16 ABS Resins: The Relation between Composition and Rheological Behavior ANTONIO CASALE, ANTONINO MORONI, and CESARE SPREAFICO Montedison S.p.A., Centro Ricerche, Divisione Petrolchimica, 21053 Castellanza (VA), Italy
The purpose of our study was a rheological evaluation of the effect of composition on the properties of ABS resins in the molten state. Steady-state viscosity was determined over a wide range of temperatures and shear rates. The shear modulus in the molten state was determined by measurement of the diameter of the extrudate. ABS resins in the molten state behaved as an amorphous homophase polymer. The effect of the elastomer phase on the viscoelastic properties which characterize the behavior of the continuous matrix, i.e. monomer friction coefficient and molecular weight between entanglements (M ), was calculated by the application of the molecular theories. The significance of these properties in heterophase systems is discussed. e
A
crylonitrile-butadiene-styrene ( A B S ) resins are typical heterophase systems. L They consist of particles of crosslinked rubber, covered b y a layer of graft resins (elastomer phase) embedded in a continuous matrix (resin phase). It is w e l l known that the elastomer phase has a major effect on the A B S polymers. V a r y i n g the structure of the elastomer phase changes the properties of the resin. A n increase i n the elastomer phase induces the following changes i n physical properties: (a)
i n the solid state, blend rigidity decreases while toughness
increases;
and (b) i n the molten state, viscosity increases ( J , 2, 3, 4) while the elastic component (defined as extrudate expansion) decreases ( 5 ) , w h i c h agrees with findings for other heterophase systems (3, 6). Relatively little information is available on the rheological behavior of A B S polymers (4, 5, 7, 8, 9, 10). Particularly little work of fundamental nature seems to have been done on the relation between A B S rheological properties and their composition, probably because of their complex structure. It is there fore difficult, on the basis of the published data, to develop a rational theory on the effect of the dispersed particles on the flow behavior of the blend. A number of papers were published on the viscoelastic properties of two-phase 172
16.
CASALE E T A L .
173
ABS Resins
systems with different morphology (blends of two homopolymers or of a polymer and a block copolymer). Considering their morphology and the effect of the elastomer phase on the properties below T , A B S polymers may be expected to behave i n the molten state (a) as a suspension with high solids concentration or as a colloidal dispersion ( 4 ) , a n d (b) as a blend of a homopolymer and a diluent or of two homopolymers. In our laboratory a systematic study was made with the aim of relating the composition of A B S polymers and their rheological properties. T h e findings enabled us to advance a hypothesis on A B S flow behavior and on the role of the elastomer phase. They also suggested a rheological criterion for polymer compatibility. Finally, on the basis of a method described previously ( I I ) , it was possible to use the rheological data to predict A B S processability. This paper, the third i n a rheotechnics series, reports on rheological studies performed at γ higher than 10" sec" . In two-phase polymer systems, Rosen and Rodriquez (3) hypothesized anomalous behavior (yield shear stress) at γ = 10" . Recently, Zosel (9) verified this hypothesis experimentally for A B S systems. A study of the rheological behavior of A B S resins at very l o w γ is now i n progress at our laboratories. T h e correlation between rheological be havior and moldability w i l l be discussed i n detail i n a subsequent paper. g
1
1
1
Experimental A B S polymers were made by mechanically mixing the resin phase ( S A N ) and the elastomer phase. This last phase was prepared by grafting the same monomers onto a preformed rubber. B y this procedure, the composition and molecular weight of the resins could be very accurately controlled before mixing. T h e following variables were included: an elastomer phase content of 0 - 4 0 % , an acrylonitrile content of the S A N of 2 0 - 3 3 % , and a weight average molecular weight of the S A N of 68,000-150,000. Table I lists the composition of the A B S resins used i n this study. The melt viscosity data were obtained i n a conventional manner using an Instron capillary rheometer over the temperature range of 1 8 0 - 2 4 0 ° C . T h e capillary had a 90° entry angle, and it was 5 c m long and 0.125 c m i n diameter. The well known equations were used to calculate apparent viscosity. Shear rates at the wall, calculated assuming a Newtonian fluid, were corrected for nonparabolic velocity profile using the Rabinowitch equation. N o correction was made for entrance effect because of the length-to-diameter ratio of the capillary used. The diameters of the frozen extrudates were measured b y micrometer after annealing. T h e data are reported as the ratio of diameter of extrudate to diameter of capillary. Basic Background Non-Newtonian viscosity, η, is expressed i n molecular theories for amor phous one-phase polymer as η/ηο = F (Τ λ) (1) where η is the Newtonian viscosity, γ is the shear rate, and λ is a quantity proportional to the terminal relaxation time. T h e function, F , depends on the theory considered (Bueche or Graesslay) while its argument is still the same, λ, according to the modified Rouse theory, is a function of the structure and of the molecules, and it is expressed b y the friction coefficient. 0
174
COPOLYMERS,
Table I. ABS Sample
1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 β
POLYBLENDS,
AND COMPOSITES
Composition of the A B S Resins
Elastomer Phase,
%
Resin Phase, % AN
Resin Phase,
26
68,000
26
117,300
26
145,700
26
226,000
20
117,000
33
119,000
0 20 30 40 0 20 30 40 0 20 30 40 0° 20 30 40 0 20 30 40 0 20 30 40
M
w
S A N with broader molecular weight distribution.
