Note: Descriptions are shown in the official language in which they were submitted.
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SHRINK FI~MS AND METHOD FOR MAKING FILMS HAVING MAXIMUM HEAT SHRINK
This invention relates to an improved method for preparing
orienting polyolefin films. In particular, this invention relates to a
method for biaxially orienting polyolefin films wherein stretching or
orientation conditions are defined to maximize the unrestrained shrink
response of the film. The invention also relates to a shrink film and
to a shrink film made by the novel method wherein in each case the
shrink film comprises as the shrink control layer at least one
homogeneously branched ethylene interpolymer having a polymer density of
less than 0.9l g/cc.
Food items such as poultry, fresh red meat and cheese, as
well as nonfood industrial and retail goods, are packaged by various
heat shrink film methods. There are two main categories of heat shrink
films - hot-blown shrink film and oriented shrink film. Hot-blown
shrink film is made by a hot-blown simple bubble film process and
oriented shrink film is made by elaborate processes known as double
bubble, tape bubble, trapped bubble or tenter framing. Heat shrink film
can be monoaxial or biaxial oriented.
The shrink packaging method generally involves placing an
article(s) into a bag (or sleeve) fabricated from a heat shrink film,
then closing or heat sealing the bag, and thereafter exposing the bag to
sufficient heat to cause shrinking of the bag and intimate contact
between the bag and article. The heat can be provided by conventional
heat sources, such as heated air, infrared radiation, hot water,
combustion flames, or the like. Heat shrink wrapping of food articles
helps preserve freshness, is attractive, hygienic, and allows closer
inspection of the quality of the packaged food. Heat shrink wrapping of
industrial and retail goods, which is alternatively referred to in the
art and herein as industrial and retail bundling, preserves product
cleanliness and also is a convenient means of bundling and collating for
- - accounting purposes.
The biaxial heat-shrink response of an oriented polyolefin
film is obtained by initially stretching fabricated film to an extent
several times its original dimensions in both the machine and transverse
directions to orient the film. The stretching is usually accomplished
while the fabricated film is sufficiently soft or molten, although cold
drawn shrink films are also known in the art. After the fabricated film
is stretched and while still in a stretched condition, the stretch
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orientation is frozen or set in by quick quenching of the film.
Subsequent application of heat will then cause the oriented film to
relax and, depending on the actual shrink temperature, the oriented film
can return essentially back to its original unstretched dimensions,
i.e., to shrink relative to its stretched dimension.
While the temperature at which a particular polymer is
sufficiently soft or molten is a critical factor in various orientation
techniques, in general, such temperatures are ill-defined in the art.
Disclosures pertaining to oriented films that disclose various polymer
types (which invariably have varying polymer crystallinitie_ and melting
points), simply do not define the stretching or orientation temperatures
used for the reported comparisons. Nor do such disclosures disclose
whether the particular orientation temperature used corresponded to an
optimum temperature in regards to the reported shrink responses or other
15 desired properties for a particular polymer. US Patent 9,863,769 to
Lustig et al., WO 95/00333 to Eckstein et al., and WO 9~/07954 to Garza
et al., are three examples of disclosures wherein stretching or
orientation temperatures are ill-defined or unspec~fied.
The direct effect of density or crystallinity on shrink
response and other desired shrink film properties such as, for example,
impact resistance, are known, for example, from WO 95/08441. That is,
even where the orientation temperature is presumably constant, lower
density polymer films will show a higher shrink response and improved
impact resistance. However, the effect of density, crystallinity and
compositional homogeneity on optimum orientation temperature is not
known. In the prior art, there are only general rules of thumb or
generalized teachings relating to suitable stretching conditions. ~or
example, in commercial operations, it is often said that the temperature
at which the film is suitably soft or molten is just above its
respective glass transition temperature, in the case of amorphous
polymers, or below its respective melting point, in the case of semi-
crystalline polymers.
- _ An example of teaching that's beyond ordinary rules of thumb
(but is nevertheless fairly generalized~ is provided by Golike in US
35_ Patent 4,597,920. Golike teaches orientation should be carried out at
temperatures between the lower and higher melting points of a copolymer
of ethylene with at least one Cg-Clg a-olefin. Golike specifically
teaches that the temperature dif~erential is at least 10~C, however,
Golike also specifically discloses that the full range of the
temperature differential may not be practical because, depending on the
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particular equipment and technique used, tearing of the polymer film may
occur at the lower end of the range. At the higher limit of the range,
Golike teaches the structural integrity of the polymer film begins to
suffer during stretching ~and ultimately fails at higher temperatures)
because the polymer film then is in a soft, molten condition. See, US
Patent 4,597,920, Col. 4, lines 52-68 bridging to Col. 5., lines 1-6.
The orientation temperature range defined by Golike (which is based on
higher and lower peak melting points) generally applies to polymer
blends and heterogeneously branched ethylene/a-olefin interpolymers,
i.e compositions having two or more DSC melting points, and does not
apply at all to homogeneously branched ethylene/~-olefin interpolymers
which have only a single DSC melting point. Golike also indicates that
a person of ordinary skill can determine the tear temperature of a
particular polymer and discloses that for heterogeneously branched
interpolymers having a density o~ about 0.920 g/cc, the tear temperature
occurs at a temperature above the lower peak melting point. See, US
Patent 4,597,920, Col. ~, Example 4. However, Golike does not teach or
suggest how a person of ordinary skill in the art of shrink film can
optimize the orientation process as to stretching temperature at a given
stretching rate and ratio to maximize the shrink response.
Hideo et al. in EP 0359907 A2 teach the film surface
temperature at the starting point o~ stretching should be within the
range of from 20~C to about 30~C below the melting temperature of the
polymer as determined in regards to the main DSC endothermic peak.
While such teaching is considered applicable to homogeneously branched
ethylene/a-olefin interpolymers having a single DSC melting peak, the
- prescribed range is fairly general and broad. Moreover, ~lideo et al do
not provide any specific teaching as to the optimum orientation
temperature for a particular interpolymer respecting heat shrink
response, nor any other desired film property.
An example of generalized teachings pertaining to
homogeneously branched ethylene/a-olefin interpolymers is provided in WO
~ 95/08441. In the Examples of this disclosure, several different
homogeneously branched substantially linear ethylene/a-olefin
interpolymers were studied and compared to one heterogeneously branched
ethylene/a-olefin interpolymers. Although the homogeneously branched
substantially linear ethyLene/a-olefin interpolymers had densities that
varied from about 0.896 to about 0.906 g/cc, all of the interpolymers
(including the heterogeneously branched linear ethylene/a-olefin
interpolymer, Attane~ 4203, supplied by The Dow Chemical Company, which
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had a density of 0.905 g/cc~ were oriented at essentially the same
orientation temperatures. Reported results in WO 95/08441 disclose
three general findings -- (l) at an equivalent polymer density,
substantially linear ethylene/a-olefin interpolymers and heterogeneously
branched linear ethylene/~-olefin interpolymers have essentially
equivalent shrink responses (compare Example 21 and Example 39 at pages
15-16), (2) shrink responses increase at lower aensities and constant
orientation temperatures, and (3) as orientation temperature increases,
orientation rates increase. Furthermore, careful study of the Examples
and unreported DSC melting point data for the interpolymers reported on
in WO 95/0844r indicate for the Examples disclosed in WO 95/08441 that,
at a given stretching rate and ratio, there is a preference for
orienting multilayer film structures at orientation temperatures above
the respective DSC melting point of the polymer employed as the shrink
control layer.
