A Study of Nano-Mechanical and Macro-MechanicalProperties of Ethylene Vinyl Acetate
Jaehyun Kim,1 Janson Wang,2 Ho-Jong Kang,1 Frank Talke21 Department of Polymer Science and Engineering, Center for Advanced Functional Polymers,Dankook University, Seoul 140–714, Korea
2 Center for Magnetic Recording Research, University of California, San Diego, La Jolla, California 92093
Nano- and macro-mechanical properties of ethylenevinyl acetate (EVA) with various amounts of vinyl ace-tate (VA) have been investigated. Nano-mechanicalproperties (modulus and hardness) were obtainedusing nano-indentation measurements while macro-mechanical properties were determined using tensiletest measurements. A decrease in Young’s modulusand hardness was observed with increasing VA contentfor both nano- and macro-mechanical measurements.An increase in Young’s modulus and hardness wasobserved as a function of the draw ratio keeping theVA content constant. The difference between macro-and nano-mechanical properties as a function of VAcontent and draw ratio is explained in terms of crystal-linity and chain orientation. POLYM. ENG. SCI., 48:277–282, 2008. ª 2007 Society of Plastics Engineers
INTRODUCTION
Nano-mechanical properties of polymeric materials [1–
4] are becoming an importance design issue in the contact
performance of surfaces. In bio applications, nano-
mechanical properties affect tissue growth and cell prolif-
eration [5]. The performance of stents [6] in coronary
revascularization depends on the availability of suitable
drug carrier coating materials for consistent drug release
and structural stability. Soft elastomeric polymers [6] such
as poly (styrene-isobutylene-styrene) triblock copolymer
and ethylene vinyl acetate (EVA) copolymer are being
considered as drug carrier materials for drug eluting
stents. Thus, the understanding of the nano-mechanical
properties of these materials for nano and macro applica-
tions is of great importance.
The measurement of mechanical properties of polymers
on the nano scale has been limited in the past by the lack
of accurate instrumentation. Recently, atomic force mi-
croscopy (AFM) and nano-indentation analysis have
become available to study nano-mechanical properties of
thin films. Two experimental approaches related to nano-
indentation measurements [7–10] have been considered.
One is the ‘‘imaging method’’ where the characteristics of
plastically deformed residual indentations are analyzed.
The other approach is the ‘‘compliance method’’ where
the force-displacement characteristics are measured. For
polymeric materials, the compliance method is generally
used to avoid having to deal with relaxation effects. To
use the compliance method, one must consider the effect
of contact conditions including tip geometry, contact ge-
ometry, loading rate, and temperature.
In this investigation, we have studied the nano- and
macro-mechanical properties of various EVAs using
nano-indentation instrumentation and macro scale tensile
testing. In addition, the relationship between macroscopic
structure and nano-mechanical properties is investigated
as a function of the draw ratio used during the manufac-
turing process. The difference between macro and nano-
mechanical properties as a function of VA content and
draw ratio is explained in terms of cystallinity and chain
orientation.
EXPERIMENTAL
EVA with various amounts of vinyl acetate (VA) was
obtained (Hyundai Petroleum Chem. (SEETEC: EF221,
EF443, VS420) and Mitsui (EVAFLEX EVA150)). The
VA content was 3, 12, 21, and 33 wt%, respectively, and
the melt index was 0.6, 1.1, 2.0, and 30.0 g/min, respec-
tively. EVA films were prepared by compression molding.
Films with draw ratios between one and five were manu-
factured using a tensile tester (Lloyd LR-10K) equipped
with a 258C isothermal chamber. The Young’s modulus
and the hardness of the EVA films were determined on
the macro scale using a commercially available tensile
tester and a Shore hardness tester (TECLOCK: JIS K
6301 A). To obtain modulus and hardness on the nano
scale, a commercially available nano-indenter (Hysitron
Correspondence to: Ho-Jong Kang; e-mail: [email protected]
DOI 10.1002/pen.20883
Published online in Wiley InterScience (www.interscience.wiley.com).
VVC 2007 Society of Plastics Engineers
POLYMER ENGINEERING AND SCIENCE—-2008
Triboindenter, Fig. 1A) was used. In addition, the surface
roughness was determined using an atomic force micro-
scope (Digital Instruments NanoScope1 IIIa) while the
crystallinity was obtained with a differential scanning cal-
orimeter (METTLER TOLEDO DSC822e). A polarized
light microscope with tilting compensator (Leitz: POL-12)
was used to measure the birefringence and orientation of
the drawn films.
