1995_scott and macosko_morphology development during the initial stages of polymer-polymer blending

7/27/2019 1995_Scott and Macosko_Morphology Development During the Initial Stages of Polymer-polymer Blending

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M o r p h o l o g y d e v e l o p m e n t i n i n i t i a l s t a g e s o f b l e n di n g ." C. E. S c o t t a n d C . W . M a c o s k o

Tab l e 1 Summaryof model blend systems nvestigated.Minor phaseis 20wt% PA

MixerT~ wall set

Major (major) temperatureSystem phase (°C) Type (°C)

Polystyrene/nylon PS 99" Non-reactive 200PS-Ox 107" Reactive 200

Styrene-maleic SMA 138b Reactive 200anhydride/nylon

Copolyester/nylon PETG 81" Non-reactive 200Polycarbonate/nylon PC 157" Non-reactive 230

"Dow Chemical product literature, Vicat softeningpointbDifferential scanning calorimetry, 10°Cmin-eEastman Chemical product literature

a melt index of 7.0. The styrene-maleic anhydridecopolymer (SMA) was ARCO Chemical's Dylark 290. Ithas a repor ted maleic anhydride content of 17 wt%. Theamorphous copolyester (glycol-substituted poly(ethyleneterephthalate), PETG) was Kodar Copolyester 6763,

provided by Eastman Chemical Company. It has areported glass transition temperature of 81°C andmolecular weight M, = 26000. The po lycarbonate (PC)was Dow Chemical Calibre 300-6. This is a general-purpose polycarbonate with a reported melt index of 6.0and Vicat softening temperature of 157°C. All materialswere dried overnight under vacuum at 80°C before thecompounding of blends.

The blends were prepared using a Haake Rheomix 600batch mixer with a Haake System 90 drive. All materialswere fed to the mixer in pellet form. The set temperaturedepended on the system being blended, as specified inT a b l e 1 . Roller blades were employed at 50 rpm. At this

rotor speed, the maximum drag flow (neglecting pressureflow) shear rate in the mixer was 65s -1. This wascalculated on the basis of the rotor speed and theminimum gap between the rotor tip and the mixingchamber wall. The blends prepared were composed of20 wt% dispersed phase unless otherwise specified. Pelletsof the components were mixed by hand in a cup beforeblending in the mixer. The mass of material charged tothe mixer was chosen so that a constant volume of 54 cm 3was achieved for each sample. The densities wereestimated using the data given by van Krevelen 19.

At t=0, the pellet mixture was fed through a chuteinto the preheated mixing chamber under a constant forceof 5 kg. The stock temperature of the materials and the

mixing torque were measured during the blendingprocess. After the specified time of mixing, the rol lerblades were stopped and the mixing chamber pulled offthe blades, leaving most of the material attached to theblades. The sample was cut from the large gap region ofthe blades and dropped directly into a bath of liquidnitrogen to freeze the morphology. The time required

between stopping the mixer and dropping the sample

into the liquid nitrogen bath was 10 to 15s. Based onstandard heat-transfer calculations2°, once the sample isdropped into liquid nitrogen it takes approximately 30 sfor the centreiine of the sample to be cooled to below100°C. This should be sufficient to fix the morphology.Mixing times investigated were 1.0, 1.5, 2.0, 2.5, 3.0, 4.0,

5.0, 7.0 and 15.0 min of mixing. Time zero correspondsto the start of feeding of the pellets to the mixer. The

time required to feed the entire charge of pellets varied

with the system, as specified below.For the model polystyrene/nylon system, the non-

reactive blend was formed by mixing PS with PA andthe reactive blend was formed by mixing PS-Ox with PA.The rheologies of the PS and PS-Ox are nearly identical,as demonstrated by Scott 13. The time required to feedthe entire charge into the mixer, averaged over the runs

made, was 24.5 s and it varied from 22 to 26 s. The blendswith an SMA matrix finished feeding in an average of37 s; the PETG matrix in 34 s; and the PC mat rix in 38 s.

