Recent research on
the stability of entanglements of polymer melts and on the correlation between
viscosity improvement during processing and entanglement stability led to the
discovery of a new property of the liquid state of polymers which is not
explained by the current models in polymer physics: it is called “sustained
orientation”. In simple terms, by manipulation of the stability of
entanglements, it is possible to create and maintain quasi-stable at high
temperature in an amorphous polymeric melt (say 120 oC above Tg) a certain state of orientation that was
induced by a mechanical deformation. This discovery is discussed in details in
recent publications [1-6] and appears to contradict our current understanding
of polymer relaxation properties.
The manipulation of
entanglements is done by Rheo-Fluidification [4-6]. Under certain Rheo-Fluidification
processing conditions, the viscosity reduction of the melt induced by the
combination of shear-thinning and strain softening can be preserved in the
pellets granulated at the exit of the disentangling processor. These “treated”
pellets display sustained-orientation, i.e. a lower viscosity when they are
reheated in a melt flow indexer, or in a dynamic rheometer after they have been
compressed into disks.
The Rheo-Fluidification Processor (Figs 19, 7, 8).
In a Rheo-Fluidification
processor a melt extrudes continuously through treatment zones where it is
submitted to a combination of shear-thinning and strain-softening via the use
of cross-lateral shear vibration superposed to pressure flow (originated by an
extruder feed). Fig. 19 shows 2 treatment stations (11) and (22) for a melt
flowing from left to right to exit at the end of the processor, where its
viscosity is measured continuously by an in-line rheometer and it is water
cooled, granulated into pellets. The “treatment” in stations (11) and (22) is
sketched in Fig.7: the melt flows through (from left to right) a gap “3” where
the upper gap surface is static and the lower surface is rotated and
(optionally) oscillated. Both surfaces contain small ribs, “12”, detailed in
Figs 8a and 8b, which create local vibrational extensional flow by squeezing
and un-squeezing the melt as it is being swept forward helicoidally. Shear
thinning is controlled by the shear rates, which add up vectorially from all
types of flow (longitudinal and cross-lateral, vibratory or not), and strain
softening is controlled by frequency and the strain amplitude of the
cross-lateral shear component. The
rotation of the rotor in station 1 and station 2 were in opposite directions, a
situation which we casually called “comb to the left-comb to the right”, referring
to the sweeping induced by the ribs in relation to the rotation direction.
Sustained Orientation (Figs 20,21).
Fig. 20 shows the
viscosity of the exiting melt (PMMA) just after it has been “treated”, i.e. at
the end of the second station of the two-station Rheo-Fluidizer shown in Fig.
19. Temperature profiles are different in both stations. Rotation speeds and
vibrational frequencies are also different in both treatment stations . The rheometer that measures the exiting melt
viscosity is not far from the last treatment zone, but still, it takes the melt
about 2 minutes to get there, and, at that temperature, which is 120 oC
above the Tg of the polymer, the melt relaxation time is very small,
of the order of 0.001 sec for that molecular weight. . The
x-axis in Fig. 20 is the extrusion time,
different parameters being tested until a “successful processing window” is
apparently found, seen as the final value of viscosity, 500 Pa-s , which is 1/3
of the extrapolated “un-sheared” in-line viscosity 1,500 Pa-s. In
Fig. 20, we observe a substantial drop of the melt viscosity at t= 20 min and
another one at t=75 min, which is obtained by just changing the processing
parameters of vibration in the treatment zones. Even if it is understood from
theory that the melt viscosity could be decreased in the treatment zone
under a different set of shear-thinning and strain softening conditions, the
same theory predicts that the thermal-mechanical history should be erased
totally in the 2 minutes times it takes to reach the rheometer, 2 minutes being
120,000 times greater than the longest relaxation time. One should not be able
to observe any viscosity change at the exit of the Rheo-Fluidizer unless the
orientation of the melt induced by shear-thinning and strain softening can be
sustained 120 oC above Tg, an impossible proposition according to our current
understanding of polymer physics!