To determine the nano-mechanical properties of the
films, the compliance method was applied [7]. A Berko-
vich indenter was used to indent the EVA surface with
gradually increasing applied load (5 � 102 � 1 � 104 mN).To eliminate relaxation effects of viscoelastic materials,
the indenter was held at contact load for 5 s after reach-
ing Pmax. Thereafter, the indenter was gradually with-
drawn. The loading and unloading rate used in this study
was between 0.1 and 2.0 mN/s. After making indentation
on the surface of the sample, the scanning was performed
with same tip on a 5 � 5 mm2 surface of the indented
area with 256 � 256 data points about 1 min after inden-
tation. About 256 s was taken for the completion of each
AFM image.
The load and displacement curves (Fig. 1B) were
determined during the loading and unloading process. The
contact stiffness [7] was calculated from the slope of the
unloading curve (dP/dh) using
S���hmax
¼ dP
dh
� �hmax
¼ 2bffiffiffip
p Er
ffiffiffiffiffiffiffiffiffiAmax
p; (1)
where b is a constant related to the tip geometry, Amax is
the projected area at the maximum penetration, and Er is
the reduced modulus given by
1
Er
¼ ð1� v2ÞE
þ ð1� v2i ÞEi
: (2)
In Eq. 2, Ei and ni are the elastic modulus and Pois-
son’s ratio of the indenter, while E and n are the elastic
modulus and Poisson’s ratio of the film specimen, respec-
tively.
Nano-indentation testing was also used to determine
the nano-hardness of the EVA films. The nano-hardness is
given by
H ¼ Pmax
Amax
¼ Pmax
ch2c(3)
where Pmax is the peak load; Amax is the projected contact
area at maximum penetration hmax; c is 24.5 for a perfect
Berkovich indenter, and hc is the plastic displacement. To
make a comparison between nano-hardness and Shore
hardness, both hardness measurements were normalized
based on the hardness value of EVA with 3 wt% VA con-
tent, respectively.
RESULTS AND DISCUSSIONS
Figure 2 shows AFM images of impressions remaining
in EVA with 3, 12, and 21 wt% VA content, respectively,
after performing nano-indentations with a load of 500 mN.We observe that the indentations become broader and that
their depth increases as the VA content increases, i.e., the
response of the material becomes more elastic with
increasing VA content because the stiffness or elastic
modulus decreases with increasing VA content. In addi-
tion, we note that the footprint of the residual indents
decreases gradually with time. Clearly, the decrease of
the footprint is a function of the content of VA in EVA.
To minimize the effect of viscoelastic changes on the
evaluation of nano-mechanical properties, we have used
the load/unload curves for the determination of the me-
chanical properties. Figure 3 shows typical load/unload
curves for EVA as a function of the applied force. It is
apparent that the nano-indentation response of EVA is a
function of the VA content. EVA with a high amount of
VA has a lower Pmax and a greater displacement even
though the same load is applied. The deviation between
Pmax and the applied load was pronounced with increasing
FIG. 1. Schematic of nano-indentation test: (A) geometry of indentation by a tip (B) load and displacement
curve.
278 POLYMER ENGINEERING AND SCIENCE—-2008 DOI 10.1002/pen
VA content. Thus, increasing the VA content causes the
material to become ‘‘softer’’ and to exhibit more elastic
characteristics. This should be confirmed from the consid-
eration of the ratio of the elastic area to the area under
the loading curve in Fig. 3. Figure 3A and B illustrate the
effect of applied force on the load–displacement curve.
We observe that the penetration depth increases with an
increase in the applied force. Furthermore, the difference
between Pmax and the applied force decreases and the
slope of the unloading curve is independent of the applied
force. This allows the calculation of the nano-modulus as
shown in Fig. 4.
FIG. 2. AFM images of plastic impression remaining in EVA after nano-indentation with applied load of
500 mN: VA content is (A) 3 wt%, (B) 12 wt%, and (C) 21 wt%.