The quenched samples from all of the model systemswere placed in the special Soxhlet cups 13. These are

aluminium cups with a piece of filter paper clamped overthe top of the cup. The paper allows the solvent anddissolved polymer to pass through, but keeps the

insoluble dispersed phase particles in the cup. The paperis reasonably dense but contains holes on the order of1/~m. Some of the smaller particles in the distributionare expected to be lost through these holes during theextraction procedure. However, many particles smaller

than 1/~m are retained by the filter, as demonstrated inseveral micrographs that follow. The matrix was easilyremoved in the Soxhlet using methylene chloride for 7days. For samples that were quenched at short mixingtimes there was essentially one chunk of material left inthe aluminium cup after the Soxhlet procedure. Thischunk was simply transferred to an SEM stub forinspection. For samples that were quenched at longmixing times, the contents of the aluminium cup afterthe Soxhlet procedure was a milky suspension of nylonparticles in methylene chloride. A drop of this suspensionwas placed on an SEM stub and the methylene chloridewas allowed to evaporate. These samples were sputtercoated with 10nm of gold using a Denton DV-502 A.

The SEM micrographs were obtained with a JEOL 84011HRSEM at 5 kV.

For the model systems with SMA, PETG and PCmatrices the sample morphology at 7.0 min of mixing wasobtained from fracture surfaces. Small bars (approx.25 mm by 10 mm by 3 mm) of the sample were prenotchedwith a razor blade and placed in liquid nitrogen for

10min. The bars were then fractured by bending. Thesesamples were coated by evaporation with 30 nm of carbonand then 10nm of gold before observation.

MODEL EXPERIMENTS: DISCUSSION OF THETECHNIQUE

The morphology of polymer blends is very complex inthe initial stages of compounding where an enormousreduction in phase size is taking place. It is difficult toinvestigate such complex morphologies using standardmicroscopic techniques, which involve observation of a

fracture surface or a microtomed slice of the sample. Inorder to overcome the difficulties with sample observa-tion, a model experiment was developed that allowsdissolution of one polymer phase. Generally it is morepracticable to dissolve the matrix phase and leave thedispersed phase. The blend components are selected suchthat one can be easily dissolved by a solvent that has noeffect on the other component. Soxhlet extraction

provides for a very gentle treatment of the sample andfor that reason it is the technique of choice for removal

POLYMER Volume 36 Number 3 1995 463



Mo rpho logy development in ini t ial stages of blending. C. E. Sco tt and C. W. Macosko

6000

50OO

.~ 4000

3000

@

[. 2000

~ . ~ l o o o

0

F e e d i n g T i m e

0 5 10 15

Time (rain.)

Mixing torque as a function of time for: (©) PS/PA blendsi g u r e 2

(data from seven runs are overlaid); (+ ) PS-Ox/PA blend (data from

two runs are overlaid)

of the matrix phase. By this means, one of the blend

components may be removed from the system while theother retains the morphology that it exhibited within theblend. This phase may then be examined directly withscanning electron microscopy. Entire particles of the

phase may then be examined, not just a two-dimensionalslice or a fracture surface. This allows for observation ofcomplex morphological structures because the largedepth of field of SEM enables observation of the wholeparticle. Tilting or shadowing may be used to obtain thecomplete dimensions of the particles.

The major disadvantage of this method is that it mustbe ensured that the dispersed phase retains its shape after

dissolution of the matrix. This may be achieved by usinga dispersed phase that is glassy at room temperature ora crosslinked dispersed phase. The solvent must becarefully selected to dissolve the matrix completelywithout swelling or partially dissolving the dispersedphase. Also, there is loss of some spatial information. Along thread of dispersed phase may break up into a seriesof droplets dur ing blending. The relative positions of thedroplets are lost using this procedure.

Model systems with a nylon dispersed phase werechosen for these experiments to minimize the possibilityof artifacts due to sample preparation. The PA has aglass transition temperature of 136°C. The matrices used

may be dissolved using methylene chloride, which has aboiling point of 40°C. Nylon is very resistant to this

solvent. Also, both reactive and non-reactive systems areavailable. The functionality of the PA will react with thefunctionalities of both the PS-Ox and the styrene--maleic

anhydride copolymer, SMA,

MODEL EXPERIMENTS: RESULTS

A summary of the systems investigated in this study ispresented in Table 1. The blend systems that wereinvestigated most thoroughly were a PS/PA non-reactiveblend and a comparative PS-Ox/PA reactive blend.Scott 13 demons trates that the PS and PS-Ox arerheologically matched at high shear rates and discussesthe interracial chemistry in the reactive blend. The torque

and material temperature time curves for these experi-ments are given in Figures 2 and 3, respectively. The