Other disclosures that set forth orientation information
regarding homogeneously branched ethylene polymers yet do not specify
orientation conditions relative to respective lowest stretch
temperatures include EP 0 600425Al to Babrowicz et al. and EP 0 587502
A2 to Babrowicz et al.
Accordingly, although there are general rules and general
disclosure as to suitable orientation temperatures for biaxially
orienting polyolefins, there is no specific information as to optimum
orientation conditions as a function of polymer type and, more
importantly, there is no specific information for homogeneously branched
interpolymers which do no~ possess the "lower and higher melting peaks"
- required by the Golike method. Also, while there is some fragmented
information regarding stretching at different orientation temperatures,
there is no specific information as to the m~;mll~ shrink response at
the lowest possible orientation temperature at a given stretching rate
and ratio for homogeneously branched interpolymers in general, and
especially no useful information for homogeneously branched
~ interpolymers having densities less than 0.91 g/cc. Still further,
although homogeneously branched ethylene interpolymers offer a variety
of other useful property advantages, at equivalent densities above about
0.907 g/cc, the shrink response of multilayer film structures which
contain a homogeneously branched ethylene interpolymer as the shrink
control layer, is generally viewed as essentially equivalent to
multilayer film structures which contain a heterogeneously branched
interpolymer as the shrink control layer. See FI~. 3.
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It is an object of the present invention to provide a method
for defining the optimum stretching or orientation temperature to
maximize the unrestrained (free) shrink response of polyolefins in
general. It is a particular object of the present invention to provide
a method for maximizing the shrink response of homogeneously branched
ethylene interpolymers such that the m~ m shrink potential of these
interpolymers can be obtained for the particular stretching rate,
stretching ratio and type of orientation equipment employed. Another
object of the invention is to provide improved shrink film structures
containing homogeneously branched ethylene interpolymers of densities
less than 0.9l g/cc as the shrink control layer.
In accordance with the present invention, we have discovered
that defining the lowest temperature for stretching or orienting
polyolefins results in a maximized shrink response, that such optimum
orientation temperature varies with polymer density and/or
cryst~lli ni ty, and that residual crystallinity determinations using
differential scanning calorimetry (DSC) indicate homogeneously branched
ethylene interpolymers have lower residual crystallinities at so-defined
optimum stretching or orientation temperatures relative to
heterogeneously branched ethylene interpolymers having equivalent
densities. As such, while such optimum orientation conditions can be
also determined by trial and error approaches, the DSC residual
crystallinity methodology provides a systematic way to efficiently
identify such conditions. The result of these discoveries is whereas
shrink response was previously thought to strictly follow polymer
density and to be independent of the polymer homogeneity, with this new
method of orientatiOn, at equivalent densities, homogeneously branched
interpolymers have been found to provide utterly unexpected,
dramatically superior shrink responses relative to heterogeneously
branched interpolymers, especially at polymer densities less than 0.9l
g/cc.
In particular, we have discovered a method of making a heat
- shrinkable polyolefin film comprising the steps of
a. fabricating a polyolefin film structure in substantially
unoriented form, and
b. stretching the polyolefin film at a selected stretçhing
rate, stretch ratio and stretching temperature, wherein
the selected stretching temperature is no more than or
equal to 5~C above the lowest stretch temperature for the
polyolefin film structure and for the selected stretching
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rate and stretching ratio and wherein the polyolefin film
structure comprises at least one ethylene polymer having
a polymer density less than 0.915 g/cc.
Another aspect of the present invention is a heat shrinkable
film structure which comprises, as the shrink control layer, at least
one homogeneously branched ethylene interpolymer having a polymer
density of less than O.91 g/cc, wherein the film structure is
characterized as having a shrink response at least lO percent greater
than the shrink response of a second film structure which comprises a
heterogeneously branched ethylene interpolymer as the shrink control
layer and wherein the ~ilm structure and the second film structure are
fabricated and stretched under essentially the same conditions and the
homogeneously branched and heterogeneously branched interpolymers have
essentially the same polymer density and I2 melt index.
Still another aspect of the present invention is a heat
shrink polyolefin film structure prepared by a method which comprises
the steps of
~a) fabricating a polyolefin film structure in
substantially unoriented form, and
(b) thereafter stretching the fabricated polyolefin
film structure at a selected stretching rate, stretch ratio
and stretching temperature, wherein the selected stretching
temperature is below the melting point of the film and is a
temperature no more than or equal to 5~C above the lowest
stretch temperature for the selected stretching rate and
stretching ratio, and
~ wherein the film structure comprises, as the shrink control layer, at
least one homogeneously branched ethylene interpolymer having a polymer
density of less than about O.9l g/cc.
While the present invention allows stretching operations
in general to maximize the unrestrained shrink potential of a
particular polymer, the benefits o~ this invention are particularly
useful for those common commercial instances where the orientation
temperature capabilities of the stretching operation are essentially
fixed. For a particular orientation temperature capability (and
stretch ratio and stretching rate), this invention allows the
systematic identification of the optimum interpolymer rather than
haphazard selections and miscalculations which waste time and can
lead to over-engineered, more costly interpolymers. Another benefit
of present invention is the novel method allows direct shrink
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response comparisons irrespective of polymer differences such as
density; that is, the novel method is a form of comparative
~ standardization that can facilitate the development of com~ercial
shrink films.
FIG. l is a first heat DSC curve illustrating the
residual crystallinity portion of a heterogeneously branched polymer
r~m~;nlng at 100~C which is a temperature below the various melting
peaks of the polymer illustrated.
FIG. 2 is a x/y plot illustrating the shrink response of
heterogeneously branched ethylene polymers and homogeneously branched
ethylene polymer as a function of polymer density. The data used to
generate the plot are reported in Table 2 hereinbelow. The
heterogeneously branched ethylene polymer samples range in polymer
densities from about 0.907 to about 0.932 g/cc while the
homogeneously branched ethylene polymer samples range from about O.9l
to about 0.918 g/cc.
FIG. 3 is another x/y plot illustrating the shrink
response of heterogeneously branched ethylene polymers and
homogeneously branched ethylene polymer as a function of polymer
density. The data used to generate this plot are reported in Table 2
hereinbelow. The heterogeneously branched ethylene polymer samples
range in polymer densities from about 0.907 to about 0.932 g/cc while
the homogeneously branched ethylene polymer samples range from about
0.887 to about 0.918 g/cc.
The double bubble and trapped bubble biaxial orientation
methods can be simulated on a laboratory scale using a T. M. Long
stretcher which is analogous to a tenter frame device. This device can
orient polyolefin films in both the monoaxial and biaxial mode at
stretching ratios up to at least 5:l. The device uses films having an
original dimension of 2 inches x 2 inches Biaxial stretching is
usually performed by stretching in the machine direction and transverse
direction of the film simultaneously, although the device can be
- operated to stretch se~uentially.
The residual crystallinity of polyolefin interpolymers
measured using a DSC partial area method can be used to characterize the
nature of polyolefin film at the stretch temperature. A stretching
temperature 5~C above, preferably 3~C above, more preferably 2.5~C above
the lowest stretch temperature (defined herein below) is considered
herein to be the optimum or near-optimum stretching or orientation
temperature for the particular film. Stretching temperatures higher
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than 5~C above the lowest stretch temperature are not considered part of
the present invention because such invariably yield lower shrink
responses for a particular stretching rate, stretching ratio and shrink
temperature. Stretching temperatures less than 2.5~C above the lowest
stretch temperature are not preferred because they tend to yield
inconsistent results, although such inconsistencies tend to depend on
specific equipment and temperature control capabilities.
