Multiphase Polymers: Blends and Ionomers - American Chemical

Chapter 5

Fractionated Crystallization in Incompatible Polymer Blends H. Frensch, P. Harnischfeger, and B.-J. Jungnickel

Downloaded by COLUMBIA UNIV on June 14, 2013 | http://pubs.acs.org Publication Date: July 21, 1989 | doi: 10.1021/bk-1989-0395.ch005

Deutsches Kunststoff-Institut, D—6100 Darmstadt, Federal Republic of Germany The crystallization of the minor component in incompat i b l e polymer blends starts sometimes at d i s t i n c t l y larger undercoolings than in the pure polymer, and proceeds in several separated steps. After a short survey on the history of the effect in the available literature, the several types and the origin of this "fractionated crystallization" as observed in some selected systems are described. The information on the blend which can be deduced from the effect i s discussed, and the consequences for the blend processing and properties are investigated.

The p r o p e r t i e s of i n c o m p a t i b l e polymer blends depend to a l a r g e extent on the mutual d i s p e r s i o n of the components, on the supermolecular s t r u c t u r e w i t h i n the phase of a s i n g l e component, and on the structure of the interface. These structural parameters, i n turn, depend on the processing or mixing conditions and on the strength of the thermodynamic incompatibility of the components as well. These boundary conditions together with the cooling rate control also the s o l i d i f i c a t i o n process of a melt. C r y s t a l l i z a t i o n i s a special s o l i d i f i c a t i o n process. I t i s well known that the c r y s t a l l i z a t i o n of a blend component can d i f f e r remarkably from that of the corresponding pure material. As i n polymers i n general, the course of c r y s t a l l i z a t i o n i n blends i s governed by e q u i l i b r i u m thermodynamics and by k i n e t i c boundary c o n d i t i o n s as well. These boundary conditions change remarkably with blending, thus causing a l o t of technically important and s c i e n t i f i c a l l y interesting effects. Depending on the blend components under consideration, they can o r i g i n i n a large number of physical and physico-chemical phenomena, among them also some having non-equilibrium thermodynamic bas i s . For melt-compatible polymer blends, among other e f f e c t s , a variation of the c r y s t a l l i z a t i o n process due t o a l t e r e d n u c l e a t i o n and growth c o n d i t i o n s has been r e p o r t e d (1-11). The usual mixing induced melting point depression has been observed too (1-5,11,12). 0097-6156/89A)395-0101$07.25A) © 1989 American Chemical Society

In Multiphase Polymers: Blends and Ionomers; Utracki, L., et al.; ACS Symposium Series; American Chemical Society: Washington, DC, 1989.


102

MULTIPHASE POLYMERS: BLENDS AND IONOMERS

For polymer blends with extended i n t e r f a c i a l regions, these phenomena were reported to apply also to the bulk of the phase borders (12-16). The investigation of incompatible polymer blends revealed, e.g., the i n d u c t i o n of s p e c i f i c c r y s t a l m o d i f i c a t i o n s (16), the r e j e c t i o n , engulfing and deformation of the dispersed component by the growing spherulites of the matrix material (4,17-19), and nucleation at the i n t e r f a c e (4,20,23). Polymer blends c o n t a i n i n g one component as f i n e l y d i s p e r s e d droplet suspension exhibit sometimes the phenomenon of "fractionated c r y s t a l l i z a t i o n " which originates in primary nucleation of isolated melt p a r t i c l e s by units of d i f f e r e n t nucleating species (24-33). This phenomenon resembles to some extent the c l a s s i c a l droplet c r y s t a l l i zation i n which c r y s t a l l i z a t i o n i s i n h i b i t e d u n t i l homogeneous nucleation occurs (34-39). Fractionated c r y s t a l l i z a t i o n proceeds stepwise at greatly different undercoolings, these steps sometimes being separated by more than 60°C (25,33). I t i s the aim of the present paper to give a short survey of the history and the available l i t e r a ture on t h i s p a r t i c u l a r blend c r y s t a l l i z a t i o n phenomenon, and to describe and interprete a l l kinds of manifestations of i t as observed so far, considering p a r t i c u l a r l y the authors work. It w i l l turn out that fractionated c r y s t a l l i z a t i o n exhibits occasionally a number of interesting and surprising new effects, the o r i g i n of which w i l l be discussed. F i n a l l y , the q u a l i t a t i v e and quantitative conclusions that can be drawn from the effects on the system under consideration are summarized. History C r y s t a l l i z a t i o n i s a phase t r a n s i t i o n that i s c o n t r o l l e d by nuc l e a t i o n and growth processes (38). The n u c l e i which are necessary f o r the onset of c r y s t a l l i z a t i o n can c o n s i s t , on the one hand, of small c r y s t a l s of the c r y s t a l l i z i n g material i t s e l f which are s t a t i s t i c a l l y c r e a t e d through thermodynamical f l u c t u a t i o n s i n the melt (homogeneous nuclei). On the other hand, heterogeneities due to the presence of i m p u r i t i e s or other m a t e r i a l s can act as n u c l e i too (heterogeneous nuclei). Each nucleating species, i n p a r t i c u l a r those of the d i f f e r e n t impurities, i s characterized by a s p e c i f i c undercooling AT (at non-isothermal c r y s t a l l i z a t i o n ) , or induction time A t (at isothermal c r y s t a l l i z a t i o n ) at which i t induces remarkable cryst a l l i z a t i o n . By experience, the undercooling at which homogeneously nucleated c r y s t a l l i z a t i o n s e t s i n ( A T ) i s u s u a l l y much g r e a t e r than those of the many heterogeneous nucleating agents ( AT^M. Real c r y s t a l l i z a t i o n processes, therefore, are generally heterogeneously induced i f the substances under i n v e s t i g a t i o n are not s p e c i a l l y p u r i f i e d . If several different heterogeneities are present, however, only that with the lowest attributed undercooling A T i s nucleat i n g l y active: when started by the described process ("primary nucleation"), c r y s t a l l i z a t i o n instantly spreads over the whole available material v i a "secondary nucleation" before another type of heterogen e i t y can become e f f i c i e n t . The dynamics of the l a t t e r process depends for a given substance only on the temperature. These necessary preceeding remarks are v a l i d not only for polymers but for a l l kinds of substances. h

h

n Q

h e


5. FRENSCHETAL.

103 Crystallization in Incompatible Polymer Blends

I t was i n 1880 when the Dutchman Van Riemsdyk r e p o r t e d that small gold melt droplets s o l i d i f i e d at much larger undercoolings than the bulk material (40). Similar observations were made later for many other metals, i n d i c a t i n g t h i s to be a basic c r y s t a l l i z a t i o n phenomenon (41-43). T u r n b u l l and Cech (44) i n v e s t i g a t e d i t i n d e t a i l . They remarked that the delayed s o l i d i f i c a t i o n of the droplets (with sizes between lOum and lOOum) a f t e r slow c o o l i n g of t h e i r melt covered a broad temperature i n t e r v a l with a lower l i m i t T that was t y p i c a l for the material under investigation. The approximate r e l a t i o n h o