material temperature of this system levels out toapproximately 203°C at 6min of mixing. The slowapproach to the steady-state temperature that is seenin Figure 3 is typical of all the blends studied here

owing to the processing condit ions selected. During theearly stages of blending, significant temperature gradientsare expected within the sample due to heat-transferlimitations. The implications of this are explored further

in the 'Discussion' section below. Based on the viscositydata obtained for PA and PS the viscosity ratio,~llPA)/q(PS), at the steady-state temperature and 65 s-tis 14. The torque curves for the reactive PS-Ox/PA blendfollow fairly closely the curves for the PS/PA non-reactive

blends. The only evidence for reaction is a slightly highertorque just after the melting peak. Based on the mixingtorques, it appears that the extent of reaction betweenthe two phases has not become significant at 1.0 min ofmixing. However, the morphologies of the blends at7.0 rain of mixing clearly indicate an interfaciat reaction,as demonstrated below.

The morphologies observed at 1.0 min of mixing in the

non-reactive PS/PA blends are presented in Figure 4. Themicrographs show the nylon phase, which remains afterthe polystyrene matrix has been removed. Three largechunks of nylon left after dissolution of the matrix fromthe three pieces of the blend are shown in a low-

magnification picture in Figure 4a. Each of these chunkswas originally a single nylon pellet. The piece in the upperright-hand corner of the picture is a relatively undeformedpellet of nylon. The pieces in the upper left and lowermiddle of the picture each have a region that is relativelyundeformed and another region that has been greatlydeformed during the mixing process. Figure 4b shows ahigher-magnification picture of the piece in the lowermiddle of Figure 4a. In the area of extensive deformation

the material has been stretched out into many ribbons orsheets. Closer examination of these ribbons reveals thatmany of them have holes in them. Higher-magnificationpictures of some of these holes are shown in Figure 4c.These holes are of the order of 10/~m in diameter. Insome areas there are no or only a few holes in the sheet.In other areas the size and concentration of the holes issufficient to form more of a lace-type structure. Figure4c shows evidence of the fragile lace structure breaking

210

200

o

190

180

170

160

F " Ti150

0 5 10 15

T i m e (min.)

Figure 3 Material temper atures as a function of time for (©) PS/PA

blends (da ta fro m seven runs are overlaid); ( + ) PS-O x/PA blend (datafrom two runs are overlaid)

464 POLYMER Volume 36 Number 3 1995



Mo rpho logy development in ini t ial stages of blending." C. E. Sco tt and C. W. Macosko

Figure 4 SEM microg raphs for PS/PA , 1 .0 min mixing. (a) Low -magn if icat ion picture of three large pieces of PA dispersed phase af ter extract ionof PS m atr ix phase. (b) Close-up of one o f the pieces from (a) . (c) Holes a nd lace st ructure observed in r ibb ons an d sheets o f the dispersed phase.(d) Broke n lace st ructure and smal l spherical par t icles o f the dispersed phase

Figure 5 Mo rpho logy of dispersed-phase par ticles f rom PS/P A blend,1.5 min mixin g

u p i n s o m e r e g i o n s o f t h e s a m p l e . Figure 4d s h o w s

h i g h - m a g n i f i c a t i o n p i c t u r e s i n a r e g i o n o f th i s s a m e

s a m p l e . T h e r e a r e m a n y s m a l l i r r e g u l a r l y s h a p e d p i e c e s ,w h i c h a p p e a r t o h a v e o r i g i n a t e d f r o m t h e b r e a k i n g o f

t h e l a c e s t r u c t u r e . T h e r e a r e a l s o m a n y s m a l l s p h e r i c a l

p a r t i c l e s o f d i a m e t e r s f r o m 0 . 5 t o 3 # m . A s s h o w n b e l o w

in Figures 6 t o 9 , t hese sma l l sphe r i ca l pa r t i c l e s a r e t he

s a m e s i z e a s t h e p a r t i c l e s f o u n d i n t h e b l e n d a t 2 . 5 t o

1 5. 0 m i n o f m i x i n g . T h e r e s u l ts g i v e n h e r e a r e f o r a b l e n d

w i t h 2 0 w t % o f PA . S i m i l a r re s u lt s w e r e o b t a i n e d f o r a

b l e n d w i t h 5 w t % P A i n P S .