We discovered that heterogeneously branched ethylene
interpolymers possess higher residual cryst~lli n; ties at their
respective optimum orientation temperature relative to homogeneously
branched ethylene interpolymers. Heterogeneously branched interpolymers
having a density in the range of from O.9O to 0.93 g/cc have residual
crystallinities at respective optimum orientation temperatures of from
20 to 24 percent, while homogeneously branched ethylene interpolymers
having a density in the range of from 0.895 to 0.91 g/cc have residual
crystallinities at their respective optimum orientation temperatures of
from 14 to 17 percent.
"Stretched" and "oriented" are used in the art and herein
interchangeably, although orientation is actually the consequence of a
film being stretched by, for example, internal air pressure pushing on
the tube or by a tenter irame pulling on the edges of the film.
The term "lowest stretch temperature" as used herein means
the temperature below which the film either tears and/or stretches
unevenly for a given stretching rate and stretching (draw) ratio during
2~ the stretching operation or step o~ an orientation technique. The
lowest stretch temperature is below the melting point of the film, it is
a temperature at or below which the film can not be stretched uniformly
(i.e., without the occurrence of banding or thick and thin spots), and
it is a temperature at or below which the film tears for a particular
stretching rate and stretch ratio.
Practitioners will appreciate that to m~X; m; ze the
orientation imparted and therefore the shrink response, the objective is
to operate as close to the lowest stretch temperature as their equipment
and capabilities will allow whether or not the significant stretching or
orientation is accomplished in a single step or by a combination of
sequential steps.
Additionally,-practitioners will appreciate that the optimum
or near-optimum stretching temperature for maximized shrink response at
a given shrink temperature will interrelate with stretching rate and
ratio. That is, while a particular stretching temperature will be
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.
optimum or near-optimum at one combination of stretching rate and ratio,
the same stretching temperature will not be optimum or near-optimum at a
different combination of stretching rate and ratio.
Practitioners will also appreciate that to obtain the
m~imll~ shrink response from the orientation frozen into the film, the
shrink temperature should match or exceed the stretching temperature.
That is, reduced shrink temperatures do not allow full relaxation or
shrinkage of the film. However, excessive shrink temperatures can
diminish film integrity.
Practitioners will further appreciate that for a given
combination of stretching temperature, stretching rate and stretching
ratio, increases in the shrink temperature to the point of film
integrity failure will yield higher shrink response performance and
higher levels of shrink tension.
Stretching temperatures in the range of from 50 to about
120~C, especially from 55 to 110~C, more especially from 60 to 95~C, and
most especially from 65 to 90~C are suitable in the present invention.
Shrink temperatures in the range of from 70 to 140~C,
especially from 80 to 125~C, and more especially from 85 to 100~C are
suitable in the present invention.
The term "residual crystallinity" is used herein to refer
the crystallinity of a polymer film at a particular stretching
temperature. Residual crystallinity is determined using a Perkin-Elmer
DSC 7 set for a first heat at 10~C/min. of a water-quenched, compression
molded ~ilm sample of the polymer. The residual crystallinity for an
interpolymer at a particular temperature is determined by measuring heat
of ~usion between that temperature and the temperature of complete
melting using a partial area technique and by dividing the heat of
fusion by 292 Joules/gram. The heat of fusion is determined by computer
integration of the partial area using Perkin-Elmer PC Series Software
Version 3.1. An example of the residual crystallinity determination and
calculation is shown in ~IG. 1.
- The term "shrink control layer" is used herein to refer to
the film layer that provides or controls the shrink response. Such a
layer is inherent to all heat shrink films. In a monolayer heat shrink
film, the shrink control layer will be the film itself. In a multilayer
heat shrink film, the shrink control layer is typically the core or an
inside film layer and is typically the thickest film layer. See, for
example, Wo 95/08441.
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The term "substantially unoriented form" is used herein in
reference the fact that some amount of orientation is usually imparted
to a film during ordinary fabrication. As such, it is meant that the
fabrication step, in itself, is not used to lmpart the degree of
orientation required for the desired or required shrink response. The
present invention is thought to be generally applicable to operations
where the fabrication and orientation steps are separable and occur
simultaneously. However, the present invention is preferably directed
to an additional and separate orientation step which is required beyond
the making of tube, sock, web or layflat sheet whether or not such is
soft, molten, or irradiated before substantial orientation is imparted.
The terms "homogeneous ethylene interpolymer,"
"homogeneously branched ethylene interpolymer" and "narrow short chain
distribution" are used in the conventional sense in reference to an
ethylene interpolymer in which the comonomer is randomly distributed
within a given polymer molecule and wherein substantially all of the
polymer molecules have the same ethylene to comonomer molar ratio. The
terms refer to an ethylene interpolymer that is characterized by a
relatively high short chain branching distribution index (SCBDI) or
composition distribution branching index (CDBI). That is, the
interpolymer has a SCBDI greater than or equal to 50 percent, preferably
greater than or equal to 70 percent, more preferably greater than or
equal to 90 percent and essentially lack a measurable high density
(crystalline) polymer fraction.
SCBDI or CDBI is defined as the weight percent of the
polymer molecules having a comonomer content within 50 percent of the
median total molar comonomer content and represents a comparison of the
monomer distribution in the interpolymer to the monomer distribution
expected for a Bernoullian distribution. The SCBDI of an interpolymer
can be readily calculated from data obtained from techniques known in
the art, such as, for example, temperature rising elution fractionation
(abbreviated herein as "TREF") as described, for example, by Wild et
- al., Journal of Polymer Science, Poly. Phys. Ed., Vol. 20, p. 441
(1982), or in US Patent g,798,08l, or by L. D. Cady, "The Role of
Comonomer Type and Distribution in LLDPE Product Performance," SPE
Regional Technical Conference, Quaker Square Hilton, Akron, Ohio,
October 1-2, pp. 107-ll9 (I985). However, the preferred TREF technique
does not include purge quantities in SCBDI calculations. More
preferably, the monomer distribution of the interpolymer and SCBDI are
determined using C NMR analysis in accordance with techniques described
--10--
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in US Patent 5,292,845 and by J. C. Randall in Rev. Macromol. Chem.
Phys., C29, pp. 201-317.
The terms "heterogeneous," "heterogeneously branched" and
"broad short chain distribution" are used herein in the conventional
sense in reference to a linear ethylene interpolymer having a
comparatively low short chain branching distribution index. That is,
the interpolymer has a relatively broad short chain branching
distribution. Heterogeneously branched linear ethylene interpolymers
have a SCBDI less than 50 percent and more typically less than 30
percen..
The term "homogeneously branched linear ethylene
interpolymer" means that the interpolymer has a homogeneous lor narrow)
short branching distribution but does not have long chain branching.
That is, the ethylene interpolymer has an absence of long chain
branching and a linear polymer backbone in the conventional sense of the
term "linear." Such interpolymers can be made using polymerization
processes (e.g., as described by Elston in USP 3,645,992) which provide
uniform (narrow) short branching distribution (i.e., homogeneously
branched). In his polymerization process, Elston uses soluble vanadium
catalyst systems to make such polymers, however others such as Mitsui
Petrochemical Corporation and Exxon Chemical Company have used so-called
single site catalyst systems to make polymers having a similar
homogeneous structure. Homogeneously branched linear ethylene
interpolymers can be prepared in solution, slurry or gas phase processes
using hafnium, zirconium and vanadium catalyst systems. Ewen et al. in
U.S. Pat. No. 4,937,299 describe a method of preparation using
metallocene catalysts.