FIG. 3. Load-displacement curves for EVA with applied loads of (A)
500 mN and (B) 1000 mN.FIG. 4. Nano-mechanical properties of EVA as a function of indenta-
tion depth: (A) Young’s modulus (B) hardness.
DOI 10.1002/pen POLYMER ENGINEERING AND SCIENCE—-2008 279
The nano-modulus and the nano-hardness of EVA were
obtained from Eq. 1 using the slope of the unloading
curve. The results are plotted in Fig. 4. We observe that
EVA with 3 wt% VA has the highest value of Young’s
modulus and hardness. Increasing the amount of VA
results in a decrease of the Young’s modulus and hard-
ness. The VA content in EVA affects the viscoelastic
properties as well as the crystalline structure of EVA.
Increasing the VA content causes an increase in the elas-
tomeric behavior and retards the formation of crystalline
structures in EVA. Thus, EVA shows a ‘‘soft’’ character-
istic. This behavior is more pronounced in the modulus
results shown in Fig. 4A. A similar result was reported by
Briscoe et al. [4] for various polymeric materials. They
found that soft polymers do not show a dependence of
mechanical properties on penetration depth while rigid
polymers exhibit a high dependence of mechanical prop-
erties on penetration depth. They postulated that this
effect is due to a localized modification of the material
surface due to oxidation and aging in the atmosphere. In
our study, all EVA samples were prepared in the same
way and maintained at the same conditions. Therefore,
oxidation or aging in atmosphere does not seem to
explain our results. We conjecture that our results are
related to the modification of the morphology of the sur-
face layer during EVA film processing due to the pres-
ence of VA. Crystalline structures near a surface differ
from crystalline structures of the material in the center
because of cooling. Two common modifications of spher-
ulitic texture found in polymeric materials are the so-
called row structure and the transcrystalline layer. For
crystalline EVA, and EVA with low amount of VA, trans-
verse growth equivalent to that along the radius of a
spherulite is likely to occur. This deformed crystalline
structure is likely to affect the nano-mechanical properties
near the surface region. Our results indicate that the influ-
ence of deformed crystalline structures is more dominant
in hardness measurements than modulus measurements.
In Fig. 5A, the nano-modulus for a 3000 nm indenta-
tion depth is plotted together with the macro modulus
measured in a tensile test at a strain rate of 100 mm/min.
Both the nano and the macro modulus values are seen to
decrease with increasing VA content. It is well known
that mechanical properties of polymers depend strongly
on the macro structures developed during polymer pro-
cessing. The surface roughness, crystalline structure, and
chain orientation are typical macroscopic properties devel-
oped during polymer processing. Since the films were pre-
pared by compression molding using the same mold, the
surface roughness of the films is very similar (Fig. 6)
even though their crystallinity levels vary. Birefringence
measurements show that chain orientation is not devel-
oped in unoriented EVA films by compression molding.
This result was also confirmed in our differential scanning
calorimetry (DSC) data shown in Fig. 6. EVA with high
VA content has lower relative crystallinity (35 J/g) than
EVA with low VA content (120 J/g). The decrease of the
Young’s modulus as a function of the VA content seems
to be related to the increasing difficulties of forming
FIG. 5. Macro- and nano-mechanical properties of EVA as a function
of VA content: (A) modulus (B) normalized hardness.
FIG. 6. Surface roughness and crystallinity of EVA as a function of
VA content.
280 POLYMER ENGINEERING AND SCIENCE—-2008 DOI 10.1002/pen
macroscopic crystalline structures with increasing VA
content. On the other hand, the effect of macroscopic struc-
ture on the nano-modulus is reduced because the intentions
are performed at a scale much smaller than the dimensions
of macroscopic structures. In other words, the deformation
due to a nano-indentation does not affect the deformation
of the macroscopic structure. Figure 5B shows the effect
of VA content on hardness. Both the macro- and the
nano-hardness decrease with increasing VA content. It is
not possible to compare the hardness values directly,
since the measurement of macro-hardness is a function of
the method used for the hardness measurement. In a
Shore hardness measurement a value of 0 is obtained if
the indenter penetrates the EVA completely, and a read-
ing of 100 is obtained if no penetration occurs. On the
other hand, nano-hardness values are related to the meas-
ured pressure of a nano tip indenter (Eq. 3). To compare
the measurements, we made a relative comparison of
hardness values by normalizing the hardness data based
on the hardness value of EVA with 3 wt% VA content.