T h e m o r p h o l o g i e s o b s e r v e d a t 1 .5 m i n o f m i x in g a r e

p r e s e n t e d i n Figure 5. T h e r e a r e m a n y i r re g u l a r ly s h a p e dp a r t i c l e s , w h i c h a p p e a r t o h a v e o r i g i n a t e d f r o m d i s i n t e -

g r a t i o n o f t h e l ac e s t r u c t u r e s h o w n a b o v e . T h e r e a r e a l s o

m a n y h i g h l y e lo n g a t e d p a r ti c le s a s w e ll as m a n y n e a r l y

s p h e r i c a l p a rt i c le s . N o t e t h a t t h e d i a m e t e r s o f th e s p h e r e s

a n d e l o n g a t e d p a r t i c l e s a r e i n t h e r a n g e o f 0 .5 t o 3 / ~m .

T h e m o r p h o l o g i e s o b s e r v e d a t 2 .5 t o 1 5 .0 m i n o f m i x i n g

a r e s h o w n i n Figures 6 t o 9 . T h e p a r t i c l e s o b s e r v e d a t

t h e s e m i x i n g t i m e s a r e n e a r l y a l l s p h e r i c a l . A f e w d r o p

b r e a k - u p s t r u c t u r e s m a y a l s o b e o b s e r v e d . T h e p a r t i c l e

d i a m e t e r s r a n g e f r o m 0 . 5 t o 3 / ~ m .

T h e m o r p h o l o g i e s o b s e r v e d a t 1 .0 m i n o f m i x i n g i n t h e

r e a c t i v e P S - O x / P A b l e n d a r e q u a l i t a t i v e l y s i m i l a r t o

t h o s e o b s e r v e d i n t h e P S / P A n o n - r e a c t i v e b l e n d s . T w o

p i c t u r e s f r o m t h i s s a m p l e a r e s h o w n i n Figure 10. A g a i n ,m a n y r i b b o n s o r s h e e t s w i t h h o l e s a n d l a c e s t r u c t u r e s

P O L Y M E R V o l u m e 3 6 N u m b e r 3 1 9 9 5 4 6 5



M o r p h o l o g y d e v e l o p m e n t i n i n i ti a l s t a ge s o f b l e n d i n g : C . E . S c o t t a n d C . W . M a c o s k o

Figure 6 Dispersed-phase par ticles f rom PS/PA blend, 2 .5 min mixing

that are present in blends at short mixing times. The

diameters of the dispersed nylon phase at 7 min of mixingare in the range of 0.1 to 0.3#m. This is five to tentimes smaller than the si ze of the dispersed phases

observed in the non-reactive PS/PA blends due tothe chemical reaction at the interface. However, the

morphologies here at short mixing times are qualitatively

similar to those observed in other non-reactive blends

shown in this section. Evidently, the initial mechanismof morphology development is little influenced by the

interracial reaction.The PETG/PA blend is not expected to be reactive

under the processing conditions used here. There is a

potential for the formation of a reactive blend in this casebecause of the ester-amide interchange reaction between

the two polymers. However, this interchange reaction isusually only observed in the presence of a catalyst and

at substantially higher temperatures than those usedhere 21'22. The PETGA/PA blend exhibited structures

including sheets with holes and strands (Figure 12).Compared to the other systems investigated here, thereappeared to be more strand structures in this particularblend. The fracture surface of a blend at 7.0 min of mixing

Figure 7 Dispersed-phase par ticles f rom PS /PA blend, 3 .0 min mixing

are evident. The behaviour is qualitatively the same asthat in the PS/PA non-reactive blends. However, at longmixing times, the dispersed phase s ize in the PS-Ox/PAblends is about 0.1 to 0.3pm, five to ten times smallerthan that in the PS/PA blends, due to the influence of

the chemical reaction at the interface.Similar experiments have been pursued with other

matrices in order to determine if these observations holdin other blend systems. The morphologies observed at

1.0 min of mixing with matrices of SMA, PETG and PCwere qualitatively similar to those observed with the

polystyrene matrices.The maleic anhydride functionality of the SMA reacts

with the amine functionality of the amorphous nylon toform an amide or cyclic imide. This has been demon-strated directly through infra-red studies and indirectlythrough examination of reactive and non-reactive blendmorphologies 13. Figure 1 a shows a graphic example ofthe sheeting structures that were observed in an SMA/PAblend. Strand and spherical structures are also shown inFigure l lb from a different portion of the sample at a

higher magnification. These two micrographs againillustrate the variety of dispersed-phase shapes and sizes