T

c

0 >

T

=

- ho

T

c

0 )

/

5


o )

( 1 )

( T ^ : m e l t / c r y s t a l e q u i l i b r i u m temperature) held for most metals. The o b s e r v a t i o n could be e x p l a i n e d q u a n t i t a t i v e l y by assuming an exponential law for the c r y s t a l induction rate, and by consideration of the d i s t r i b u t i o n of the nucleating heterogeneities ( a l l of which were assumed to be equally e f f i c i e n t ) over the large number of dropl e t s of d i f f e r e n t size (41,44). The heterogeneous nucleation a c t i v i t y i s terminated at that temperature where homogeneous nucleation sets i n . The course of the c r y s t a l l i z a t i o n before the onset of homogeneously nucleated c r y s t a l l i z a t i o n , t h e r e f o r e , r e f l e c t s the number d e n s i t y of the h e t e r o g e n e i t i e s and the s i z e d i s t r i b u t i o n of the droplets. From the homogeneous nucleation, f i n a l l y , conclusions can be drawn on the interface parameters of the c r y s t a l s (35-39,45). The c r e a t i o n of comparably s t a b l e suspensions of s u f f i c i e n t l y s m a l l polymer d r o p l e t s i s not easy. I t was t h e r e f o r e only i n 1959 when such i n v e s t i g a t i o n s and o b s e r v a t i o n s were r e p o r t e d ctlso f o r polymers. Price (34) found that droplets of polyethylene oxide (PEO) with an average diameter of 5pm c r y s t a l l i z e d only at an undercooling of 65°C. T h i s value i s t o compare w i t h the maximally a t t a i n a b l e undercooling of 20°C i n bulk. T u r n b u l l , P r i c e et a l . (36,37,46) prepared rather stable suspensions (particle diameter some um) of nalkanes, p o l y e t h y l e n e (PE) and polypropylene (PP) i n a thermodynamically inert l i q u i d . They observed the delayed c r y s t a l l i z a t i o n at a low cooling rate of the molten polymer droplets i n a l i g h t microscope and stated t h e i r course to be continuous. Their explanation was the same as f o r the time and temperature dependent course of the c r y s t a l l i z a t i o n i n the l a t t e r . Those droplets which do not contain at least one heterogeneous impurifying p a r t i c l e undergo homogeneous nuc l e a t i o n . The f i n e r the d i s p e r s i o n , the more t h i s n u c l e a t i o n type dominates. Some years later, Koutsky et a l . (35) reported on another b a s i c experiment i n t h i s f i e l d . They sprayed s o l u t i o n s of s e v e r a l polymers on a s i l i c o n o i l f i l m which was spread over a glass plate. During spraying, the solvent evaporated, and polymer droplets of the s i z e of (1...100)um arose. The o b s e r v a t i o n of the c r y s t a l l i z a t i o n during slow c o o l i n g ( c o o l i n g r a t e : 0.1°C/min) of t h e i r melt i n a microscope r e v e a l e d c l e a r l y f o r the f i r s t time that the c r y s t a l l i z a t i o n did not occur continuously but i n d i s t i n c t steps which were i n i t i a t e d at d i s t i n c t l y d i f f e r e n t undercoolings. Most of the materia l , however, c r y s t a l l i z e d again at a maximally possible undercooling which was t y p i c a l for the polymer. This undercooling varied between about 50°C (for PE) and 100°C (for PP) and was a t t r i b u t e d again t o homogeneous nucleation (PE, PP), or to the nucleating e f f i c i e n c y of the interface between the s i l i c o n o i l and the polymer melt (for PEO, polyoxymethylene (POM), i s o t a c t i c polystyrene (iPS)), respectively.


104


The e x p l a n a t i o n of the observed c r y s t a l l i z a t i o n temperature d i s t r i b u t i o n with peaks at d i s t i n c t temperatures i s obvious: the spectrum of undercoolings at which the several c r y s t a l l i z a t i o n steps occur r e f l e c t s immediately the e f f i c i e n c y spectrum of the several nucleating heterogeneous species. It should be recalled that only that heterogeneous species with the lowest s p e c i f i c undercooling AT^ i s active even when several different kinds of them are present. If the dispersion of the polymer i s so fine that not every droplet contains at least one p a r t i c l e of the usually active species (that i s , i f the degree of d i s p e r s i o n i s of the order of magnitude of the number density of those heterogeneous p a r t i c l e s ) , that heterogeneous species with the second lowest s p e c i f i c undercooling AT can become active and so on. The polymer droplets under investigation so far were immersed i n a l i q u i d t h a t d i d not i n f l u e n c e c r y s t a l l i z a t i o n . The d e s c r i b e d e f fects, however, should occur also i f they are inserted i n a surrounding of a s o l i d or molten second polymer, that i s , i n polymer blends where the polymer under investigation i s the minor component, dispersed as f i n e l y as p o s s i b l e i n the matrix polymer. Such m a t e r i a l s should be advantageous i n that sense that they enable the investigat i o n of the c r y s t a l l i z a t i o n by usual t h e r m o a n a l y t i c a l techniques, e.g. by DSC. This, i n turn, would allow a more precise description of the effects. I t i s surprising that such investigations have not been performed for a long time and that corresponding observations were made rather by accident. Only i n 1969, Lotz & Kovacs (26), and O'Malley et a l . (27) r e p o r t e d independently that the PEO component i n a PEO/PS block copolymer c r y s t a l l i z e d i n two steps as revealed by DSC measurements, one of them at about the same temperature at which the PEO homopolymer c r y s t a l l i z e s , and the other at -20°C. The l a t t e r step o c c u r r e d only i f the PEO was d i s t i n c t l y the minor component and f i n e l y dispersed i n the PS matrix, and was attributed to homogeneously nucleated c r y s t a l l i z a t i o n i n the sense of the foregoing considerat i o n s . Romankevich et a l . (24) made a s i m i l a r o b s e r v a t i o n w i t h r e spect to the POM component i n a POM/PE block copolymer. The additional peak was about 13°C below that temperature at which POM usually c r y s t a l l i z e s . This experimental result was confirmed later by Jungn i c k e l et a l . (32,47,48) who observed up to four c r y s t a l l i z a t i o n peaks f o r POM i n a blend with PE and who i n t r o d u c e d the term " f r a c tionated c r y s t a l l i z a t i o n " for the effect. The r e l a t i v e intensity of the several peaks changed with the degree of dispersion of the POM i n the PE matrix. Tsebrenko (31) found also up to three c r y s t a l l i z a t i o n peaks for the POM component i n blends with an ethylene/vinyl acetate copolymer, and explained i t by the delayed c r y s t a l l i z a t i o n of the POM in the interphase between the phase separated components. Hay et a l . (29) , and B a i t o u l et a l . (30) i n v e s t i g a t e d blends of PE and PS, and they found the c r y s t a l l i z a t i o n of the PE component to proceed i n two steps, one at about 100°C and another at about 80°C (29), or 70°C (30) , provided this component was the minor phase and formed insert i o n s i n the PS with a diameter of t y p i c a l l y 10pm. The a d d i t i o n a l peak was attributed to homogeneous c r y s t a l l i z a t i o n as i t did Turnb u l l , P r i c e et a l . (36,37) i n t h e i r d r o p l e t experiments. Y i p et a l . (25,33) i n v e s t i g a t e d the c r y s t a l l i z a t i o n of PP blended w i t h a SBS block copolymer. 'They found two additional c r y s t a l l i z a t i o n peaks at 74°C and at 44°C i f the PP was the minor component. They remarked, moreover, that the usual peak at 107°C disappeared completely when