The term "homogeneously branched linear ethylene
interpolymer" does not refer to high pressure branched polyethylene
which is known to those skilled in the art to have numerous long chain
branches.
Typically, the homogeneously branched linear ethylene
-- interpolymer is an ethylene/a-olefin interpolymer, wherein the a-olefin
is at least one C3-c20 a-olefin (e.g., l-propylene, l-butene, l-pentene,
4-methyl-l-pentene, l-hexene, l-octene and the like), preferably wherein
at least one of the a-olefins is l-octene. Most preferably, the
ethylene/a-olefin interpolymer is a copolymer of ethylene and a C3-C20 a-
olefin, especially an ethylene/C4-C6 a-olefin copolymer. Commercial
examples of homogeneously branched linear ethylene/a-olefin
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interpolymers are sold by Mitsui Chemical under the designation "TAFMER"
and by Exxon Chemical under the designation "EXACT".
Heterogeneously branched VLDPE and LLDPE are well known
among practitioners of the linear polyethylene art. They are prepared
using conventional Ziegler-Natta solution, slurry or gas phase
polymeri~ation processes and coordination metal catalysts as described,
for example, by Anderson et al. in U.S. Pat. No. 4,076,698. These
conventional Ziegler-type linear polyethylenes are not homogeneously
branched, do not have any long-chain branching and have a linear polymer
backbone in the conventional sense of the term "linear." Also, these
polymers do not show any substantial amorphism at lower densities since
they inherently posses a substantial high density (crystalline) polymer
fraction. At densities less than 0.90 g/cc, these materials are more
difficult to prepare than homogeneously branched ethylene polymer and
are also more difficult to pelletize than their higher density
counterparts. At such lower densities, heterogeneously branched VLDPE
pellets are more tacky and have a greater tendency to clump together
than their higher density counterparts.
The terms "ultra low density polyethylene" (ULDPE), "very
2~ low density polyethylene" (VLDPE) and "linear very low density
polyethylene" (LVLDPE) have been used interchangeably in the
polyethylene art to designate the polymer subset of linear low density
polyethylenes having a density less than or equal to 0.915 g/cc. The
term "linear low density polyethylene" (LLDPE) is then applied to those
linear polyethylenes having a density above 0.915 g/cc. As used herein
and in the conventional sense, these terms indicate that the polymer has
a heterogeneous short chain branching distribution and linear polymer
backbone. Commercial examples of heterogeneously branched VLDPE
polyolefins suitable for use in the present invention include ATTANE~
~LDPE polymers supplied by the Dow Chemical Company and FLEXOMER~ VLDPE
polymers supplied by ~nion Carbide Corporation.
Although the novel method of the present invention is useful
- for preparing shrink film structures comprised of heterogeneously
branched ethylene polymers, and homogeneously branched ethylene polymers
and these polymers are also suitable for claimed shrink film prepared by
the novel method, not all of the above polymers are suitable for use in
the novel film of the present invention. That is, while the method and
the shrink film by the method are generally applicable to all of the
above ethylene polymers, only homogeneously branched substantially
linear ethylene polymers and homogeneously branched linear ethylene
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polymers are suitable for all aspects of the present invention including
the novel film.
~ The term "substantially linear ethylene/a-olefin
interpolymer" is used herein to refer to homogeneously branched
ethylene/a-olefin interpolymers that contain long chain branches as well
as short chain branches attributable to homogeneous comonomer
incorporation. The long chain branches are of the same structure as the
backbone of the polymer and are longer than the short chain branches.
The polymer backbone of substantially linear a-olefin polymers is
10 substituted with an average of 0.01 to 3 long chain branch/1000 carbons.
Prefer~ed substantially linear polymers for use in the invention are
substituted with from 0.01 long chain branch/1000 carbons to 1 long
chain branch/1000 carbons, and more preferably from 0.05 long chain
branch/1000 carbons to 1 long chain branches/1000 carbons.
Long chain branching is defined herein as a chain length of
at least 6 carbons, above which the length cannot be distinguished using
13C nuclear magnetic resonance spectroscopy. The long chain branch can
be as long as about the same length as the length of the polymer
backbone to which it is attached= To~g shain bra~Ghe_ are obvLQ~sly of
greater length than of short chain branches resulting from comonomer
incorporation.
The presence of long chain branching can be determined in
ethylene homopolymers ~y using 13C nuclear magnetic resonance (NMR)
spectroscopy and is quantified using the method described by Randall
25 (Rev. Macromol. Chem. Phys., C29, V. 2&3, p. 285-297).
As a practical matter, current 13C nuclear magnetic
resonance spectroscopy cannot determine the length of a long chain
branch in excess of six carbon atoms. However, there are other known
techniques useful for det~r~in;ng the presence of long chain branches in
ethylene polymers, including ethylene/1-octene interpolymers. Two such
methods are gel permeation chromatography coupled with a low angle laser
light scattering detector (GPC-LALLS) and gel permeation chromatography
coupled with a differential viscometer detector (GPC-DV). The use of
these techniques for long chain branch detection and the underlying
theories have been well documented in the literature. See, for example,
Zimm, G.H. and Stockmayer, W.H., J. Chem. Phys., 17, 1301 (1949) and
Rudin, A., Modern Methods of Polymer Characterization, John Wiley &
Sons, ~ew York (1991) pp. 103-112.
A. Willem deGroot and P. Steve Chum, both of The Dow
40 Chemical Company, at the October 4, 1994 conference of the Federation of
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Analytical Chemistry and Spectroscopy Society (FACSS) in St. Louis,
Missouri, presented data demonstrating that GPC-DV is a useful technique
for quantifying the presence of long chain branches in substantially
linear ethylene interpolymers. In particular, deGroot and Chum found
that the level of long chain branches in substantially linear ethylene
homopolymer samples measured using the Zimm-Stockmayer equation
correlated well with the level of long chain branches measured using l3C
NMR.
Further, deGroot and Chum found that the presence of octene
does not change the hydrodynamic volume of the polyethylene samples in
solution and, as such, one can account for the molecular weight increase
attributable to octene short chain branches by knowing the mole percent
octene in the sample. By deconvoluting the contribution to molecular
weight increase attributable to l-octene short chain branches, deGroot
and Chum showed that GPC-DV may be used to quantify the level of long
chain branches in substantially linear ethylene/octene copolymers.
deGroot and Chum also showed that a plot o~ Log(I2, Melt
Index) as a function of Log(GPC Weight Average Molecular Weight) as
determined by GPC-~V illustrates that the long chain branching aspects
(but not the extent of long branching) of substantially linear ethylene
polymers are comparable to that of high pressure, highly branched low
density polyethylene (L~PE) and are clearly distinct from ethylene
polymers produced using Ziegler-type catalysts such as titanium
complexes and ordinary homogeneous catalysts such as hafnium and
vanadium complexes.
The substantially linear ethylene/~-olefin interpolymers
used in the present invention are-a unique class of ~ompounds that are
further defined in US Patent 5,272,236, serial number 07/776,130 filed
October 15, l99l and in US patent 5,278,272, serial number 07/939,281
filed September 2, 1992.