Figure 8 Dispersed-phase par t icles f rom PS/PA blend, 7 .0 rain mixing

Figure 9 Dispersed-phase par ticles f rom PS/PA blend, 15.0 min mixing

466 POLYMER Volume 36 Number 3 1995



M orp ho logy developm ent in in i t ia l s tages of b lend ing. C. E. Sco t t and C. W. Ma cosko

m o r p h o l o g i e s t h a t a r e o b s e r v e d a t s h o r t m i x i n g t i m e s

s u g ge s t a n i n i ti a l m e c h a n i s m o f m o r p h o l o g y d e v e l o p m e n t

w h i c h i s s u m m a r i z e d i n Figure 14. T h i s m e c h a n i s m

i n v o l v e s t h e f o r m a t i o n o f s h e e ts o r r i b b o n s o f t h e

d i s p e r s e d p h a s e i n t h e m a t r ix , w h i c h a r e d r a w n o u t o f al a r g e m a s s o f t h e d i s p e r s e d p h a s e . H o w t h e s h e e t o r

r i b b o n i s f o r m e d i s n o t c l e a r . I t c o u l d b e f o r m e d a n d

d r a w n o u t i n t h e f l o w f i e l d i n s i d e t h e m i x e r . S u c h a

s t r u c t u r e c o u l d a l s o f o r m w h e n a l a r g e p i ec e o f t h ed i s p e r s e d p h a s e i s d r a g g e d a c r o s s a h o t s u r f ac e , su c h a s

t h e m i x e r w a l l , t h e m i x e r b l a d e , o r a h o t t e r p o l y m e r

p a r t i c le i n t h e m i x e r . F o r m a t i o n o f si m i l a r s h ee t s h a sb e e n o b s e r v e d b y S u n d a r a r a j 23 i n t w o - c o m p o n e n t m o l t e n

sys tem s sub jec t to shea r . Ow ing to the e f fec ts o f f low and

in te r rac ia l t ens ion these shee t s a re uns tab le and ho les

beg in to fo rm in them . The h o les a re , o f course , f i ll ed

w i t h t h e m a t r i x p h a s e , w h i c h s u r r o u n d s t h e s h e e t o n

e i t h e r si d e. W h e n t h e h o l e s in t h e s h e e t o r r i b b o n a t t a i n

a suf f i c ien t s i ze and c on cen t ra t io n a f rag i l e l ace s t ruc ture

i s f o r m e d , w h i c h b e g i n s t o b r e a k a p a r t d u e t o f l o w a n d

in te r fac ia l fo rces in to i r regu la r ly shaped p ieces . These

p i e ce s a r e o f a p p r o x i m a t e l y t h e d i a m e t e r o f t h e p a r t ic l e st h a t a r e g e n e r a t e d i n t h e b l e n d a t l o n g m i x i n g t i m e s .T h e s e i r r e g u l a r p i e c e s c o n t i n u e t o b r e a k d o w n u n t i l a l l

Figure 10 Structure of the PA d ispersed phase from a PS-Ox/PAblend, 1.0 min mixing

s h o w s t h a t t h e p a r t i c l e d i a m e t e r s a r e i n t h e r a n g e o f 0 .2 t o

l # m . T h i s i s a b o u t t h e s a m e d i a m e t e r a s t h e s t r a n d s i n

Figure 12b.S t r u c t u r e s in t h e P C / P A b l e n d a t 1 . 0 m i n o f m i x i n g

a r e s h o w n i n Figure 13. The ho les in the shee t s he re

a p p e a r t o b e s o m e w h a t i r r e g u l a r l y s h a p e d . T h e l i q u i d

n i t r o g e n f r a c t u r e s u r f a c e o f th e b l e n d a t 7 .0 m i n o f m i x i n gs h o w s t h a t t h e d i a m e t e r s o f th e d i s p e r s ed p h a s e v a r y

f rom 0 . 2 to 0 . 7 /~m . As shown by Figure 13, m a n y n e a r l y

sphe r ica l pa r t i c le s o f th i s s ize were obse rved in the b len da t 1 . 0 m i n o f m i x i n g .