As is well known, viscoelastic functions measured at different temperatures, molecular weights, and concentrations can be superposed i n a single master curve reduced to a reference temperature, T , molecular weight M , and con centration C , b y using appropriate shift factors a , a , and a (12). Equation 1 then becomes η /η ο = F (y ατθίδα c) (2) 0
0
0
T
M
c
Obviously, viscosity can be reduced not only to η , but also to a different value of viscosity. It is also possible to use the shear stress instead of the viscosity. Molecular M e a n i n g of a . T h e effect of temperature on viscosity is re lated to its effect on the friction coefficient, w h i c h , i n turn, depends on the fractional free volume according to the equation: 0
T
In ζ = In ζ
ΐη
+ j-
(3)
where β is a numerical constant close to unity, and ζ is the inherent friction coefficient that is independent of temperature and molecular weight beyond a limit value. It is assumed that the fractional free volume at any temperature, /, increases linearly w i t h temperature above the glass temperature T according to the equation: / = /„ + α (Τ - T ) (4) g
g
where f is the fractional free volume at T , and a is the thermal coefficient of expansion of the fractional free volume above T . Both f and a can be calcu lated by the W i l l i a m s - L a n d e l - F e r r y ( W L F ) equation if the shift factor a is known. The applicability of the time-temperature superposition principle, not only to homopolymers but even to A B S polymers, has been demonstrated by (J
g
g
g
T
16.
CASALE E T A L .
175
ABS Resins
Scalco, Huseby, and Blyler ( 8 ) , Zosel ( 9 ) , and Bergen a n d Morris ( 1 0 ) . Prest and Porter (13) applied the same principle to homopolymer blends [poly ( 2,6-dimethylphenylene oxide )-polystyrene]. Recently some papers were published on triblock copolymers of styrene-butadiene-styrene and on their blends w i t h polybutadiene (14, 15). Triblock copolymers can be considered heterophase material as the different constituent blocks are thermodynamically incompatible w i t h each other, and, consequently, polystyrene domains are en closed i n polybutadiene (continuous matrix). T h e findings indicate that these systems are i n general thermorheologically complex, so that the shift factor a depends not only on temperature but also on time. These conclusions have been extrapolated to other two-phase systems. Molecular M e a n i n g of a . T h e dependence of the reduced viscosity of melt polymers on molecular weight is given b y r
M
[XW
(Μ V
\a*]
M
4
m
where a is the square end-to-end distance per number of monomer units. Since a is not dependent on molecular weight, a is only a function of the molecular weight ratio. Molecular M e a n i n g of a . W i t h homophase polymers, the effect of the diluent on viscosity is attributable to: the change i n the friction coefficient, ζ; SL change i n the entanglement factor, Q due to a modification of the entangle ment spacing; and a possible change i n the a . T h e last factor is negligible i n comparison w i t h the others. T h e relative importance of the first two factors depends on test temperature; at temperatures much higher than T , the effect of the second factor is dominant (12). 2
2
M
r
N
2
g
Viscosity Dependence on Molecular Weight, Temperature, and Elastomer Phase Content The shear stress-shear rate data were treated as follows. (a) T h e curves obtained at different temperatures for each S A N of different molecular weight (distribution and acrylonitrile content being equal) and the relative blends were reduced to a reference temperature T = 210°C by the reduced variables method (16). T h e log of the shift factor a plotted versus the reciprocal of the temperatures gave a straight line. T h e activation energy for viscous flow ( Δ Ε ) was calculated from the slope of the line (17). The activation energies for the different blends are listed i n Table II. ΔΕ was independent of S A N molecular weight, and it decreased linearly as the elasto0
T
Τ
Τ
Table I I .