Substantially linear ethylene/a-olefin interpolymers differ
significantly from the class of polymers conventionally known as
homogeneously branched linear ethylene/a-olefin interpolymers described,
for example, by ~lston in US Patent 3,645,992, in that substantially
linear ethylene interpolymers do not have a linear polymer backbone in
the conventional sense of the term "linear." Substantially linea~
ethylene/a-olefin interpolymers also differ significantly from the class
of polymers known conventionally as heterogeneously branched traditional
Ziegler polymerized linear ethylene interpolymers (for example, ultra
low density polyethylene, linear low density polyethylene or high
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density polyethylene made, for example, using the technique disclosed by
Anderson et al. in US Patent 4,076,698 and utilized by Golike as
described in US Patent 4,597,920, in that substantially linear ethylene
interpolymers are homogeneously branched interpolymers. Substantially
linear ethylene/a-olefin interpolymers also differ significantly from
the class known as free-radical initiated highly branched high pressure
low density ethylene homopolymer and ethylene interpolymers such as, for
example, ethylene-acrylic acid (EAA) copolymers and ethylene-vinyl
acetate (EVA) copolymers, in that substantially linear ethylene
interpolymers do not have equivalent degrees of long chain branching and
are made using single site catalyst systems rather than free-radical
peroxide catalysts systems.
Single site polymerization catalyst, (for example, the
monocyclo-pentadienyl transition metal olefin polymerization catalysts
described by Canich in US Patent 5,026,798 or by Canich in US Patent
5,055,438) or constrained geometry catalysts (for example, as described
by Stevens et al in US Patent 5,064,802) can be used to prepare
substantially linear ethylene polymers, so long as the catalysts are
used consistent with the methods described in US Patent 5,272,236 and in
US Patent 5,278,272 Such polymerization methods are also described in
PCT/US 92/08812 (filed October 15, 1992) However, the substantially
linear ethylene interpolymers are preferably made by using suitable
constrained geometry catalysts, especially constrained geometry
catalysts as disclosed in US Application Serial Nos : 545,403, filed
July 3, 1990; 758,654, filed September 12, 1991; 758,660, filed
September 12, 1991; and 720,041, filed June 24, 1991
- Suitable cocataLysts for use herein include but are not
limited to, for example, polymeric or oligomeric aluminoxanes,
especially methyl aluminoxane or modified methyl aluminoxane (made, for
example, as described in US Patent 5,041,584, US Patent 4,544,762, US
Patent 5,015,749, and/or US Patent 5,041,585) as well as inert,
compatible, non-coordinating, ion forming compounds. Preferred
cocatalysts are inert, non-coordinating, boron compounds.
The polymerization conditions for manufacturing the
substantially linear ethylene interpolymers used in the present
invention are preferably those useful in the continuous solution
polymerization process, although the application of the present
invention is not limited thereto Continuous slurry and gas phase
polymerization processes can also be used, provided the proper catalysts
and polymerization conditions are employed. To polymerize the
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substantially linear interpolymers useful in the invention, the single
site and constrained geometry catalysts mentioned earlier can be used,
but for substantially linear ethylene polymers the polymerization
process should be operated such that the substantially linear ethylene
interpolymers are formed. That is, not all polymerization conditions
inherently make the substantially linear ethylene polymers, even when
the same catalysts are used. For example, in one embodiment of a
polymerization process useful in making substantially linear ethylene
interpolymers, a continuous process is used, as opposed to a batch
process.
The substantially linear ethylene interpolymer ~or use in
the present invention is characterized as having
(a) a melt flow ratio, Ilo/I2 2 5.63,
(b) a molecular weight distribution, MW/Mn, as determined by gel
permeation chromatography and defined by the equation:
(MW/Mn) S (I10/I2) - 4.63,
(c) a gas extrusion rheoloqy such that the critical shear rate
at onset o~ surface melt ~racture for the substantially linear ethylene
interpolymer is at least 50 percent greater than the critical shear rate
at the onset of sur~ace melt fracture for a linear ethylene
interpolymer, ~herein the substantially linear ethylene interpolymer and
the linear ethylene interpolymer comprise the same comonomer or
comonomers, the linear ethylene interpolymer has an I2 MW/Mn and
density within ten percent o~ the substantially linear ethylene
interpolymer and wherein the respective critical shear rates o~ the
substantially linear ethylene interpolymer and the linear ethylene
interpolymer are measured at the same melt temperature using a gas
extrusion rheometer, and
(d) a single differential scanning calorimetry, DSC, melting
30 peak between -30 and 150~C.
The substantially linear ethylene interpolymers used in this
_ invention are homogeneously branched interpolymers and essentially lack
a measurable "high density" ~raction as measured by the TREF technique
(i.e., have a narrow short chain distribution and a high SCBD index).
The substantiaLly linear ethylene interpolymer generally do not contain
a polymer fraction with a degree of branching less than or equal to 2
methyls/lOOO carbons. The "high density polymer ~raction" can also be
described as a polymer fraction with a degree o~ branching less than 2
methyls/lO00 carbons.
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The substantially linear ethylene interpolymers for use in
the novel method and the film made from the novel method of the present
invention are interpolymers of ethylene with at least one C3-C20 a-
olefin and/or C4-Clg diolefin. Copolymers of ethylene and an a-olefin
of C3-C20 carbon atoms are especially preferred. The term
"interpolymer" is used herein to indicate a copolymer, or a terpolymer,
or the like, where, at least one other comonomer is polymerized with
ethylene to make the interpolymer.
Suitable unsaturated comonomers useful for polymerizing with
ethylene include, for example, ethylenically unsaturated monomers,
conjugated or non-conjugated dienes, polyenes, etc. Examples of such
comonomers include C3-C20 a-olefins as propylene, isobutylene, l-butene,
l-hexene, 4-methyl-l-pentene, l-heptene, l-octene, l-nonene, l-decene,
and the like. Preferred comonomers include propylene, l-butene, l-
hexene, 4-methyl-l-pentene and l-octene, and l-octene is especially
preferred. Other suitable comonomers include styrene, halo- or alkyl-
substituted styrenes, tetrafluoroethylene, vinylbenzocyclobutane, l,4-
hexadiene, l,7-octadiene, and cycloalkenes, e.g., cyclopentene,
cyclohexene and cyclooctene.
Determination of the critical shear rate and critical shear
stress in regards to melt fracture as well as other rheology properties
such as "rheological processing index" (PI), is performed using a gas
extrusion rheometer (GER~. The gas extrusion rheometer is described by
M. Shida, R.N. Shroff and ~.V. Cancio in Polymer Engineering Science,
25 Vol. 17, No. ll, p. 770 (1977), and in "Rheometers for Molten Plastics"
by John Dealy, published by Van Nostrand Reinhold Co. ~l982) on pp. 97-
99. GER experiments are performed at a temperature of about 190~C, at
nitrogen pressures between about 25Q to about 5500 psig using about a
0.0754 mm diameter, 20:l L/D die with an entrance angle of about 180~.
For the substantially linear ethylene polymers described herein, the PI
is the apparent viscosity (in kpoise) of a material measured by GER at
- - an apparent shear stress of about 2.15 x lQ6 dyne/cm2. The substantially linear ethylene polymer for use in the invention are ethylene
interpolymers having a PI in the range of O.Ol kpoise to 50 kpoise,
preferably 15 kpoise or less. The substantially linear ethylene
polymers used herein have a PI less than or equal to 70 percent of the
PI of a linear ethylene interpolymer (either a conventional Ziegler
polymerized interpolymer or a linear homogeneously branched interpolymer
as described by Elston in US Patent 3,645,992) having an I2, MW/Mn and
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density, each within ten percent of the substantially linear ethylene
interpolymer.
