the form of
Equation 11-2 was fit by MLR
bo, bi and bz are coefficients to be deterrnined.
The main assumptions used in linear regression are:
i. Al1 of the experimental error is present in the dependent variable. That is, it is
assumed that the independent variable is determined much more precisely than the
dependent variable.
ii. The experimental enor associated with each value of the dependent variable is
independent of the error in any other value.
iii. The errors of the dependent variables are normally distributed.
iv. The error variance ofthe dependent variables is unchanged over the entire range of
the variables.
In this work, linear regression is used to obtain calibration curves. For this type of
application the main consideration is the "goodness of fit" of the equation to the data and
the emor in the predicted values obtained when the equation is used. The "goodness of
fit" measures used were: the multiple correlation coefficient squared, the correlation
coefficient, the standard error of calibration, and the plot of residuals. To quanti@ the
error in the predicted values the 95% confidence interval of the predicted values about the
fitted line are calculated. Also, a "standard error of prediction" is used. Each of these
various measures are described in tum in the following paragraphs.
The multiple correlation coeflicient squared (R*). R~ is used as a measure of the
degree of variation or scattering which is explained by the linear regression model. It is
defined as:
YAi = predicted value from the fit
Yi = expenmental value
y' = xyi/n = mean experimental value
n = number of samples
A value of R~ equal to one shows that the regression model explains 100% of the scatter
in the measured data. The closer is R~ to one, the better the fit of the equation to the
data.
The correlation coefficient (r ). r is a measure of the dependence between the two
variables. For variables x and y the correlation coefficient is calculated as follows:
The value of r can range from -1 to +l. A value of -1 indicates a perfect negative
correlation; a value of +1 indicates a perfect positive correlation; a value of O indicates
no correlation. When r = O the variables are independent
The standard error of calibration (SEC) (also known as the standard error of the
estimate). This quantity is the standard error of calibration (SEC). This statistic is the
standard deviation for the residuals due to differences between the actual values and the
predicted values for samples within the calibration. It indicates the "typical scatter" of the
data around the fitted line. The lower the value of SEC the better the fit. The SEC statistic
is a usehl estimate of the theoretical best precision obtainable fiom the specific set of
data used to develop the calibration equation.
The Root Mean Square Error of Calibration (RMSEC) was used. It is defined by
Equation 11-5 :
n = number of sarnples in calibration set
The plot of residuals. A residual shows the deviation between the observed value and
the predicted value. It is defined as:
Values of ei are plotted versus the estimated values from the regession equation. The
residuals must Vary in size and sign (k) in a random manner about the predicted value if
the equation selected for the mode1 is a good approximation to the true relationship
between the dependent variable and independent variable. Furthermore, residual plots
help to detect the nonlinear nature of the relationship.
Standard Error of Prediction (SEP). The SEP is also temied the standard error of
performance, also known as the standard deviation for the residuals due to differences
between the actual value and the predicted value for samples outside of the calibration set
using a specific calibration equation.
The Root Mean Square Error of Prediction (RMSEP) was used to indicate the prediction
ability of the calibration equation. It is defined by Equation 11-7:
m = number of samples in prediction set
Confidence intervals for predicted values. The 95% confidence interval for the
predicted mean value of the dependent variable for a given value of the independent
variable given by:
The interpretation of this quantity is that if the experiments were repeated to provide new
data values and the fit and 95% confidence interval re-calculated and this was done many
times to give many different sets of intervals, then 95% of these intervals would contain
the true mean value of the dependent variable at that specific value of the independent
variable. If only one prediction of the dependent variable is made at that specific value
then the probability that the calculated interval at this point will contain the true mean is
0.95.
If we consider prediction of an actual observed value of the dependent variable at the
specific value of the independent variable then we rnust consider that the observation also
has error associated with it. Then the 95% confidence interval is given by:
Finally, if we consider prediction of an actual observed value which is the rnean of q
observed values of the dependent variable then the 95% confidence interval is:
Appendix III: Bagley plots
Figure IIL1: End correction determination for PP 1 a) different flow rate ( Q = I O ~ m3/sec) b) upper and lower 95% confidence intervals for Q=0.16~10~ (m3/sec) 11451.
Figure 111.2: End correction determination for PP3 a) different flow rate (Q= 1 o4 m3/sec) b) upper and lower 95 % confidence intervals for Q=O. 15x 1 o4 (m3/sec) [145].
Appendix IV: DSC calibration curves
The equations for the calibration curves are as follows:
a HDPElPP2 blends
HDPE2PA-6 blends
CHDpE = 2.212 + 0.5433 * AHm
PP2/PA-6 blends
where, C and A are the concentration and the area under the melting peak respectively.
Table N.1: Calibration results for immiscible polymer blends using DSC.
Sy s tem
HDPElIPP2
blends
HDPE2/PA-6
blends
PP2lPA-6
blends
Mode1
Eq.N-1 SLR
Eq.lV-2 SLR
Eq.IV-3 MLR
Eq.IV-4 SLR
Figure IV.1: Area under melting peaks as a function of composition for HDPElPP2 blends.
Eq.IV-5 SLR
3
0.989
0.993
0.997
0.996
O. 994
RMSEC
(wt%)
3.041
2.445
i .574
1.852
W S E P
(wt%)
3.322
2.177
2.623
2.165
1.973 2.809
HDPE (Wh)
(a)
Figure IV.2: Area under melting peak versus composition a) HDPEUPA-6 blends b) PP2/PA-6 blends.
O 10 20 30 40 50 60 70 80 90 100 110
Estimated (wPhH DPE)
Figure IV.3: Plots of residual for HDPEl/PP2 blends using DSC a) SLR model (area under melting peak of PP) b) SLR model (area under melting peak of HDPE) c) MLR mode1 (area under melting peak of PP and HDPE).
4 Cali bration Sampies a
O 10 20 30 40 50 60 70 80 90 100 170
Estimated (WhHDPE) (a)
Validation Sarnpies
O 10 20 30 40 50 60 70 80 90 100 110
Estirnated (wt%PP)
Figure IV.4: Residual plot fiom SLR Mode1 a)HDPE2/PA-6 blends b)PP2/PA-6 blends.