Activation Energy, AE , and Shear M o d u l u s , G , of the Different A B S V T
% Rubber
M
68.000 117^00 145,700 226,000 a 6
W
b
Δ
ET
30.6 29.2 29.0 30.3
80
20
0
SAN
G 50.0 49.5 51.6 35.0
Λ
Κτ
25.6 23.0 29.0 26.7
G
Δ
120.0 129.0 132.0 82.0
ET
23.3 23.0 27.0 24.8
ΔΕτ is given in kcal/mole, and G in dyne/cm Χ 1 0 . Sample with different molecular weight distribution. 2
-4
40
G 192.0 207.0 200.0 135.0
Δ
ET
21.0 18.4 23.9
G 387.0 345.00 375.0 266.0
176
COPOLYMERS,
POLYBLENDS, A N D COMPOSITES
mer phase content was increased. AE , of the pure elastomer phase, obtained by extrapolation, was 5 - 1 0 kcal/mole, w h i c h is close to the value for rubber. This finding contrasts w i t h that of H u g u e t and Paxton (4); they reported that the activation energy for the pure elastomer phase was 0 kcal/mole. 2
1
ELASTOMER PHASE %
20^i
Ί
Figure 1.
^
2
1
1
SAN (M = 68,000) and ABS blends w
a. Master curves reduced to 210°C;
and b. shift factor vs. elastomer phase content
Figure 2. a. Master curves reduced to 210°C
N ^ S E C
ABS blends
and 0% elastomer phase content; and b. shift factor vs. molecular weight
16.
CASALE E T A L .
177
ABS Resins
lgfra a .a SEC r
Figure 3. ABS master curves reduced to 210°C SAN
R
M
and 0% elastomer phase content
M w = 68,000
(b) It is well known that i n polymer-plasticizers systems, flow curves for different concentrations can be superimposed (16). Consequently, the same procedure was applied to A B S polymers. The master curves at T = 2 1 0 ° C for each of the resins and their A B S blends were reduced to a reference content ( 0 % ) b y shifting the experimental curve along the abscissa (see Figure l a ) . A g a i n a straight line was obtained when the log of the shift factor a was plotted versus elastomer phase content. F r o m the slope of the line, the effect of the rubber phase on polymer viscosity (defined ΔΕ b y analogy w i t h AE ) was calculated (see Figure l b ) . 0
R
Β
T
(c) B y superimposing the master curves at T — 2 1 0 ° C and graft content 0 % for S A N of three different molecular weights, a third shift factor a and the relative slope AE were obtained (Figure 2 ) . This procedure was very satisfactory, considering that i n a single master curve 500 experimental data, referred to four rubber phase contents, three molecular weights (distribution being equal), a n d five temperatures were superimposed (Figure 3 ) . The same procedure was applied to one S A N of different molecular weight distribution and to three samples w i t h different acrylonitrile contents i n order to determine the effect of these variables on the rheological behavior of A B S polymers. The following conclusions were drawn from the findings. (a) A B S polymers in the molten state behave as one-phase polymers with regards to viscosity dependence on temperature, molecular weight, and elasto mer phase content. Viscosity dependence on temperature agrees w i t h the findings of Scalco et al. (8), Zosel ( 9 ) , and Bergen and Morris (10). T h e first authors performed tensile stress relaxation experiments on a commercial A B S polymer, and they concluded that the temperature dependence on the viscoelastic behavior of polyblends was quite similar to that of many amorphous homopolymers above their T . O n the other hand, recent evidence indicates that the time-temperature superposition principle cannot be applied without modification to block copolymers and their blends w i t h homopolymers (14, 15). According to our data, this concept cannot be extended to A B S polymers over our experimental range of graft content ( 0 - 4 0 % ). 0
M
M
g
Figure 4.