An apparent shear stress versus apparent shear rate plot is
used to identify the melt fracture phenomena and quantify the critical
shear rate and critical shear stress of ethylene polymers. According to
R~m~ rthy in the Journal cf Rheology, 30(2), 337-357, 1986, above a
certain critical flow rate, the observed extrudate irregularities may be
broadly classified into two main types: surface melt fracture and gross
melt fracture.
Surface melt fracture occurs under apparently steady flow
conditions and ranges in detail frDm loss of specular film gloss to the
more severe form of "sharkskin." ~erein, as determined using the above-
described GER, the onset of surface melt fracture (OSMF) is
characterized at the beginning of losing extrudate gloss at which the
surface roughness of the extrudate can only be detected by 40x
magnification. The critical shear rate at the onset of surface melt
fracture for the substantially linear ethylene interpolymers is at least
50 percent greater than the critical shear rate at the onset of surface
melt ~racture of a linear ethylene interpolymer having essentially the
same I2 and Mw/Mn.
Gross melt fracture occurs at unsteady extrusion flow
conditions and ranges in detail from regular (alternating rough and
smooth, helical, etc.) to random distortions. For commercial
acceptability and maximum abuse properties of films, coatings and
profiles, surface defects should be minimal, if not absent The
critical shear stress at the onset o~ gross ~elt fracture for the
- substantially linear ethylene interpolymers used in the invention, that
is those having a density less than 0.91 g/cc, is greater than 4 x 1O6
dynes/cm2. The critical shear rate at the onset of surface melt
fracture (OSMF) and the onset of gross melt fracture (OGMF) will be used
herein based on the changes of surface roughness and configurations of
the extrudates extruded by a GER. Preferably, in the present invention,
the substantially linear ethylene interpolymer will be characterized by
its critical shear rate, rather than its critical shear stress.
Substantially linear ethylene/~-olefin interpolymers, like
other homogeneously branched ethylene/~-olefin interpolymers that
consist of a single polymer component material, are characterized by a
single DSC melting peak. The single meltinq peak is determined using a
differential scanning calorimeter standardized with indium and deionized
water. The method involves about 5-7 mg sample sizes, a "first heat" to
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about 180~C which is held for about 4 minutes, a cool down at about
10~/min. to about -30~C which is held for about 3 minutes, and heat up
at about 10~C/min. to about 150~C for the "second heat". The single
melting peak is taken from the "second heat" heat flow vs. temperature
curve. Total heat of fusion of the polymer is calculated from the area
under the curve.
For substantially linear ethylene interpolymers having a
density of 0.875 g/cc to 0.91 g/cc, the single melting peak may show,
depending on equipment sensitivity, a "shoulder" or a "hump" on the low
melting side that constitutes less than 12 percent, typically, less than
9 percent, and more typically less than 6 percent of the total heat of
fusion of the polymer. Such an artifact is observable for other
homogeneously branched polymers such as Exact~ resins and is discerne&
on the basis of the slope of the single melting peak varying
monotonically through the melting region of the artifact. Such an
artifact occurs within 34~C, typically within 27~C, and more typically
within 20~C of the melting point of the single melting peak The heat
of fusion attributable to an artifact can be separately determined by
specific integration of its associated area under the heat flow vs.
temperature curve.
The molecular weight determination is deduced by using narrow
molecular weight distribution polystyrene standards (from Polymer
Laboratories) in conjunction with their elution volumes. The equivalent
polyethylene molecular weights are determined by using appropriate Mark-
Houwink coefficients for polyethylene and polystyrene (as described byWilliams and Ward in Journal of Polymer Science, Polymer Letters, Vol. 6,
(621) 1968) to derive the following equation:
Mpolye~ylene = a * (Mpolys~rene) .
In this equation, a = 0.4316 and b = 1Ø Weight average molecular
weight, Mw, and number average molecular weight, Mn~ is calculated in the
usual manner according to the ~ollowing formula:
M~ wi(Mi~ ; where wi is the weight fraction of the molecules with
molecular weight Mi eluting from the GPC column in fraction i and j = 1
when calculating Mw and j = -1 when calculating Mn~
For the homogeneously branched ethylene interpolymers used
in the present invention, the MW/Mn is preferably less than 3, mo~e
preferably less than 2.5, and especially from 1.5 to about 2.5 and most
especially from 1.8 to 2.3.
Substantially linear ethylene interpolymers are known to
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have excellent processability, despite having a relatively narrow
molecular weight distribution. Surprisingly, unlike homogeneously and
heterogeneously branched linear ethylene polymers, the melt flow ratio
(I1o/I2~ of substantially linear ethylene interpolymers can be varied
essentially independently of the molecular weight distribution, MW/Mn~
Accordingly, the preferred ethylene a-olefin interpolymer for use in the
present invention is a substantially linear ethylene interpolymer.
Homogeneously branched substantially linear ethylene
interpolymers are available from The Dow Chemical Company as Affinity
polyolefin plastomers, and as Engage polyolefin elastomers.
Homogeneously branched substantially linear ethylene polymers can be
prepared by the continuous solution, slurry, or gas phase polymerization
of ethylene and one or more optional a-olefin comonomers in the presence
of a constrained geometry catalyst, such as is disclosed in European
Patent Application 416,815-A.
The density of the polyolefin polymer (as measured in
accordance with ASTM D-792) for use in the présently claimed method is
generally greater than 0.85 g/cc, especially ~rom 0.86 g/cc to 0.93
g/cc, more preferably, from about 0.88 g/cc to 0.92 g/cc and most
2C preferably, from 0.88 to 0.91. ~hen used as the shrink control polymer
layer of the shrink film, the preferred polymer density of the
polyolefin polymer is less than 0.915 g/cc. The density of the
homogeneously branched ethylene polymer for use in all aspects o~ the
present invention is less than 0.91 g/cc, generally in the range of 0.85
to 0.91 g/cc, preferably less than 0.907 g/cc, more- preferably less than
or equal to 0.905 g/cc, most preferably less than or equal to 0.902
g/cc, and especially in the range of 0 880 to 0.90 g/cc.
The molecular weight of polyolefin polymers is conveniently
indicated using a melt index measurement according to ASTM D-1238,
Condition 190~C/2.16 kg (formerly known as "Condition E" and also known
as I2). Melt index is inversely proportional to the molecular weight of
- the polymer. Thus, the higher the molecular weight, the lower the melt
index, although the relationship is not linear. The melt index for the
polyolefin polymers use~ul herein is generally from 0.01 g/10 min. to 20
g/10 min., preferably from 0.01 g/10 min. to 10 g/10 min., and
especially from 0.1 g/10 min. to 2 g/10 min.
Other measurements useful in characterizing the molecular
weight of substantially linear ethylene interpolymers and homopolymers
involve melt index determinations with higher weights, such as, for
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r.h h { ~
. ' r ~ C --r r r ~ r r ~. r ~ ~ -
common example, ASTM D-1238, Condition 190~C/10 kg (formerly known as
"Condition N" and also known as I1o). The ratic of a higher weight melt
index determination to a lower weight determination is known as a melt flow
ratio, and for measured I1o and the I2 melt index values the melt flow
ratio is conveniently designated as I10/I2. For the substantially linear
ethylene polymers used to prepare the films of the present invention, the
melt flow ratio indicates the degree of long chain branching, i.e., the
higher the I1o/I2 melt flow ratio, the more long chain branching in the
polymer. The I1o/I2 ratio of the substantially linear ethylene polymers is
preferably at least 7, and especially at least 9.