h e



5.

Crystallization in Incompatible Polymer Blends 105

FRENSCHETAL.

the PP content dropped s u f f i c i e n t l y , and/or became f i n e l y enough dispersed. The peak at 74°C i s located at the same temperature where the main c r y s t a l l i z a t i o n occurred i n the d r o p l e t experiments of T u r n b u l l , P r i c e et a l . (36,37), and Koutsky et a l . (35). The authors, therefore, attributed i t to homogeneous c r y s t a l l i z a t i o n . The peak at 44°C was assumed to be due to the s o l i d i f i c a t i o n of PP i n a s m e c t i c chain alignment. T h i s " c r y s t a l m o d i f i c a t i o n " u s u a l l y a r i s e s i f the melt i s quenched (49). This interpretation, however, i s questionable since the homogeneously nucleated c r y s t a l l i z a t i o n should occur at the highest attainable undercooling i f the droplets have a monomodal size d i s t r i b u t i o n . The experimental facts of Yip et al., on the contrary, h i n t at t h a t the main c r y s t a l l i z a t i o n step i n the r e f e r r e d d r o p l e t experiments (35-37) i s not due to homogeneous n u c l e a t i o n as the authors c l a i m e d but r a t h e r by a c u r r e n t l y unknown heterogeneously nucleating species or surface. This survey on the l i t e r a t u r e indicates that only few data are available on the droplet c r y s t a l l i z a t i o n phenomena i n i n c o m p a t i b l e polymer blends. Moreover, these observations are partly not completely explained, and, where explained, these explanations are partly not s a t i s f y i n g or contradictory. In the next chapter, therefore, experimental r e s u l t s f o r some s e l e c t e d systems as i n v e s t i g a t e d by the authors are presented with the aim to show a l l faces and properties of f r a c t i o n a t e d c r y s t a l l i z a t i o n i n d e t a i l , and t o c o n t r i b u t e t o a better understanding of the o r i g i n of the effect. Theoretical

Considerations

The o v e r a l l c r y s t a l l i z a t i o n k i n e t i c s i n a dispersed system i s determined by several processes. F i r s t , there i s the time and temperature dependent n u c l e a t i n g e f f i c i e n c y of a certain heterogeneous species which, in turn, causes a time and temperature dependent course of the c r y s t a l l i z a t i o n that i t induces. This process depends on the number density of the nucleatingly active heterogeneous species under consideration (which must not necessarily be constant but i n turn can be time dependent) and the dispersion of the c r y s t a l l i z i n g polymer. I t s graph i s smooth without maxima or shoulders i f the mentioned d i s t r i b u t i o n f u n c t i o n s are monomodal. I t can be c h a r a c t e r i z e d by an undercooling AT , e.g., by that temperature at which a given percentage of the material has c r y s t a l l i z e d , or by the temperature of the highest c r y s t a l l i z a t i o n r a t e . I t i s d i f f i c u l t to d e r i v e a n a l y t i c a l expressions for t h i s function which take i n t o account a l l r e l e v a n t s t r u c t u r a l parameters, i n p a r t i c u l a r the s i z e d i s t r i b u t i o n of the droplets, the d i s t r i b u t i o n of the nucleating heterogeneities over the s e v e r a l d r o p l e t s , the c o o l i n g r a t e (with the l i m i t i n g value zero, i.e. i s o t h e r m a l c r y s t a l l i z a t i o n ) , the r e a c t i o n r a t e (turning i n t o e f f i c i e n c y of the n u c l e a t i n g s i t e s ) , and the c r y s t a l growth r a t e i t s e l f . Some considerations and calculations i n t h i s respect can be found, e.g., i n (36,37) and (45). An e v a l u a t i o n of t h i s f u n c t i o n , b a s i c l y , a l l o w s e s t i m a t i o n of the i n t e r f a c e energies between the polymer melt and the heterogeneity and the polymer c r y s t a l faces, respectively. h

I f , a d d i t i o n a l l y , s e v e r a l d i f f e r e n t nucleating heterogeneities are present, the o v e r a l l c r y s t a l l i z a t i o n i s a complicated superposition of the contributions of them a l l . Frequently, the c r y s t a l l i z a t i on that i s caused by a certain heterogeneity i s completed within time


106


or - for nonisothermal conditions - temperature intervals which are s m a l l i n comparison to the d i f f e r e n c e s i n the i n d u c t i o n times or s p e c i f i c undercoolings, r e s p e c t i v e l y , of the s e v e r a l subsequently a c t i v e h e t e r o g e n e i t i e s . The s u p e r p o s i t i o n of t h e i r c o n t r i b u t i o n s , then, leads to a stepwise o v e r a l l c r y s t a l l i z a t i o n . These repeatedly mentioned s p e c i f i c undercoolings can be estimated i n the following way. The f r e e energy 0 of formation of a r e c t a n g u l a r nucleus of c r i t i c a l size at an undercooling AT i s proportional to a "specific i n t e r f a c i a l energy difference" Ay (38V ps — 0*


Ay A

Ay

~

p

/(AT) .