ABS extrudate expansion vs. τ at different temperatures and elastomer phase contents SAN
M w = 145,000
(b) Consequently, the molecular theories developed for one-phase poly mers (27,18, 19) can be applied to polyblends. However, the reduced viscosity is a function not only of S A N molecular weight and of temperature, but also of rubber phase content and type. (c) N o n - N e w t o n i a n behavior at the same shear rate increases w i t h molecular weight (as expected) and rubber content. ( d ) The effect of the molecular weight distribution of the S A N is of funda mental importance. As expected, the master curve of a broad distribution sam ple could not be superimposed on the others. The effect of acrylonitrile content is very slight over the considered range. Extrudate Expansion Extrudate expansion data, B, are defined as the ratios of extrudate diame ters to capillary diameters. Β data obtained at different temperatures were plotted as a function of shear stress (see for example Figure 4 for S A N w i t h M. = 145,000). F r o m Β data, the shear modulus of the molten polymers, G — w/yify be calculated. T is the shear stress at the w a l l and y is the recoverable elastic strain. In fact, Β was recently related quantitatively to the y corresponding to r (20) by the equation: w
T
R
c
a
n
W
R
w
" • - ï ' . [ ( ' + à ) " - è ] It was shown that the theoretical values calculated by Equation 6 are i n good agreement over the range of practical interest w i t h the empirical relation: Β = 1 + 0.155
y
R
(7)
16.
CASALE E T A L .
179
ABS Resins
These findings led to the following conclusions: (a) Plots of Β vs. T were linear over the range T = (1-40) Χ 1 0 d y n e / c m , w h i c h agreed with the findings of Rosen and Rodriguez for a different heterophase system ( 6 ) . (b) Temperature d i d not effect extrudate expansion over the experi mental range of 1 8 0 - 2 4 0 ° C , shear stresses being equal. (c) The molecular weight of the S A N (distribution being equal) had no noticeable effect on the extrudate expansion of the resins and of the blends at the same elastomer phase content. (d) A s expected, Β of the broader distribution sample was higher; the shear modulus, G , (see Table II and Figure 5) was not dependent on tempera ture, shear stress, or S A N molecular weight over the experimental ranges. 5
w
W
2
"*"θ
ϊδ
25
35
40"
55
60
C L A S T O M C R PHASE R %
Figure 5.
Log of the shear modulus vs. the elas tomer phase content
(e) Progressive addition of elastomer phase definitely decreased the ex pansion, w h i c h agreed w i t h previous reports on A B S resins (5) and high impact polystyrene (2). H a n ( 5 ) , on the basis of his previous data on blends of polystyrene and polypropylene, suggested that melt elasticity goes through a maximum at a certain blending ratio. W e , however, have observed that the diameter of the elastomer phase extrudate was almost equal to the capillary diameter, even if it was difficult to collect precise data because of the poor consistency.
180
COPOLYMERS,
POLYBLENDS,
A N D COMPOSITES
Discussion As has already been stated, the verified possibility of extending the re duced variables principles to A B S resins makes it possible to treat these typical heterophase systems as blends of amorphous homophase polymers and plasticizers. One possible explanation is that over the experimental γ range it is not possible to separate the contributions of the two different phases, and the mate rials w i l l behave as homophase polymer. I n fact, long-time molten polymer rheology experiments measure viscoelastic processes over the entire molecule, and, as a consequence, molecular compatibility is evaluated ( 1 3 ) . O n the other hand, high frequency and/or low temperature tests involve the main chain as well as the side chains of the polymer system; the segmental miscibility of the polymer-polymer system is then evaluated. It is important i n experimental measurements of polymer compatibility to evaluate the actual size of the volume subject to the test. As was already stated (see Figure 6 ) , the temperature dependence of the shift factor a is a function of the elastomer phase content. T h e strong effect of the rubber content on the temperature dependence of the shift factor a could be explained b y an increase i n free volume of the S A N resin induced b y the elastomer phase, as was suggested b y Prest and Porter (13) for polystyrenepoly (phenylene oxide) blends. In order to verify this hypothesis, log a experi mental data for S A N and relative blends were used to calculate the W L F parameters and, i n turn, the free volumes (f ) at the reference temperature ( T ) and the thermal expansion coefficients (a) b y the equation: T
T
T
0
0
16Ô
vfô
ϊδο
ϊθδ
20Ô
210
220
Sô
2S0
2feo
T C M P E R A T U R C °C
Figure 6.