P R O P O S E D M E C H A N I S M O F IN I T IA L

M O R P H O L O G Y D E V E L O P M E N T

T h e m o d e l e x p e r i m e n t s t h a t h a v e b e e n p r e s e n t e d s h o wr e m a r k a b l y s i m i l a r t y p e s o f s t r u c t u r e s a t s h o r t m i x i n gt im es . F ive d i f fe ren t m a t r i c es have been inve s t iga ted . In

e a c h c a s e , t h e r e i s c l e a r l y a d r a m a t i c r e d u c t i o n i n t h e

p h a s e s i z e d u r i n g t h e m e l t i n g o r s o f t e n i n g p h a s e o fc o m p o u n d i n g . T h i s o b s e r v a t i o n i s co n s i s te n t w i t h f i n d i n g si n s e v e ra l o t h e r s y s t e m s9-17 . Ke y aspec t s o f the s t ru c ture

o f t h e d i s p e r s e d p h a s e a r e s h a r e d b y a l l o f t h e m o d e l

b l e n d s y s t e m s t h a t w e r e i n v e s t i g a te d .O n t h e b a s is o f t h e o b s e r v a t i o n s p r e s e n t e d a b o v e , t h e

Figure 11 (a) Sheet structures of the PA dispersed phase n a SMA /PAblend, 1.0min mixing. (b) S trand and ribbon structures of the PAdispersed phase in a SMA /PA blend, 1.0min mixing




Mo rph olog y development in ini t ial stages of blending." C. E. Sco tt and C. W. Macosko

Holes form In ther ibbon due to interracialInstability.

The lace br eaks into Irregularly shapedpieces w ith diam eters appreetinmtelyequa l to the u l t imate partk le s/ze.

m

10 Ixm

The s ize and concentra t ionof holes becomes sufficien tfor a lace s tructure to beformed.

F i g u r e 14 Proposed mechanism for initial morphology development in polymer blends

m

2 p.m

D r o p a n d c y l i m l e r b r e a k u pform the f ina l s p h e r i c a lparticles.

use parameters such as stress, deformation rate, or total

strain, which are more closely related to materialdeformation processes. However, owing to the complexnature of the flow field in this mixer and the couplingwith heat-transfer considerations, the distribution of suchparameters throughout the sample is unknown. The timeof mixing is the only parameter that is accurately knownfor every portion of the sample. Experiments utilizingmore simple flow fields will be required in order to in-vestigate the effects of the more fundamental parameters.

Heat-transfer considerations are critical in determiningthe morphology development at short mixing times.There is substantial variation in temperature, stress andstrain from one location in the sample to another.Extensive deformation and domain size reduction occur

in the regions that have been heated above their softeningor melting points and have been exposed to high stresses.Interfacial tension effects become more important as thephase sizes are reduced. In addition, the effects ofcompatibilizing interfacial reactions will not becomesignificant until large amounts of interfacial area havebeen created. The SMA/PA system reported in thepresent work is a fast reacting system. However, it

exhibited the same mechanism of morphology develop-ment at short times as that observed in similar non-reactive systems.

In the experiments reported here, the solid mixture ofpellets was quickly heated in the mixer. The cold pelletshave been force fed into a hot mixer. This is intended to

model the extrusion process where the pellet mixture istypically fed into a relatively cool feeding zone and then

quickly conveyed into a hot barrel zone. As discussed

above, localized heating of the sample is a criticalcomponent of the mechanism observed here. Thus,different results may be obtained in isothermal studiesor with different modes of sample heating. For example,Shih e t a l . 1 6 have used a much slower heating rate onblends that included a rubbery minor phase. Theyidentified several morphological transformations thatoccurred under those conditions. Such transformationswould be obscured in experiments similar to the presentstudy where heating rates are very fast and localizedheating is important.

It is not clear how general the morphologies observedhere or the mechanism proposed here may be. The

relative softening or melting rates of the blend com-

ponents could be important. For the materials invest-igated here, the relative softening rates are probably welldescribed by the glass transi tion temperatures. As shownin T a b l e 1 , in the model systems presented here the glasstransition temperature of the matrix was varied over asubstantial range. Qualitatively, the mechanism of de-formation appears to be the same for these blends inwhich the glass transition temperature of the dispersed

phase is higher than or near that of the continuous phase.However, for systems such as dispersion of rubbers inthermoplastic matrices where the glass transition tem-perature of the minor phase (rubber) is significantly lowerthan that of the major phase, a phase inversion such asthat described by Shih e t a l . 1 6 is expected to be part of

the mechanism. In this case, the minor phase softens firstand init ially becomes the continuous phase. After melting




M o r p h o l o g y d e v e l o p m e n t i n i n i t i a l s t a g e s o f b l e n d i n g . C . E. S c o t t a n d C . W . M a c o s k o

or softening of the major phase , phase inv ers ion occurs

so that the major phase becomes the continuous phase .