(2)

2

s

i s defined by (38,50):

s

1

y

p

m

- yp*" '^ - y < >

s

ps

Here, the convention y

1 2

+

Vps

(1,2)

( c )

( 3 )

-

= y

s u b s t a n c e

..i'\substance"2»

( s t a t e

o f

substance"l", s t a t e of substance"2") i s introduced; m = melt, c = c r y s t a l , s = substrate, p = polymer). y (m) and Yp (c), followingly, are the i n t e r f a c i a l energies between the nucleating substrate and the polymer melt and c r y s t a l s , respectively. y_(m,c) i s the surface free energy p a r a l l e l to the molecular chain d i r e c t i o n between the c r y s t a l and i t s own melt. The r e l a t i o n ps

Ay

S

= 2y (m,c)

p s

(4)

p

holds for homogeneous nucleation. If one assumes that for the onset of c r y s t a l l i z a t i o n , 0 /kT must be smaller than a certain c r i t i c a l value which i s independent on the material under consideration, and i f one neglects that the c r y s t a l l i zation rate depends also on the temperature dependent mobility of the c r y s t a l l i z a b l e segments, then from Equation 2 the following approximate expression follows for the r e l a t i o n between A y and the undercooling at which the nucleation by two d i f f e r e n t species " A " and " B " gives r i s e to c r y s t a l l i z a t i o n : A

Y

A

/ A

Y

S

B

(

T

A

/

T

B

)

(

A

T

A /

A

T

B

}

2

(

-

I f , i n p a r t i c u l a r , " B " designates Ay

B

5

A

)

the homogeneous nucleation,

= 2y (m,c) , p

(cf. Equation 4), and i f Equation 1 i s v a l i d , AT

o)

B

= T^ /5, and T

then A y

A

o)

fi

= (4/5)T^ ,

i s given by o)

o )

2

Ay /y (m,c) * 6 2 . 5 ( T / T ^ ) ( A T / T ^ ) . A

p

A

A

(5b)

From Equations 5, the r e l a t i v e Ay-values of the heterogeneities can be determined from the corresponding undercoolings. The inversion



5. FRENSCHETAL.

of t h i s statement i s true too. The temperature T at which a c e r t a i n heterogeneity "A" induces c r y s t a l l i z a t i o n depends on i t s A y - v a l u e (Figure 1). Usually, only that heterogeneity with the smallest A y value i s e f f i c i e n t ; v i a secondary nucleation at the created c r y s t a l s , the c r y s t a l l i z a t i o n process spreads over the whole volume once i t has started, and i t i s completed before the undercooling of the heterogeneity with the second smallest Ay-value i s reached. But exactly t h i s e f f e c t i s i n h i b i t e d i f the m a t e r i a l volume i s d i v i d e d i n t o many separated droplets. Among a large number of small polymer droplets, each of volume v , the f r a c t i o n of droplets which contain exactly z heterogeneities of the kind "A" that usually induce c r y s t a l l i z a t i o n follows a Poisson d i s t r i b u t i o n function (45): A

D


f^

A )

= ((M

(A)

Z

v ) /z!)exp(-M D

(A)

v ) D

(6)

(A) where M i s the concentration of randomly suspended heterogeneities and M ^ v i s t h e i r mean number per d r o p l e t . The f r a c t i o n of dropl e t s which c o n t a i n at l e a s t one h e t e r o g e n e i t y of the k i n d "A" i s given by ^ ^ = 1 and amounts to: v

D

z

f

z >

£

A )

>

(A)

= 1 - exp(-M v ). D

(7)

The consideration of a droplet size d i s t r i b u t i o n may somewhat modify t h i s equation. £ ^ describes that part of the droplets and, therefore, of the material that c r y s t a l l i z e s induced by heterogeneity "A". The remainder c r y s t a l l i z e s at a g r e a t e r u n d e r c o o l i n g induced by heterogeneity "B" and so on. For these further c r y s t a l l i z a t i o n steps, the same c o n s i d e r a t i o n s hold. From the r e l a t i v e i n t e n s i t y of the d i f f e r e n t c r y s t a l l i z a t i o n steps, therefore, conclusions can be drawn on the concentration of the r e s p e c t i v e h e t e r o g e n e i t i e s i f the mean size of the droplets i s known. The larger a p a r t i c l e i s , the greater i s the p r o b a b i l i t y of c o n t a i n i n g u n i t s of a c e r t a i n h e t e r o g e n e i t y and, followingly, the p r o b a b i l i t y to c r y s t a l l i z e at the usual temperature . I t i s easy to r e a l i z e that the c r y s t a l l i z a t i o n at the u s u a l temperature which i s induced by h e t e r o g e n e i t y "A" i s completely suppressed i f the r e l a t i o n z

M

( A )

v

D

>

« 1

(8)

holds. It should be pointed out that there i s no d i r e c t physical r e l a tion between the phenomenon of fractionated c r y s t a l l i z a t i o n and the number and the size of spherulites i n the pure polymer. Whereas the occurrence of fractionated c r y s t a l l i z a t i o n i s r e l a t e d t o the r a t i o between the number d e n s i t i e s of d i s p e r s e d polymer p a r t i c l e s and primary nuclei, the size and the number of spherulites are a d d i t i o nally influenced by the cooling rate and the c r y s t a l l i z a t i o n temperature. There i s , therefore, also no r e l a t i o n between the fractionated c r y s t a l l i z a t i o n and the type of the a r i s i n g c r y s t a l l i n e e n t i t i e s (complete s p h e r u l i t e s , stacks of lamellae,...) both i n the pure and in the blended material. There i s , f i n a l l y , no r e l a t i o n between the scale of dispersion which i s necessary for the occurrence of f r a c t i o nated c r y s t a l l i z a t i o n and the spherulite size i n the unblended polymer.



108


Figure 1. Plot (according to Equation 5b) of the s p e c i f i c i n t e r f a c i a l energy d i f f e r e n c e A y a g a i n s t the r e l a t i v e u n d e r c o o l i n g at which a heterogeneity nucleates the polymer. A, B : Heterogeneous nucleations; C : Homogeneous nucleation.


5.

FRENSCHETAL.


The dispersion l e v e l and the p a r t i c l e sizes are governed by the melt rheology of the blend components. Theoretical approaches as well as experiments give evidence that the minor component p a r t i c l e radius R during melt p r o c e s s i n g i s r e l a t e d to the i n t e r f a c i a l energy y between the components, to the matrix v i s c o s i t y r| and to the shear r a t e T by (1,61-63)


R - Y (Tl

T )

-

1

.