Temperature dependence of the shift factor for the ABS blends
16.
CASALE E T A L .
181
ABS Resins
90i
20.
10
ol -40
-30
Figure 7.
-20
-10
0
10
20 (T-210)°C
30
Determination of the WLF parameters at Ύ = for the ABS blends 0
40
210°C
The W L F and the free volume parameters are listed i n Table III. A s can be seen, the free volume of the blends increased as the elastomer phase content increased. Only w i t h S A N resins has it been possible to determine C / and C and, consequently, fJB and a/B. T h e findings (f /B = 0.0313 and a/B — 6.3 Χ 10" ) agree with the values reported by Ferry for polystyrene (0.033 and 0.032; 6.9 a n d 6.3 Χ 10" ) (21). T h e same parameters could be calcu lated for the different blends b y assuming that: P
2
g
4
4
JR
= f,
R
+
«R
(T
-
(9)
T, ) R
where subscript R refers to the blends. This analysis was made b y Prest and Porter (13). Table III.
W L F and Free-Volume Parameters
Rubber Content, %
Cl°
c °
fo/B
α/Β, X 10-*
fg/B
0 20 30 40
4.477 2.970 2.222 1.572
156.7 122.07 105.1 91.805
0.09699 0.1463 0.1954 0.2762
0.619 1.20 1.86 3.01
0.0313 — — —
2
182
COPOLYMERS,
POLYBLENDS, AND COMPOSITES
W i t h A B S resins, however, many difficulties arise i n applying Equation 9 since dynamic measurements revealed clearly the T peaks of the two com ponents. In particular one must decide: if f is constant below T ^ or Grubber' ° f T , and the measurement system. F o r example, it must be decided if the resin phase transition must be considered as T and if its value is changed by the elastomer phase. A c c o r d i n g to some authors, the S A N Τ could be modified by the presence of the elastomer phase. Scalco et al. reported that the T of their sample was 85 °C w h i c h was 20°C lower than the S A N Τ (8). Bergen and Morris assigned T values of 8 4 ° C and 97°C respec tively to two different A B S samples on the basis of findings i n other thermal and mechanical experiments (10). However, as is w e l l known, experimental results are a function of the system of measurement, and this also applies to T . The study of A B S free volume is very important i n relation to its role in the elastomer phase reinforcing effect; as suggested b y N e w m a n n and Strella (22, 23), the elastomer phase mainly induces a yielding in the matrix. A triaxial stress field i n the environment of the dispersed particles causes a local increase i n free volume w h i c h permits energy-consuming cold flowing phe nomena. The increase i n free volume caused by the elastomer phase w o u l d in turn bring about a decrease i n the friction coefficient and, consequently, a decrease in viscosity. Experimentally, we found an increase in A B S viscosity. As we have already stated, the effect of the diluent on viscosity, in the case of one-phase polymers, is related to changes in both the friction coefficient and the entanglement factor. F o r one-phase polymers, the latter is predominant under our experimental conditions. W e assume that the increase in blend viscosity could be explained by a change in the state of matrix entanglements induced by the elastomer phase. Bergen and Morris, on the basis of the num ber of grafted chains, suggested that the region surrounding each particle be considered a region of high chain entanglement density (10). The same idea was applied to branched polystyrene (24). The effect induced by the elasto mer phase on the matrix entanglement state could be evaluated as a decrease in average, molecular weight between entanglements, M*. B y this term we mean a molecular weight between entanglements averaged over the whole sample, taking into account both the regions w i t h and without particles. This evaluation could be made by comparing the flow curves of the different blends at equal free volume. The free volume of the S A N at T = 210°C is 0.0913. Unfortunately, the corresponding temperatures of the blends are well beyond the experimental range ( 1 4 8 ° , 155°, and 1 6 8 ° C ) , and a longer procedure must be applied to calculate the effect of rubber on M*. The shift factor at a constant temperature, a , is the constant ratio between the relaxation times at two different elastomer phase concentrations and two different friction coefficients. g
g A N
gR
t n e
v
a
m
e
gR
gR
g
ff
g
g
g
n
R
β
·
ί Τ
>
-
(Jfcr · t
r. -
(10)
The reference content, as already stated, is zero, that is M — M £ S A N - B y substituting Equation 3 i n Equation 10, we obtain: Co
ι
ίτ\
log a (T) R
=
oil 2.4 log
. 2.303 /?