Pellets were made of the
treated melt by passing it through a strand die, and cooling the strands in
water before pelletization. Two batches of 75 kg each of “treated” pellets were
prepared by this procedure and sent for testing. By this comment I mean to say
that the process was steady and could produce continuously “disentangled melts”
that turned into “disentangled pellets” (the way I summarized the experiment)..
The melt flow index (MFI)
of these frozen-in treated pellets is plotted in Fig. 21 versus the in-line
Rheometer viscosity value which varies as a function of the chosen processing
parameters. The correspondence is remarkable and means that, indeed, it is
possible to retain in the pellets a large portion of the viscosity reduction
observed in the melt due to the Rheo-Fluidification process. The orientation of
the melt induced by shear-thinning and strain softening has been preserved and
has survived re-heating in the MFI barrel still maintaining a 100% lower
viscosity than the reference (the melt with no treatment).
Furthermore, the rheological properties of those pellets were
studied, after the pellets were compressed into new samples as if they were a
new polymer grade. Disks were molded to be studied by dynamic rheometry. The
viscosity reduction observed at the exit of the Rheo-fluidizer could survive a
new heating in a molding press, and, additionally, a study of viscosity vs time
in a rheometer at a high temperature. Yet, the molecular weight was hardly
changed (~3%) to justify the viscosity reduction and it was also observed that
the viscosity gradually returned to the value it should have at the
corresponding temperature (the Newtonian value), indicating that the changes
were reversible. Actually, it took 24 hours at 235 oC for the
viscosity to return to its original value!. Again
note that the sample in Figs 20 and 21 is a linear PMMA deformed 120 oC
above its Tg, and thus, in terms of the conventional understanding of the
rheology of polymer melts in the terminal zone, our results imply that the
polymer has been retaining its orientation for a time 86.4 million times longer
than its “longest relaxation time” (calculated from the cross-over point in a
frequency sweep at the same temperature). This is totally incomprehensible in
terms of our present understanding of “entanglement”, the corner stone of long
chain physics.
These experiments could
be interpreted by stating that the entanglement
became unstable producing an increase of Me, the molecular weight
between entanglement (thus the wording which I continue to use
“disentanglement”). The instability of the entanglement lasted 24 hours! But
there is no explanation to why Me can vary independently of the relaxation
time, and be increased (or decreased) by relatively low shear forces; there is no explanation in the current
theories for an unstable entanglement network resulting in an unstable liquid
state for polymers, and for how it could be correlated to non-linear
viscoelastic effects.
The conclusions:
At least these
experiments suggest that the classical concept of Me to describe entanglements
is too simplistic and its usefulness is probably limited to the linear range of
viscoelasticity.
In the worst case scenario, the whole foundation of polymer physics,
based on its understanding of entanglements, must be reformulated. See Ref. 3.
More reading:
1. J.P Ibar, “ The Great Myths of Polymer Rheology. Part I.
2. J.P. Ibar “ The Great Myths of Polymer Rheology”
Part II. Transient and Steady State. The question of the entanglement
stability. Journal of Macromolecular Science, Part B, 49, 1148 -1258
(2010).
3. J.P. Ibar, “The Great Myths in Polymer Rheology,
Part III: Elasticity of the Network of Entanglements”, J. Macrom. Sci. Part
B, Phys. 52:222-308, 2013.
4. J.P.Ibar, “Processing polymer melts under
Rheo-Fluidification flow conditions: Part 1. Boosting shear-thinning by adding
low frequency non-linear vibration to induce strain softening.”. J.
Macromol. Sci. Part B, Phys,, 52:411-445, 2013 (publication on line November 1st 2012. DOI:
10.1080/00222348.2012.711999).
5. J.P. Ibar, “Processing polymer melts under
Rheo-Fluidification flow conditions: Part 2. Simple flow Simulation”. J.
Macromol. Sci. Part B, Phys., 52:446-465, 2013 (publication on line : November
1st 2012)
DOI: 10.1080/00222348.2012.712004)
6) J.P. Ibar, “Mixing Polymers under
Rheo-Fluidification Conditions”, Macromolecular Symposia, Special Issue, 11th International
European Symposium on Polymer Blends, 2012. In press.