Additives such as antioxidants (e.g., hindered phenolics (such
as Irganox~ 1010 or Irganox~ 1076~, phosphites (e.g., Irgafos~ 168), cling
additives (e.g., PIB), PEPQ~ (a trademark of Sandoz Chemical, the primary
ingredient of which is believed to be a biphenylphosphonite), pigments,
colorants, fillers, and the like can also be included in the polyolefin
polymers, to the extent that they do not interfere with the method and the
enhanced shrink response discovered by Applicants. The fabricated film may
also contain additives to enhance its antiblocking and coefficient of
friction characteristics including, but not limited to, untreated and
treated silicon dioxide, talc, calcium carbonate, and clay, as well as
primary and secondary fatty acid amides, silicone coatings, etc. Other
additives to enhance the film's anti-fogging characteristics may also be
added, as described, for example, in US Patent 4,486,552 (Niemann). Still
other additives, such as quaternary ammonium compounds alone or in
combination with EAA or other functional polymers, may also be added to
enhance the film's antistatic characteristics and allow packaging of
electronically sensitive goods.
Film structures of the present invention can be made using
conventional simple bubble or cast extrusion techniques, however, preferred
film structures are prepared using more elaborate techniques such as
"tenter framing" or the "double bubble," "tape bubble" or "trapped bubble"
process. The double bubble technique is described by Pahkle in US Patent
3,456,044.
Multilayer film structures of the invention can be prepared by
a coextrusion technique or a lamination technique, and can also comprises a
polymer mixture. Suitable polymer mixtures include at least one
homogeneously branched ethylene interpolymer such as the polymer mixture of
at least one homogeneously branched substantially linear ethylene
interpolymer and at least one heterogeneously branched ethylene polymer.
Further, multilayer film structures of the invention can also comprises a
barrier film layer.
The presently claimed method for preparing shrink film
structures and the novel film structures of the present invention are more
fully described in the following examples, but are not limited to the
examples shown. The homogeneously branched substantially linear
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ethylene polymers used in the following examples were prepared according
to procedures and techniques described in the Examples of U.S. Patent
5,272,236 and 5,278,272. The homogeneously branched linear ethylene
interpolymer used in the following Examples was made by the Exxon
Chemical Company.
Examples
In an experi~ent to determine the comparative shrink
response of ethylene interpolymers, a homogeneously branched
substantially linear ethylene interpolymer and a heterogeneously
branched linear ethylene interpolymer were evaluated. Example l
utilized a substantially linear ethylene/l-octene copolymer having a
density of 0.90 g/cc, a melt index (I~) of 0.8 g/lO minutes, a molecular
weight distribution (Mw/Mn) of 2.2, and a melt flow ratio (Ilo/I2) of
8.5. Example 2 utilized the polymer of Example l which had been
irradiated at 5 0 Mrad. Comparative Run 3 utilized a heterogeneously
branched ~LDPE ethylene/1-octene copolymer having a density of 0.905
g/cc, a melt index (I2) of 0.8, a molecular weight distribution (MW/Mn)
of 3.5 and a melt flow ratio (Ilo/I2) of 8. Comparative Run 4 utilized
the polymer of Comparative ~un 3 which had been irradiated at 5.~ Mrad.
Irradiation was performed on pellets of the respective interpolymers by
exposure to electron beam radiation at E Beam Services, Inc.
~Canterbury, NJ~. The DSC melting point for the non-irradiated
interpolymers was determined, all four samples were prepared into
sheeting (18.5 + 1.5 mil thick) (0.~7 mm i 0.04 mm) and subsequently
biaxially stretched using a T.M. ~ong laboratory stretching frame. The
stretching temperature utilized was a temperature below the DSC melting
point of the copolymer but 5~C above the temperature at which tearing of
the sheet occurs during stretching. These stretched sheets were tested
for unrestrained (free~ shrink at 95~C in accordance with procedures in
ASTM D-2732 by cutting four inch by four inch (lO.2 cm x lO.2 cm)
samples from each of the stretched sheets and carefully placing them
flat into the bottom of silicone-coated metal pans. The metal pans had
sides 1 inch (2.5 cm) high and were well-coated with 200 centipoise
silicone oil. The pans containing the film samples were then placed
into a forced-air convection oven at 95 C for ten minutes. After ten
minutes, the pans were removed from the oven and allowed to cool to an
ambient temperature. After cooling, the film samples were removed and
the dimensions in both the machine and transverse directions were
measured. The Vicat softening point was measured in accordance with ASTM
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Dl525. Table l summarizes the shrink response data as well as providesthe stretch ratio information for each sample:
Table l
SampleStretching Stretch Free Free
Temperature* Ratio Shrink Shrink
(C) (%) MD (~) TD
Inventive 88 3 x 3 > 35 > 35
Example l
Inventive 88 4 x 9 > 35 > 35
Example 2
Comparative Run 97 4 x 4 < 25 < 25
Comparative Run 97 4 x 4 < 25 < 25
*Stretchin~ temperature at + l~C
The data in Table l show that the sheets fabricated from the
homogeneously branched substantially linear ethylene interpolymer
exhibited superior shrink response performance (at least 14 percent
greater) over comparable sheets fabricated from the conventional
heterogeneously branched linear ethylene interpolymer. The superior
shrink response is exhibited even when the amount of biaxial stretching
is significantly lower for the Inventive Examples relative to the
comparative examples (i.e., 3 x 3 versus 4 x 4). The Inventive Examples
even show superior free shrink performance regardless of whether the
~ comparative examples were irradiated or nonirradiated prior to
orientation (stretching). The stretching temperature for the
homogeneously branched substantially linear ethylene interpolymer was
7~C below its single DSC melting point. Conversely, the stretching
temperature for the heterogeneously branched linear ethylene
interpolymer was 25~C below its highest DSC melting peak, 21~C below its
intermediate melting peak and 2~C above its lowest melting peak. The
stretching temperature for the homogeneously branched substantially
linear ethylene polymer is considered the optimum or near-optimum
stretching temperature for the polymer and for the particular stretch
ratio and stretching rate. That is, the shrink response is the m~;mllm
obtainable for the sample at a 95 C shrink temperature wherein a higher
stretching temperature would yield a reduced shrink response.
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In another evaluation, rather than utilizing the method
disclosed by Golike where ethylene polymer compositions are arbitrarily
biaxially oriented at a temperature within 10~C o~ the highest melting
polymer compound (as per polymer blends) or compositional ~raction (as
per a heterogeneously branched ethylene interpolymer), the percent
residual crystallinity as a function of temperature was determined for a
series of ethylene polymers. The series included heterogeneously
branched linear low density polyethylenes (LLDPEs) supplied by The Dow
Chemical Company under the trademark "DOWLEX;" heterogeneously branched
ultra or very low density polyethylenes (ULDPEs or VLDPEs) supplied by
The Dow Chemical Company under the trademark "ATTANE" and by Union
Carbide under the trademark "FLEXOMER;" homogeneously branched
substantially linear ethylene interpolymers supplied by The Dow Chemical
Company under the trademarks "A~INITY" and "ENGAGE;'' homogeneously
branched linear ethylene interpolymers supplied by Exxon Chemical
Corporation under the trademark "EXACT" and Mitsui Chemical Company
under the trademark "TAFMER;" and ethylene vinyl acetate (EVA)
copolymers supplied by Dupont Chemical Company under the trade~ark
"ELVAX" and Nova Polymers under the trademark "NOVAPOL."