(9)

Furtheron, the dispersed droplets are the smaller the closer to unity the v i s c o s i t y r a t i o of the components i s (62-64). T h e i r s i z e s decrease a l s o i f the f i r s t normal s t r e s s d i f f e r e n c e of the d i s p e r s e d phase becomes smaller than that of the matrix (61). The droplet size, moreover, i s i n f l u e n c e d by the tendency t o f u r t h e r break down of elongated p a r t i c l e s due to c a p i l l a r y i n s t a b i l i t i e s (61) as well as by coalescence v i a an i n t e r f a c i a l energy driven viscous flow mechanism. A l l these procedures and dependences a f f e c t the structure formation within their t y p i c a l time scales (61,62). Selected Systems Sample Materials and Preparation. We want to present here the main results of our investigations of the blend systems PE/POM (32,47,48), poly ( v i n y l i d e n e f l u o r i d e ) / polyamide-6 (PVDF/PA-6), and PVDF/poly (butylene terephthalate) (PVDF/PBTP) (Frensch, H., J u n g n i c k e l , B.-J. C o l l . Polym. S c i . , i n press). The PVDF, the PA-6, and the PE that we used were commercial materials, the l a t t e r a s t a b i l i z e r free grade. The POM and the PBTP of our investigations were especially prepared and contained no a d d i t i v e s apart from an unusually low amount of a s t a b i l i z e r . Separate i n v e s t i g a t i o n s r e v e a l e d that a l l components degraded only n e g l i g i b l y during the several preparational steps and the measurements. The dried components were melt mixed in a single-screw laboratory extruder. The extruded strands were regranulated and the extrusion cycle was repeated up to five times. After every cycle, samples were taken f o r the i n v e s t i g a t i o n s . By the repeated e x t r u s i o n , d i f f e r e n t degrees of dispersion were realized. Investigation Techniques. DSC measurements were c a r r i e d out under nitrogen atmosphere. In order to destroy the self-seeding nuclei i n the components, the samples were preheated f o r 5 min at l e a s t 35°C above the m e l t i n g p o i n t s of the higher m e l t i n g component. Then, several c r y s t a l l i z a t i o n and reheating runs were performed at a standard rate of 10°C/min. In some cases, other rates were used too. The d i s p e r s i o n s t r u c t u r e of the blends both i n the melt and i n the s o l i d s t a t e was imaged p a r t l y by l i g h t microscopy (LM), and partly by scanning (SEM) and transmission electron microscopy (TEM). Wide-angle X-ray scattering (WAXS), Infrared (IR) measurements, and torsional pendulum analysis at 1Hz were performed too. Results/Morphology. LM micrographs of the melt s t r u c t u r e of the PE-POM blends at 170°C are displayed in Figure 2. The components are phase separated. The diameter of the dispersed p a r t i c l e s of the POM, that i s the minor component i n these figures, decreases from about 20pm after one extrusion run to about 5pm after five ones.




110

Figure 2. LM micrographs of a PE/POM = 85/15 v o l . - % blend a f t e r one (a), and f i v e (b) e x t r u s i o n c y c l e s . Scale bar corresponds to 30 urn.



5. FRENSCH ET AL.


The TEM micrographs of the PVDF-PA-6 blends (Figure 3) r e v e a l that these components are phase separated too. The p a r t i c l e sizes i n the four times extruded samples vary between 0.1pm and 3pm f o r the PVDF component and 0.05pm and 2pm f o r the PA-6, depending on the composition. The greater p a r t i c l e s have a composite structure since in t h e i r turn they contain p a r t i c l e s of the matrix phase as a further l e v e l of d i s p e r s i o n . With i n c r e a s i n g e x t r u s i o n c y c l e number, the dispersion becomes finer, and the composite character, where i n i t i a l ly present, i s occasionally lost. Torsional pendulum analysis exhibits discrete relaxations at the g l a s s t r a n s i t i o n temperatures of PVDF and PA-6 at -45°C and 50°C, respectively. This also indicates incompatibility of the blend components in the amorphous phase after s o l i d i f i c a t i o n . SEM of f r a c t u r e s u r f a c e s r e v e a l phase s e p a r a t i o n i n the PVDFPBTP blends too (Figure 4). The p a r t i c l e sizes (PVDF: between 1pm and 20pm, PBTP: between 0.1pm and 2pm)) are c o n s i d e r a b l y l a r g e r than those in the PVDF-PA-6 blends of comparable composition. The p a r t i c l e diameters again decrease with increasing extrusion cycle number. R e s u l t s / C r y s t a l l i z a t i o n Behaviour/PE-POM Blends. The PE component i n this blend c r y s t a l l i z e s in a l l samples at almost the same temperature and to the same extent. In contrast, the POM, i f the minor component, c r y s t a l l i z e s i n up to three d i s t i n c t steps ("fractionated", Figure 5) at 146°C, 138°C, and 132°C, the r e l a t i v e DSC peak area of which depends on the number of extrusion cycles, that i s , on the degree of d i s p e r s i o n of that component (Figure 6). The r e l a t i v e i n t e n s i t y of the several peaks, moreover, changes with increasing DSC cycle number to that corresponding to a lower extrusion cycle number. This i n d i cates a coarsening of the dispersion with increasing dwelling time i n the melt. The o v e r a l l degrees of c r y s t a l l i n i t y of both components, however, are almost independent of the extrusion and DSC cycle number, respectively. The POM, in p a r t i c u l a r , exhibits only one melting endotherm at an almost constant temperature for a l l samples. R e s u l t s / C r y s t a l l i z a t i o n Behaviour/PVDF-PA-6 Blends. The c r y s t a l l i z a t i o n of s e v e r a l four times extruded blends (that i s , f o r constant mixing c o n d i t i o n s but v a r i a b l e composition), as s t u d i e d by DSC, i s d i s p l a y e d i n Figure 7. The c r y s t a l l i z a t i o n temperature T taken as the maximum of the c r y s t a l l i z a t i o n curve at a c o o l i n g r a t e of 10°C/min i s at 140°C f o r pure PVDF and at 178°C with a shoulder on the high temperature side for pure PA-6. The most s t r i k i n g result of these investigations has been found with the c r y s t a l l i z a t i o n behaviour of the 85/15 blend. I t i s remarkable that i t shows only one c r y s t a l l i z a t i o n exotherm at 140°C, that i s , at the T of PVDF whereas nothing happens at the usual PA-6 c r y s t a l l i z a t i o n temperature. The WAXS analyses and the DSC melting runs, however, reveal that, nevertheless, the PA-6 c r y s t a l l i z e s to the usual extent. In the 70/30 and 50/50 blends, the PA-6 c r y s t a l l i z e s i n two steps, p a r t l y at 184°C (that i s , at a somewhat higher temperature than i n the pure material) and p a r t l y at 140°C (that i s , as i n the 85/15 blend) as can be derived from comparison of the exotherm and the endotherm peak areas. In the 50/50 blend, a l s o the c r y s t a l l i z a t i o n of the PVDF component s p l i t s up i n t o two steps at 140°C and 116°C. In the 15/85 blend, f i n a l l y , the PVDF c r y s t a l l i z e s merely at the low temperature step at about 112°C. c

Q




112

Figure 3. TEM micrographs of PVDF/PA-6 blends. PVDF i s the dark phase. Scale bar corresponds to 2um (a,b), or 5um (c). a) PVDF/PA-6 = 15/85 v o l . - % , four extrusion cycles; b) PVDF/PA-6 = 75/25 v o l . - % , four extrusion cycles; c) PVDF/PA-6 = 75/25 v o l . - % , one extrusion cycle. (Reproduced with permission from r e f . 68. Copyright 1988 Steinkopff-Verlag Darmstadt.)