'o
M
2.303 Β
-\
Me JR Experimentally, we have found (see Figure 1) that R
Jo
e
g
A
X
and ζ
=
0
n
n
(11)
16.
CASALE E T A L .
183
ABS Resins
log a (T ) R
(12)
= m(T„)R
0
where m ( T ) is a function of the reference temperature, T . B y combining Equations 11 and 12, w e find that 0
0
In the iso-free-volume state, Equation 13 becomes: log M *
= log M
eR
eo
(14)
- nR
This equation reveals the dependence of the A B S average molecular weight between entanglements (M *) on the molecular weight between entanglements of the pure resin ( M ), on the amount of the elastomer phase, and on factor η = dlog a /dR w h i c h depends on the elastomer phase structure. T h e value of the constant η can be calculated b y substituting i n Equation 13 the experimental values of / , f , and m , calculated at T = 2 1 0 ° C and three different rubber contents. W i t h the elastomer phase used i n this study, CJ{
c
R
0
0
R
log M * eR
= log Λ/,„ -
2.0 · R
(15)
However, experimentally it was calculated that the value of constant η w o u l d be changed 1 0 % b y increasing the elastomer phase content from 0 % to 4 0 % . This variation was considered reasonably acceptable, considering the com plexity of the elaboration. The hypothesis of the rubber effect on the change i n M of the pure resin can be confirmed b y the extrudate expansion data. According to the rubber elasticity theory: 6o
Ρ
RT M7
(16)
F r o m Equations 14 and 16, we obtain log GR = log
-
nR = log G
€o
- nR
(17)
W i t h our graft polymer, log G Η
log Gen + 2.0 R
=
Experimental findings (see Figure 5) indicated that G the equation:
(18) R
log GR
= log G
et)
can be expressed b y
+ 2.05 R
(19)
G is then independent of S A N molecular weight (which agrees w i t h theory), and it is exponentially dependent on rubber content as predicted b y E q u a tion 17. The agreement between the experimental and predicted values of the constant of the exponents was very good. T h e Μ of the S A N copolymer, as calculated b y the extrudate expansion data, was 64,000. This value was higher than that for the polystyrene homopolymer. F r o m the value of the S A N Μ , it is possible to calculate the M * of the different blends b y Equation 15. A t an elastomer phase content of 4 0 % , M * was equal to 10,000. As we have stated above, extrudate expansion of the pure elastomer phase was negligible. A t higher graft contents, the relation between shear modulus and elastomer phase content probably could change. It is thereR
Γ ( )
Γ()
ej
CR
184
COPOLYMERS,
POLYBLENDS,
A N D COMPOSITES
fore impossible to derive the M * value for the pure graft phase b y extrapola tion. F r o m a scientific point of view, further studies must be performed i n order to understand the physical meaning of the average molecular weight between entanglements of a two-phase polymer, considering that A B S polymers are composed of highly crosslinked polybutadiene particles surrounded by a region of high chain entanglement density and a continuous matrix. A variation in the friction coefficient probably superposes on the state of entanglement variation. In any case, the constant η permits characterization of the elastomer phase independently of S A N molecular weight and content. T h e lower its value, the closer the viscoelastic behavior of the A B S polymers is to that of the pure resin. O u r method, then, provides a rheological criterion for measuring the effect of a graft copolymer on a continuous matrix. A s was discussed above, the master curves relative to a reference temperature and rubber content: eR
= F \j (α )α an] τ
