A study of nano-mechanical and macro-mechanical properties of ethylene vinyl acetate

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<ul><li><p>A Study of Nano-Mechanical and Macro-MechanicalProperties of Ethylene Vinyl Acetate</p><p>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 140714, Korea</p><p>2 Center for Magnetic Recording Research, University of California, San Diego, La Jolla, California 92093</p><p>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 Youngs modulusand hardness was observed with increasing VA contentfor both nano- and macro-mechanical measurements.An increase in Youngs 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., 48:277282, 2008. 2007 Society of Plastics Engineers</p><p>INTRODUCTION</p><p>Nano-mechanical properties of polymeric materials [1</p><p>4] are becoming an importance design issue in the contact</p><p>performance of surfaces. In bio applications, nano-</p><p>mechanical properties affect tissue growth and cell prolif-</p><p>eration [5]. The performance of stents [6] in coronary</p><p>revascularization depends on the availability of suitable</p><p>drug carrier coating materials for consistent drug release</p><p>and structural stability. Soft elastomeric polymers [6] such</p><p>as poly (styrene-isobutylene-styrene) triblock copolymer</p><p>and ethylene vinyl acetate (EVA) copolymer are being</p><p>considered as drug carrier materials for drug eluting</p><p>stents. Thus, the understanding of the nano-mechanical</p><p>properties of these materials for nano and macro applica-</p><p>tions is of great importance.</p><p>The measurement of mechanical properties of polymers</p><p>on the nano scale has been limited in the past by the lack</p><p>of accurate instrumentation. Recently, atomic force mi-</p><p>croscopy (AFM) and nano-indentation analysis have</p><p>become available to study nano-mechanical properties of</p><p>thin lms. Two experimental approaches related to nano-</p><p>indentation measurements [710] have been considered.</p><p>One is the imaging method where the characteristics of</p><p>plastically deformed residual indentations are analyzed.</p><p>The other approach is the compliance method where</p><p>the force-displacement characteristics are measured. For</p><p>polymeric materials, the compliance method is generally</p><p>used to avoid having to deal with relaxation effects. To</p><p>use the compliance method, one must consider the effect</p><p>of contact conditions including tip geometry, contact ge-</p><p>ometry, loading rate, and temperature.</p><p>In this investigation, we have studied the nano- and</p><p>macro-mechanical properties of various EVAs using</p><p>nano-indentation instrumentation and macro scale tensile</p><p>testing. In addition, the relationship between macroscopic</p><p>structure and nano-mechanical properties is investigated</p><p>as a function of the draw ratio used during the manufac-</p><p>turing process. The difference between macro and nano-</p><p>mechanical properties as a function of VA content and</p><p>draw ratio is explained in terms of cystallinity and chain</p><p>orientation.</p><p>EXPERIMENTAL</p><p>EVA with various amounts of vinyl acetate (VA) was</p><p>obtained (Hyundai Petroleum Chem. (SEETEC: EF221,</p><p>EF443, VS420) and Mitsui (EVAFLEX EVA150)). The</p><p>VA content was 3, 12, 21, and 33 wt%, respectively, and</p><p>the melt index was 0.6, 1.1, 2.0, and 30.0 g/min, respec-</p><p>tively. EVA lms were prepared by compression molding.</p><p>Films with draw ratios between one and ve were manu-</p><p>factured using a tensile tester (Lloyd LR-10K) equipped</p><p>with a 258C isothermal chamber. The Youngs modulusand the hardness of the EVA lms were determined on</p><p>the macro scale using a commercially available tensile</p><p>tester and a Shore hardness tester (TECLOCK: JIS K</p><p>6301 A). To obtain modulus and hardness on the nano</p><p>scale, a commercially available nano-indenter (Hysitron</p><p>Correspondence to: Ho-Jong Kang; e-mail: hjkang@dku.edu</p><p>DOI 10.1002/pen.20883</p><p>Published online in Wiley InterScience (www.interscience.wiley.com).