FRENSCHETAL.



5.

F i g u r e 4. SEM m i c r o g r a p h s o f PVDF/PBTP b l e n d s . S c a l e b a r c o r r e s p o n d s t o 2um (a,b), o r lOum ( c ) . a) PVDF/PBTP = 15/85 v o l . - % , f o u r e x t r u s i o n c y c l e s ; b) PVDF/PBTP = 85/15 v o l . - % , f o u r e x t r u s i o n c y c l e s ; c) PVDF/PBTP = 85/15 v o l . - % , one e x t r u s i o n c y c l e . (Reproduced w i t h p e r m i s s i o n from r e f . 68. C o p y r i g h t 1988 S t e i n k o p f f - V e r l a g Darmstadt.)




114

Figure 5. DSC c o o l i n g curves of a PE/POM = 85/15 v o l . - % a f t e r one (b) and four (c) e x t r u s i o n c y c l e s , (a): pure POM.


blend

FRENSCHETAL.



5.

Figure 6. Relative area of the POM c r y s t a l l i z a t i o n peaks at 146°C (a), 138°C (b), and 132°C (c) i n dependence on the extrusion cycle number Z.




116

Figure 7. DSC c r y s t a l l i z a t i o n curves of PVDF/PA-6 blends. Four extrusion cycles; parameter: blend composition. (Reproduced w i t h permission from r e f . 68. Copyright 1988 Steinkopff-Verlag Darmstadt.)


5. FRENSCHETAL.


Variation of the cooling rate does not affect the number and the r e l a t i v e i n t e n s i t i e s of the c r y s t a l l i z a t i o n peaks. In particular, the unusual, complete coincidence of the T i n the 85/15 blend occurs at various cooling rates ranging between 0.5°C/min and 100°C/min. The c r y s t a l l i z a t i o n dependence on the extrusion cycle number, that i s , i t s dependence on the mixing intensity, with a fixed compos i t i o n i s shown i n Figure 8. The DSC cooling curve of the four times extruded 75/25 blend exhibits again the coincident c r y s t a l l i z a t i o n of the PVDF matrix and the PA-6 d r o p l e t s at 140°C as a l r e a d y the 85/15 blend d i d (cf. Figure 7). In the one time extruded blend, additionally, a part of the PA-6, possibly the larger domains, c r y s t a l l i z e s at i t s usual T of 184°C and a part of the PVDF, probably the droplets dispersed inside the mentioned larger PA-6 domains, c r y s t a l l i z e s at 113°C. The 85/15 blend exhibits the same behaviour. With increasing number of extrusion cycles, the PA-6 domains become smaller and loose their composite character. This causes the disappearence of the usual high temperature c r y s t a l l i z a t i o n peak of the PA-6 as well as that of the low temperature c r y s t a l l i z a t i o n peak of the PVDF. The l a t t e r , therefore, i s assumed to represent the c r y s t a l l i z a t i o n of those PVDF d r o p l e t s which are i n s e r t e d i n PA-6 d r o p l e t s . Obviously, the described effects, i n p a r t i c u l a r the fractionation of the c r y s t a l l i z a t i o n , depend to a l a r g e extent on the k i n d and on the degree of the dispersion of the minor component. With increasing d i s p e r s i t i v i t y of that component, the magnitudes of i t s a d d i t i o n a l c r y s t a l l i z a t i o n peaks become stronger at the expense of the usual peak. Despite the occasional fractionation of the c r y s t a l l i z a t i o n or i t s suppression at the usual temperature, the DSC heating curves of a l l blends e x h i b i t i n a l l cases f o r both polymers a s i n g l e m e l t i n g endotherm (PA-6: one f o r both c r y s t a l m o d i f i c a t i o n s ) at an almost constant temperature ( v a r i a t i o n range s m a l l e r than 6°C) of about 175°C for the PVDF and of about 217°C/223°C for the PA-6. The r e l a t i v e c r y s t a l l i n i t y of each component a l s o does not change s i g n i f i c a n t l y with composition and extrusion cycle number. In fact, WAXSand IR-analyses show that, i n a l l samples, PA-6 c r y s t a l l i z e s mainly in the f -modif i c a t i o n at an almost constant y/a - r a t i o , and PVDF i n the a-modification. The b a s i c c h a r a c t e r of a l l o b s e r v a t i o n s i s independent on the number of the cooling/heating cycles. c


Q

R e s u l t s / C r y s t a l l i z a t i o n Behaviour/PVDF-PBTP Blends. The pure PBTP (Figure 9) c r y s t a l l i z e s at about 180°C. The c r y s t a l l i z a t i o n temperature T r i s e s remarkably to 189°C and 194°C f o r the one and f o r the four times extruded blends, r e s p e c t i v e l y , a f t e r adding 15 v o l . - % PVDF. In a s i m i l a r way, the pure PVDF c r y s t a l l i z e s at 140°C whereas the corresponding T i n the blends i s between 142°C and 148°C. The PVDF i n the 15/85 blend, e.g., c r y s t a l l i z e s at 147°C a f t e r one extrusion cycle whereas i t c r y s t a l l i z e s at 143°C after four ones. The PBTP c r y s t a l l i z a t i o n i n the four times extruded 85/15 blend i s suppressed i n a s i m i l a r manner as a l r e a d y d e s c r i b e d f o r the PA-6. In place of that, the PBTP c r y s t a l l i z e s at 147°C simultaneously with the PVDF as derived from comparison of the exotherm and the endotherm DSC peaks of the cooling and reheating runs. Again, the m e l t i n g curves show no s i g n i f i c a n t d i f f e r e n c e s t o that of the pure m a t e r i a l s , and a l l d e s c r i b e d e f f e c t s occur a l s o c

Q



118


Figure 8. DSC c r y s t a l l i z a t i o n curves of the PVDF/PA-6 = 75/25 vol.-% blend. Parameter: number of extrusion cycles Z.(Reproduced with permission from r e f . 68. Copyright 1988 Steinkopff-Verlag Darmstadt.)


FRENSCHETAL.



5.

230°C

Figure 9. DSC c r y s t a l l i z a t i o n curves of PVDF/PBTP blends. Parameters: blend composition and number of extrusion cycles Z.(Reproduced with permission from r e f . 68. Copyright 1988 SteinkopffVerlag Darmstadt.)