of the three S A N samples of different molecular weights were shifted onto the curve of the sample w i t h lowest molecular weight. T h e shift factor a (Figure 2b) is related to the molecular weight according to the equation M
log a.M = 3.0 log M - 3.0 log M (20) Equation 20 is similar to Equation 5. Its slope, however, is 3.0 instead of 4.4. The cause of this discrepancy is unknown. However, it should be remembered that a factor of 3.4 was found i n stress relaxation experiments with polystyrene and poly(a-methylstyrene) ( 2 5 ) . a
Conclusions A B S polymers i n the molten state behaved as one-phase amorphous poly mers i n shear modulus and viscosity. The elastomer phase increased the average free volume of the system as well as the value of the thermal expansion coefficient of the free volume, a. The a value could be related to the reinforcing effect of the graft polymers. The elastomer phase caused a change i n the entanglement state of the pure resin w h i c h could be evaluated as a change i n the average molecular weight between entanglements, M *. T h e compositional dependence on M * is given b y CJ
log M * eR
ei
= log M
eo
- nR
The effect on M * was greater than that on free volume; consequently the viscosity of the system was increased b y increasing the graft content. T h e change i n M * also induced an increase i n shear modulus. The constant η is a graft property dependent on its structure, but, more important, it was independent of S A N content and molecular weight. Its value was closely related to A B S viscosity and viscoelastic behavior. ej
VJ
Acknowledgments The authors wish to express their sincere appreciation to S. Baldoni for his assistance i n the experimental work and i n the elaboration of experimental data. H e l p f u l discussions w i t h G . Ajroldi are also gratefully acknowledged.
16.
CAS A L E E T A L .
Literature 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25.
ABS Resins
185
Cited
Ebneth, H . , Beohm, K., La Nuova Chim. (1969) 10, 35. Hagan, R. S., Davis, D . Α., J. Polym. Sci. Part Β (1964) 2, 909. Rosen, S. L., Rodriguez, F., J. Appl. Polym. Sci. (1965) 9, 1601. Huguet, M . G., Paxton, T. R., Amer. Chem. Soc., Div. Polym. Chem., Prepr. 11 (2), 548 (Chicago, September, 1970). Han, C. D., J. Appl. Polym. Sci. (1971) 15, 2591. Rosen, S. L., Rodriguez, F., J. Appl. Polym. Sci. (1965) 9, 1615. Dzhagarova, Ε . K., Sov. Plast. (1971) 12, 23. Scalco, E., Huseby, T. W., Blyler, L. L., J. Appl. Polym. Sci. (1968) 12, 1343. Zosel, Α., Rheol. Acta (1972) 11, 229. Bergen, R. L., Morris, H . L., Proc. Int. Rheol. Cong., 5th, Tokyo, Japan, 1970, p. 433. Casale, Α., Moroni, Α., Ronzoni, I., Polym. Eng. Sci. (1974) 14, 651. Ferry, J. D., "Viscoelastic Properties of Polymers," p. 524, Wiley, New York, 1970. Prest, Jr., W. M . , Porter, R. S., J. Polym. Sci. Part A-2 (1972) 10, 1639. Fesko, D. G., Tschoegl, N. W., J. Polym. Sci. Part C (1971) 35, 51. Choi, G., Kaya, Α., Shen, M . , Polym. Eng. Sci. (1973) 13, 231. Ferry, J. D., "Viscoelastic Properties of Polymers," p. 292, Wiley, New York, 1970. Casale, Α., Moroni, Α., Civardi, D., Int. Cong. Plast. Elast., Milan, October, 1972. Ferry, J. D . , "Viscoelastic Properties of Polymers," pp. 195 and 247, Wiley, New York, 1970. Middleman, S., "The Flow of High Polymers," Interscience, New York, 1968. Cogswell, F. N., Plast. Polym. (1970) 38, 391. Ferry, J. D., "Viscoelastic Properties of Polymers," p. 316, Wiley, New York, 1970. Newman, S., Strella, S., J. Appl. Polym. Sci. (1965) 9, 2297. Strella, S., Appl. Polym. Symp. (1968) 7, 165. Fujimoto, T., Narukawa, H . , Nagasawa, M . , Macromolecules (1970) 3 (1), 56. Ferry, J. D., "Viscoelastic Properties of Polymers," p. 415, Wiley, New York, 1970.
RECEIVED April 4, 1974.