</p><p>VVC 2007 Society of Plastics Engineers</p><p>POLYMER ENGINEERING AND SCIENCE-2008</p></li><li><p>Triboindenter, Fig. 1A) was used. In addition, the surface</p><p>roughness was determined using an atomic force micro-</p><p>scope (Digital Instruments NanoScope1 IIIa) while the</p><p>crystallinity was obtained with a differential scanning cal-</p><p>orimeter (METTLER TOLEDO DSC822e). A polarized</p><p>light microscope with tilting compensator (Leitz: POL-12)</p><p>was used to measure the birefringence and orientation of</p><p>the drawn lms.</p><p>To determine the nano-mechanical properties of the</p><p>lms, the compliance method was applied [7]. A Berko-</p><p>vich indenter was used to indent the EVA surface with</p><p>gradually increasing applied load (5 102 1 104 mN).To eliminate relaxation effects of viscoelastic materials,</p><p>the indenter was held at contact load for 5 s after reach-</p><p>ing Pmax. Thereafter, the indenter was gradually with-drawn. The loading and unloading rate used in this study</p><p>was between 0.1 and 2.0 mN/s. After making indentationon the surface of the sample, the scanning was performed</p><p>with same tip on a 5 5 mm2 surface of the indentedarea with 256 256 data points about 1 min after inden-tation. About 256 s was taken for the completion of each</p><p>AFM image.</p><p>The load and displacement curves (Fig. 1B) were</p><p>determined during the loading and unloading process. The</p><p>contact stiffness [7] was calculated from the slope of the</p><p>unloading curve (dP/dh) using</p><p>Shmax</p><p> dPdh</p><p> hmax</p><p> 2bp</p><p>p ErAmax</p><p>p; (1)</p><p>where b is a constant related to the tip geometry, Amax isthe projected area at the maximum penetration, and Er isthe reduced modulus given by</p><p>1</p><p>Er 1 v</p><p>2E</p><p> 1 v2i </p><p>Ei: (2)</p><p>In Eq. 2, Ei and ni are the elastic modulus and Pois-sons ratio of the indenter, while E and n are the elastic</p><p>modulus and Poissons ratio of the lm specimen, respec-</p><p>tively.</p><p>Nano-indentation testing was also used to determine</p><p>the nano-hardness of the EVA lms. The nano-hardness is</p><p>given by</p><p>H PmaxAmax</p><p> Pmaxch2c</p><p>(3)</p><p>where Pmax is the peak load; Amax is the projected contactarea at maximum penetration hmax; c is 24.5 for a perfectBerkovich indenter, and hc is the plastic displacement. Tomake a comparison between nano-hardness and Shore</p><p>hardness, both hardness measurements were normalized</p><p>based on the hardness value of EVA with 3 wt% VA con-</p><p>tent, respectively.</p><p>RESULTS AND DISCUSSIONS</p><p>Figure 2 shows AFM images of impressions remaining</p><p>in EVA with 3, 12, and 21 wt% VA content, respectively,</p><p>after performing nano-indentations with a load of 500 mN.We observe that the indentations become broader and that</p><p>their depth increases as the VA content increases, i.e., the</p><p>response of the material becomes more elastic with</p><p>increasing VA content because the stiffness or elastic</p><p>modulus decreases with increasing VA content. In addi-</p><p>tion, we note that the footprint of the residual indents</p><p>decreases gradually with time. Clearly, the decrease of</p><p>the footprint is a function of the content of VA in EVA.</p><p>To minimize the effect of viscoelastic changes on the</p><p>evaluation of nano-mechanical properties, we have used</p><p>the load/unload curves for the determination of the me-</p><p>chanical properties. Figure 3 shows typical load/unload</p><p>curves for EVA as a function of the applied force. It is</p><p>apparent that the nano-indentation response of EVA is a</p><p>function of the VA content. EVA with a high amount of</p><p>VA has a lower Pmax and a greater displacement eventhough the same load is applied. The deviation between</p><p>Pmax and the applied load was pronounced with increasing</p><p>FIG. 1. Schematic of nano-indentation test: (A) geometry of indentation by a tip (B) load and displacement</p><p>curve.