120



during further cooling/heating cycles. The PVDF, i n particular, as proved by WAXS measurement, c r y s t a l l i z e s i n a l l samples i n the a modification. Discussion. The c r y s t a l l i z a t i o n ot the investigated blends exhibits s e v e r a l d i f f e r e n c e s to t h a t of the pure components. An o c c a s i o n a l occurrence of douple m e l t i n g endotherms (that has not been been described hitherto) of the components during the reheating runs and the s m a l l m e l t i n g p o i n t d e p r e s s i o n s of the minor component can be explained i n the usual manner by imperfect c r y s t a l formation during the - possibly fractionated - c r y s t a l l i z a t i o n , by molecular reorganization and lamellar thickening during heating, and by polymorphism (55-60). The c o n s i d e r a b l e r i s e i n c r y s t a l l i z a t i o n temperature after adding a second component (PA-6 and PBTP i n their PVDF blends) may be due to migration of nuclei across the interface, or an altered chain mobility i n the interface (48). These issues w i l l , therefore, not be considered i n the following. Other effects, however, are related to the topic of t h i s paper and have to be treated i n d e t a i l : — s p l i t up of the c r y s t a l l i z a t i o n of the d i s p e r s e d component i n t o several d i s t i n c t steps (fractionated c r y s t a l l i z a t i o n ; POM blended with PE, both PVDF and PA-6 i n their blends); — complete i n i t i a l suppression of the c r y s t a l l i z a t i o n of the dispersed component, and subsequent f u l l y c o i n c i d e n t c r y s t a l l i z a t i o n with the matrix (both PA-6 and PBTP i n their blends with PVDF). These effects are c l e a r l y connected with the dispersion of the component under i n v e s t i g a t i o n i n t o the other, and they are enhanced i f t h i s dispersion becomes finer. The c r y s t a l growth r a t e s of PVDF, PA-6, and POM amount to at least lOum/min i n the temperature range where t h e i r c r y s t a l l i z a t i o n steps occur (6,52,67). A d i s p e r s e d p a r t i c l e , t h e r e f o r e , once nucleated, c r y s t a l l i z e s promptly and the primary rather than the secondary nucleation i s the rate-controlling factor of the c r y s t a l l i z a t i o n k i n e t i c s of the d i s p e r s e d phase. Thus, the c r y s t a l l i z a t i o n temperatures as observed i n the DSC-cooling run agree roughly with the nucleation temperatures. Discussion/Melt Rheology of PVDF-PA-6 Blends. Normally, the melt v i s cosity of a polymer decreases with increasing temperature T and with i n c r e a s i n g shear r a t e T . Since our samples were processed at T = 230°C w i t h T= (10 ...10 ) s " , the v i s c o s i t i e s amount to T\ (PVDF) = (2xl0 ...2xl0 )Pas (65) and X\ (PA-6) = (7xl0 ...2xl0 )Pas (66). A low i n t e r f a c i a l energy between PVDF and PA-6 melts which may be of the order of lmJ/m , and the close match of the blend component v i s c o s i t i e s at higher shear r a t e s may account f o r the r a t h e r s m a l l p a r t i c l e sizes of the dispersed component and for the occurrence of s u b d i s p e r s i o n s i n s i d e a p a r t i c l e . Furtheron, the s l i g h t l y g r e a t e r v i s c o s i t y of PVDF compared to that of PA-6 may, according to Equation 9, be the reason f o r that the PA-6 p a r t i c l e s are somewhat s m a l l e r than those of the PVDF. _1

3

3

1

2

2

2

D i s c u s s i o n / E s t i m a t i o n of Nuclei Concentrations. The average volumes v of the d i s p e r s e d PA-6 p a r t i c l e s i n the four times extruded PVDF/PA-6 85/15 and 75/25 blends as e s t i m a t e d from the e l e c t r o n micrographs amount to about 4 x l 0 ~ u m and 3 x l 0 ~ u m , r e s p e c t i v e l y . D

3

3

2

3


2

5. FRENSCHETAL.

Crystallization in Incompatible Polymer Blends

Since the PA-6 c r y s t a l l i z a t i o n at i t s usual temperature of about 180°C i s suppressed, and since almost a l l PA-6 droplets c r y s t a l l i z e in a low temperature step together with the PVDF, Equation 8 must be f u l l f i l l e d with as the number density of nucleating impurities a c t i v e i n PA-6 above 140°C. F o l l o w i n g l y , M* ^ should be l e s s than 10pm" . T h i s v a l u e a g r e e s r o u g h l y w i t h t h e number d e n s i t y (0.2...2)urn"** of spherulites and sheaves grown i n pure PA-6 at 140°C. On the c o n t r a r y , the PA-6 p a r t i c l e s i n the one time extruded 75/25 blend are much larger. Accordingly, the c r y s t a l l i z a t i o n peak at the i n i t i a l temperature of about 180°C i s not completely suppressed; there i s , however, a s p l i t up of the PA-6 c r y s t a l l i z a t i o n . The average volume of a PVDF p a r t i c l e i n the PVDF/PA-6 = 15/85 blend after four extrusion cycles amounts to about 8xl0" pm . Moreover, a l l of them c r y s t a l l i z e at the low temperature step. Followingl y , the number d e n s i t y M ^ of the n u c l e a t i n g i m p u r i t i e s which are active i n PVDF at 140°C i s below about 50pm since Equation 8 must be f u l l f i l l e d also for t h i s sample. In the 50/50 blend which exhibits two c r y s t a l l i z a t i o n steps there are, on the c o n t r a r y , a l s o l a r g e r PVDF p a r t i c l e s with a volume of about 4pm , the c r y s t a l l i z a t i o n of which most probably gives r i s e t o the DSC peak at 140°C s i n c e the p r o b a b i l i t y of c o n t a i n i n g a h e t e r o g e n e i t y i s g r e a t e r f o r a l a r g e p a r t i c l e . Therefore, M ' ' of the PVDF should amount a t l e a s t t o 0.2pm" i n order to f u l l f i l l the condition M ^ v s l (cf. again Equat i o n s 6-8). The average volume of dispersed PBTP p a r t i c l e s i n the PVDF/PBTP = 85/15 blend after the f i r s t and the fourth extrusion cycles amounts to about 4pm and 5x10 pm , respectively. By the same arguments as above one has to conclude from the fractionation of the c r y s t a l l i z a t i o n of the d i s p e r s e d PBTP d r o p l e t s even i n the one time extruded blend that the number density of the heterogeneities which nucleate PBTP above 148°C i s below about 0.2pm" . A