Our results show that the decrease in macro-hardness due
to increasing VA content is much less than that observed
for the nano-hardness measurements. Again, we believe
that this result is due to the fact that the nano-indentation
is performed at a nano scale while the macro-hardness
measurement is an average of the hardness on the macro
scale.
In addition to the effect of the crystalline structure on
mechanical properties, the chain orientation is another
important parameter in polymer processing that governs
mechanical properties. EVA films with 3 wt% VA were
stretched at different draw ratios. In Fig. 7, the macro-
and nano-moduli for the oriented EVA films are shown.
We observe that the macro-modulus increases strongly
with increasing draw ratio but that the nano-modulus is
almost independent of the draw ratio. This behavior can
be explained in terms of the structural changes occurring
during the draw process. In Fig. 8, relative crystallinity
and birefringence data are shown as a function of the
draw ratio. Although EVA is a rubber like material,
stress induced crystallization does not take place during
the stretching. However, molecular orientation was intro-
duced by uniaxial drawing for draw ratios between two
and three. Since EVA films are uniaxially stretched, both
the crystalline and amorphous chains gradually align
their axis along the machine direction. The birefringence
results are macroscopic scale definitions based on the av-
erage mean value of crystalline and amorphous orienta-
tion. Therefore, the birefringence values are correlated
with the macro-modulus. However, in nano-indentation
measurements, the orientation in a local area where the
nano-indentation test is conducted may differ signifi-
cantly from the average orientation indicated by the bire-
fringence. Thus, there is a high likelihood that the orien-
tation varies a depending on whether the indenter is
located in a local crystalline phase or local amorphous
phase.
FIG. 7. Macro- and nano-modulus of oriented EVA films as a function
of draw ratio.
FIG. 8. Macroscopic structures of oriented EVA films.
DOI 10.1002/pen POLYMER ENGINEERING AND SCIENCE—-2008 281
CONCLUSION
Nano-indentation measurements were carried out to
determine nano-mechanical properties of EVA with vari-
ous amounts of VA. Nano-hardness and nano Young’s
modulus data were obtained using the compliance
method. The results show that nano-mechanical properties
in a localized area are controlled by the intrinsic visco-
elastic properties and the macroscopic structures devel-
oped during polymer processing. The effect of the macro-
scopic structures on the nano-modulus is much less than
its effect on the macro-modulus. The reason for the
observed behavior is related to the fact that nano-
indentation measurements are performed in a localized
area of the macroscopic structure. Macro-hardness mea-
surements and macro Young’s modulus measurements are
averages on a scale much larger than the scale of nano-
hardness and nano Young’s modulus values, i.e., macro-
scopic measurements are averages over a large area and
do not show strong local effects.
REFERENCES
1. B.D. Beake and G.J. Leggett, Polymer, 43, 319 (2002).
2. J.F. Graham, M. Kovar, P.R. Norton, P. Pappalardo, J. Van
Loon, and O.L. Warren, J. Mater. Res., 13, 3565 (1988).
3. D. Drechsler, A. Karbach, and H. Fuchs, Appl. Phys. A, 66,S825 (1998).
4. B.J. Briscoe, L. Fiori, and E. Pelillo, J. Phys. D: Appl.Phys., 31, 2395 (1998).
5. C.S. Chen, M. Mrksich, S. Huang, G.M. Whitesides, and
D.E. Ingber, Science, 276, 1425 (1997).
6. S.V. Ranade, K.M. Miller, R.E. Richard, A.K. Chan, M.J.
Allen, and M.N. Helmus, J. Bio. Mater. Res. Part A, 71,625 (2004).
7. G.M. Pharr, W.C. Oliver, and F.R. Brotzen, J. Mater. Res.,7, 613 (1992).
8. M.F. Doerner and W.D. Nix, J. Mater. Res., 1, 601 (1986).
9. I. Sneddon, Int. J. Eng. Sci., 3, 47 (1965).
10. B.J. Briscoe, K.S. Sebastain, and M.J. Adams, J. Phys. D:Appl. Phys., 27, 1156 (1994).
282 POLYMER ENGINEERING AND SCIENCE—-2008 DOI 10.1002/pen
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