The work reported here has only dealt with one

dispersed-phase polymer. T he type of deformation that

occurs ear ly in the melt ing process may be dependent on

the polymer molecular weight and leve l of crystal l inity ,

am ong oth er things . The drawing behav iour of the

polymer is a cr i t ical e lement in the proposed mechanism.

The rh eology o f the polym er as a function of temperatureand shear may be critical. Unfortunately, the selection of

model sys tems for these particular experiments is con-

strained by several practical considerations.

The mech anism m ay also depend on the type of mixer

that is used. H owev er , Sundararaj e t a l . 15 have observed

the s ame me c hani s m o f mor pho logy de ve lopme nt w i th

a twin-screw extruder . Lindt and Gh osh tv have a lso

observed sheet formation during blending in a s ingle-

screw extruder.

The c lassic m echanism of drop break-up is certainly

important in morphology development. Extens ive l i tera-

ture exis ts concerning this mechanism and excel lent

reviews are given by Ralliso n 3-~'~' and Bentley and

Leal 3v. This w ould occur as an integral pa rt of the

mechanism that has been proposed above . I t would be

expected to grow in importance at intermediate and long

mixing t imes , Cons ideration of the problem of the

break-up of a drop or f i lament of one f luid inside anoth er

has been the top ic for a large num ber of investigations.

There is also a general consensus in the literature that

coalescence of droplets during blending m ay be important.

Unfortunately, there are relatively few published invest-

igations of this phe nom enon 38 4o.

C O N C L U S I O N S

In order to determine the morphology at short mixing

times, m od el blend systems are investigated, wh ich allow

the matrix to be dissolved away so that the dispersed

phase may be observed direct ly us ing scanning e lectron

microscopy. The dispersed phase for the model experi-

ments is an amorphous nylon. The matr ices inc lude

polystyrene, an oxa zoli ne fun ctiona l polystyrene, a

s tyr e ne -male i c anhydr ide c opo lyme r , an amor phous

copolyester and a polycarbonate . These experiments

dramatically reveal the primary m od es of particle

deform ation and the nature of the mo rph ologies at short

mixing t imes. The intermediate mo rph ologie s have been

demonstrated to be quite complex. The e lectron micro-

graphs suggest a mech anism for mo rph ology dev elop-ment at short mixing times that is summarized in F i g u r e

I 4 . Redu ction o f phase do ma in s ize begins pr imari ly from

the formation of a sheet or r ibbon o f the dispersed phase .

Owing to interfacial and flow forces, holes form in the

ribbon and grow until a lace structure is formed. This

lace is then broken down into irregularly shaped particlesand finally into nearly spherical particles.

T he in it ia l me c hani s m o f mor pho logy de ve lopme nt i s

certainly intimately connected with the melting or

softening process. Significant differences in temperature,

stress and strain lead to large differences in phase size

mor pho logy fr om loc a t ion to loc a t ion w i th in a s ample .

These results suggest that compounding equipment

and operations may be des igned to promote partic le s izereduction in the melting stage in order to take advantage

of the mechamisms outl ined here . Understanding the

me c hani s m o f mor p ho logy de ve lopm e nt may a l s o have

importan t implication s for optimiz ing the form and order

of addit ion of various co mp onen ts during blending

operations .

A C K N O W L E D G E M E N T S

The authors grateful ly acknowledge the support of aNat iona l Sc ie nc e Foundat ion Gr aduate Fe l lows hip , a

Univers i ty of Minnesota D issertation Fel low ship, and a

Plastics Inst itute of America F el lowsh ip for the support

of C. Scott at various t imes during the pursuit o f this

research. Support for this work was also provided by

DuPont and Ge ne r a l E le c tr i c .

R E F E R E N C E S

c~

0

I

2

3

14

15

16

17

18

I

2 0

2l

2 3

24

25

26

272~

29

3 0

31

32

33

3 4

35

3 6

37

38

39

4O

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