3


3

3

3

A

3

D

J

3

Discussion/Estimation of S p e c i f i c I n t e r f a c i a l Energies for PVDF. With T ^ ( a-PVDF) = 459K (53) and T ^ ( ?-PA-6) = 50OK (52), and using Equation 1, one gets T (PVDF) « 94°C, and T (PA-6) s 12 7°C. Both these values are well below the T observed for the components in the PVDF/PA-6 blends which, therefore, i n every case c r y s t a l l i z e nucleated heterogeneously. At a low c o o l i n g r a t e of l°C/min, the temperatures of the fractionated c r y s t a l l i z a t i o n of PVDF are 148°C and 119°C. With y ( m , c ) = 9.7mJ/m (6,54) we obtain with Equation 5b f o r the s p e c i f i c i n t e r f a c i a l energy d i f f e r e n c e s of the two nucleating heterogeneities A y s(148) / A PVDF,s(119) / • o)

o )

hQ

hQ

Q

2

pVDF

=

p

v

=

l l m J

V

D

3

8

m

J

m

2

a

n

d

F

m

Discussion/Coincidence of C r y s t a l l i z a t i o n Temperatures. Let us f i r s t c o n s i d e r the PVDF/PA-6 blend. In view of the n o n a l t e r e d T of PVDF, we suppose that the PVDF c r y s t a l l i z a t i o n induces the PA-6 c r y s t a l l i z a t i o n r a t h e r than v i c e versa. Hence, the j u s t c r e a t e d c r y s t a l s of the PVDF matrix act as n u c l e a t i n g h e t e r o g e n e i t y f o r the PA-6. The A y-value between PVDF c r y s t a l s and PA-6 melt, obviously, i s smaller than that of a l l other heterogeneities which are present i n PA-6 to a s u f f i c i e n t extent except, possibly, the species "A". I t s associated s p e c i f i c undercooling, moreover, must be so s m a l l that the PVDF crystals can induce the c r y s t a l l i z a t i o n of the PA-6 from the instant of their own creation. c


121

122


The coincident c r y s t a l l i z a t i o n of the PVDF matrix and dispersed PBTP p a r t i c l e s i n the 85/15 blend (Z = 4) t a k e s p l a c e a t (142...148)°C, that i s , above the T of pure PVDF. I t i s not c l e a r whether the PVDF or the PBTP c r y s t a l l i z e s f i r s t . In either case, the nucleation of the f i r s t c r y s t a l l i z i n g component may be induced either by a s p e c i e s of n u c l e a t i n g h e t e r o g e n e i t i e s or by the molten second blend component. The newly created c r y s t a l s of the one component, then, act immediately as nuclei for the c r y s t a l l i z a t i o n of the other in the same manner as already described for the PVDF/PA-6 blends. Q


Summary, Conclusions

and Outlook

It was the aim of the present paper to show that c r y s t a l l i z a t i o n i n i n c o m p a t i b l e polymer blends can e x h i b i t a l o t of p e c u l i a r e f f e c t s beside the c l a s s i c a l well known physical and physico-chemical phenomena. The e f f e c t s considered here, i n p a r t i c u l a r , are due to the dispersion structure of such blends, and to the changes in the cryst a l l i z a t i o n n u c l e a t i o n c o n d i t i o n s which are such caused. They are important from a physical, a material s c i e n t i f i c , and a technological point of view as well. The phenomena on which has been r e p o r t e d here are l i n k e d to a droplet-like dispersion of the component under investigation. Theref o r e , they are u s u a l l y e x h i b i t e d only by the minor component i n incompatible polymer blends. Both components, however, can exhibit the e f f e c t s simultaneously i f the dispersion i s composed such that, in turn, a part of the matrix material i s included into the p a r t i c l e s of the dispersed component. These mentioned e f f e c t s are mainly (1) s p l i t up of the c r y s t a l l i z a t i o n into several d i s t i n c t steps, the temperature of which can d i f f e r by several ten degrees; (2) i n h i b i t i o n of the c r y s t a l l i z a t i o n at the usual temperature; (3) c o i n c i d e n c e of the c r y s t a l l i z a t i o n of both components at t h a t temperature at which one of them u s u a l l y c r y s t a l l i z e s , or at another temperature; (4) occasionally homogeneously nucleated c r y s t a l l i z a t i o n . B a s i c l y , the e f f e c t s are caused by the n u c l e a t i n g a c t i v i t y of d i f f e r e n t inhomogeneities. They can become more complicated i f the second component, i n p a r t i c u l a r their just created c r y s t a l s , acts as a c r y s t a l l i z a t i o n n u c l e a t i n g inhomogeneity. Such a - in some cases mutual - nucleating a c t i v i t y i s hidden under usual conditions. F i n a l ly, some blends, e.g. that of PBTP and PVDF, exhibit a most complicated mutual nucleation behaviour: the molten f i r s t component acts as n u c l e a t i n g s u b s t r a t e f o r the second one and becomes then i t s e l f nucleated by the newly created c r y s t a l s of that component. The t e c h n i c a l importance of the d e s c r i b e d e f f e c t s i s obvious. The lowering of the s o l i d i f i c a t i o n temperature, or the broadening of the s o l i d i f i c a t i o n temperature range by o c c a s i o n a l l y s e v e r a l ten degrees influences remarkably the r h e o l o g i c a l boundary c o n d i t i o n s during p r o c e s s i n g . They d e l i v e r , f u r t h e r , an a d d i t i o n a l connection between the degree of d i s p e r s i t i v i t y and the r h e o l o g i c a l m a t e r i a l parameters. The changes in material properties which are such caused are an open q u e s t i o n although the a v a i l a b l e i n v e s t i g a t i o n s do not indicate a strong change in supermolecular structure by the delayed crystallization. The links between the degree and the l e v e l of d i s p e r s i t i v i t y , on


5. FRENSCHETAL.


the one hand, and the type and strength of the p a r t i c u l a r fractionated c r y s t a l l i z a t i o n effect that i t causes, on the other hand, allow an at least q u a l i t a t i v e characterization of the f i r s t by the l a t t e r via suitable reference measurements. By the several effects, moreover, an estimation of s p e c i f i c properties of the components l i k e the absolute or r e l a t i v e amounts of the d i f f e r e n t c r y s t a l l i z a t i o n inducing heterogeneities, and their nucleating e f f i c i e n c y are possible. A determination of the interface energies of the faces of the c r y s t a l s i s a l s o p o s s i b l e i f only the u n d e r c o o l i n g can be reached at which, f i n a l l y , homogeneously nucleated c r y s t a l l i z a t i o n starts. Acknowledgment


F i n a n c i a l support of the A r b e i t s g e m e i n s c h a f t I n d u s t r i e l l e r F o r schungsvereinigungen, grant No. 6015 and 6697, i s acknowledged.

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RECEIVED November 11, 1988


Multiphase Polymers: Blends and Ionomers - American Chemical

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