The polymer samples were made into 30 mil sheets on a
standard cast ~ilm extrusion line by cast extrusion and quick quenching
with a chill roll. The melt temperature of the cast extrusion at the
die was set at 4~0~F (249~C) and the chill roll temperature was set at
75~F (24~C). DSC first heat at 10~C/min. was determined on tap water-
quenched, compression molded thin films of each polymer sample
(unextruded) to simulate the quenching encountered by the extrusion cast
- sheets. The temperature that corresponded to 22.5 weight percent
absolute residual crystallinity was determined for each polymer sample.
The evaluation proceeded by initially setting the stretching
temperature in the T M. Long stretcher such that 22.5 weight percent
residual crystallinity was maintained ~or each polymer sample. The so-
defined stretching temperature was substantially less than the
~ temperature o~ the lowest melting peak of the sample. If a given sample
sheet could not oriented (i.e., tore during stretching and/or stretched
unevenly), the stretching temperature was raised for subsequent sample
sheets of the same polymer sample in 3~C intervals until or such that a
corresponding sample sheet~could be consistently and uniformly oriented
at a stretching ratio o~ 4.5 x 4.5 and a stretching rate 5 inches/second
(12.7 cm/second). The ~irst interval of a higher temperature where the
polymer sample could be consistently and uniformly oriented was taken as
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the optimum stretching temperature for the particular polymer sample,
stretch ratio and stretching rate and, as such, provides the highest
residual crystallinity that the particular polymer sample sheet could be
oriented.
Tafmer~ A4090 could not be uniformly stretched apparently
due its relatively hlgh melt index. Similarly, Tafmer~ P0480 was not
oriented due to its very low density. The optimum orientation
temperatures for all of the polymer samples evaluated as well as the DSC
melting points, the residual crystallinity at the optimum stretching
temperature for tap water quenched film samples and the Vicat softening
temperatures for the various polymer samples are reported in Table 2.
With the exception of Example 9, all the stretching temperatures
reported in Table 2 were 3~C above the respective lowest stretch
temperature for each polymer sample. The stretching temperature
reported in Table 2 for Example 9 is more than 3~C above its lowest
stretch temperature.
For the stretching step, initial sample dimensions of 2" x
2" (5.1 cm x 5.l cm) were used. A stretching (draw) ratio of 4.5 x 4.5
and a stretching rate of 5 inches/ second (12 7 cm/s) were used. The
samples were preheated to the identified optimum or near-optimum
stretching temperature in the T. M. Long stretcher for 3 minutes. The
hot air was deflected so as not to impinge on the sample directly to
avoid hot spots on the sheets. Sheets were stretched at the highest
possible level of residual crystallinity ~i.e., at their respective
optimum or near-optimum stretching temperatures) to maximize the shrink
response potential of the sheets for the above stretching ratio and
- stretching rate.
The hot-water shrinkage at 90~C for the biaxially oriented
sheets are also shown in Table 2. Shrinkage values were obtained by
measuring the unrestrained shrink in a water bath kept at 90~C. The
samples were cut into 12 cm x 1.27 cm pieces. The samples were marked
10 cm. from one end for identification. Each sample was completely
~~ immersed in the water bath for five seconds and then removed. Film
shrinkage was obtained from the calculation using ASTM method D 2732-83.
The average of four samples was calculated and the data are also
reported in Table 2. Since the samples were equi-biaxially oriented
(i.e., 4.5 x 4.5), the shrinkage in the machine and transverse
directions were equivalent as expected. Also, the shrink response at
90C for the various heterogeneously branched and homogeneously branched
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ethylene interpolymers as a function of polymer density are shown in
FI~. 2 and FIG. 3.
The orientation (stretching) temperatures shown in Table 2
represent the low end of the orientation window for each sample. The
high end of an orientation window is usually just below the higher
meltir.g peak of the polymer. Thus, it can be concluded from Table 2
that heterogeneously branched ethylene interpolymers have a much broader
orientation window than homogeneously branched ethylene interpolymers
(i.e., Af~inity, Engage and Exact resins).
Golike in US Patent 4,597,920 teaches orientation should be
carried out between the lower and higher melting points of a
heterogeneously branched copolymer or polymer blend. The DSC melting
information for Dowlex~ LLDPE 2045, Attane~ ULDPE 4201, Attane~ ULDPE
4203 and Affinity~ substantially linear interpolymer PL 1880 are also
provided Table 2. Table 2 indicates, contrary to the teachings of
Golike, that the heterogeneously branched Dowlex~ and Attane~ polymers
can be oriented for a maximized shrink response at stretching
temperatures below their respective lower melting peaks. As set forth
above, the homogeneously branched Affinity~ polymer sample has a single
DSC melting point and, as such, Golike's teachings are not specifically
applicable to such polymers. However, it is noted that the
homogeneously branched Affinity~ polymer sample can also be oriented for
a maximized shrink response at a stretching temperature below its
respective lower melting peak.
Moreover, Table 2 indicates heterageneously branched LLDPE
and ULDPE polymers having a density in the range from O 907 g/cc to
0.937 q/cc can be oriented at a maximum residual crystallinity of from
20 weight percent to 24 weight percent and that the m~; mllm residual
crystallinity for optimum or near-optimum orientation of these polymers
is predominately influenced by polymer crystallinity or crystalline
polymer fractions. Table 2 also indicates homogeneously branched
- ethylene interpolymers in the density range from 0.899 g/cc to 0.918
g/cc can be oriented at a m~; mllm residual crystallinity of from 14
weight percent to 17 weight percent. These m~;mllm residual
crystallinity differences indicate that, at least as to shrink response,
heterogeneously branched ethylene interpolymers differ completely~from
homogeneously branched ethylene interpolymers.
Additionally, whereas FIG. 2 indicates that at polymer
densities greater than O.9l g/cc, heterogeneously branched polymers show
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a higher shrink response than homogeneously branched polymers at an
equivalent density, FIG. 3 indicates that at densities less than O.9l
g/cc for interpolymers stretched in accordance with the present
invention, homogeneously branched ethylene polymers show greater than or
equal to lO percent greater, especially 15 percent greater, more
especially 20 percent greater and most especially 25 percent greater
shrink response than heterogeneously branched ethylene polymers having
equivalent densities in that range and fabricated (including
orientation) at essentially the same conditions. Such is particularly
true where the interpolymer densities are less than 0.907 g/cc, more
especially equal to or less than 0.905 g/cc, more especially less than
or equal to 0.902 g/cc and most especially equal to or less than O.90
g/cc.
The shrink response of the homogeneously branched ethylene
polymers as shown in FIG. 3 is especially surprising since reasonable
extrapolation of the data shown in FIG. 2 as well as the data disclosed
in WO Publication 95/08441 suggest the shrink response of homogeneously
branched ethylene polymers are expected to be inferior to or, at best,
essentially equivalent to heterogeneously branched ethylene polymers at
densities less than O.9l g/cc.
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W O 97/30111 PCT~US97/02S85
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