</p><p>278 POLYMER ENGINEERING AND SCIENCE-2008 DOI 10.1002/pen</p></li><li><p>VA content. Thus, increasing the VA content causes the</p><p>material to become softer and to exhibit more elastic</p><p>characteristics. This should be conrmed from the consid-</p><p>eration of the ratio of the elastic area to the area under</p><p>the loading curve in Fig. 3. Figure 3A and B illustrate the</p><p>effect of applied force on the loaddisplacement curve.</p><p>We observe that the penetration depth increases with an</p><p>increase in the applied force. Furthermore, the difference</p><p>between Pmax and the applied force decreases and theslope of the unloading curve is independent of the applied</p><p>force. This allows the calculation of the nano-modulus as</p><p>shown in Fig. 4.</p><p>FIG. 2. AFM images of plastic impression remaining in EVA after nano-indentation with applied load of</p><p>500 mN: VA content is (A) 3 wt%, (B) 12 wt%, and (C) 21 wt%.</p><p>FIG. 3. Load-displacement curves for EVA with applied loads of (A)</p><p>500 mN and (B) 1000 mN.FIG. 4. Nano-mechanical properties of EVA as a function of indenta-</p><p>tion depth: (A) Youngs modulus (B) hardness.</p><p>DOI 10.1002/pen POLYMER ENGINEERING AND SCIENCE-2008 279</p></li><li><p>The nano-modulus and the nano-hardness of EVA were</p><p>obtained from Eq. 1 using the slope of the unloadingcurve. The results are plotted in Fig. 4. We observe that</p><p>EVA with 3 wt% VA has the highest value of Youngs</p><p>modulus and hardness. Increasing the amount of VA</p><p>results in a decrease of the Youngs modulus and hard-</p><p>ness. The VA content in EVA affects the viscoelastic</p><p>properties as well as the crystalline structure of EVA.</p><p>Increasing the VA content causes an increase in the elas-</p><p>tomeric behavior and retards the formation of crystalline</p><p>structures in EVA. Thus, EVA shows a soft character-</p><p>istic. This behavior is more pronounced in the modulus</p><p>results shown in Fig. 4A. A similar result was reported by</p><p>Briscoe et al. [4] for various polymeric materials. They</p><p>found that soft polymers do not show a dependence of</p><p>mechanical properties on penetration depth while rigid</p><p>polymers exhibit a high dependence of mechanical prop-</p><p>erties on penetration depth. They postulated that this</p><p>effect is due to a localized modication of the material</p><p>surface due to oxidation and aging in the atmosphere. In</p><p>our study, all EVA samples were prepared in the same</p><p>way and maintained at the same conditions. Therefore,</p><p>oxidation or aging in atmosphere does not seem to</p><p>explain our results. We conjecture that our results are</p><p>related to the modication of the morphology of the sur-</p><p>face layer during EVA lm processing due to the pres-</p><p>ence of VA. Crystalline structures near a surface differ</p><p>from crystalline structures of the material in the center</p><p>because of cooling. Two common modications of spher-</p><p>ulitic texture found in polymeric materials are the so-</p><p>called row structure and the transcrystalline layer. For</p><p>crystalline EVA, and EVA with low amount of VA, trans-</p><p>verse growth equivalent to that along the radius of a</p><p>spherulite is likely to occur. This deformed crystalline</p><p>structure is likely to affect the nano-mechanical properties</p><p>near the surface region. Our results indicate that the inu-</p><p>ence of deformed crystalline structures is more dominant</p><p>in hardness measurements than modulus measurements.</p><p>In Fig. 5A, the nano-modulus for a 3000 nm indenta-</p><p>tion depth is plotted together with the macro modulus</p><p>measured in a tensile test at a strain rate of 100 mm/min.</p><p>Both the nano and the macro modulus values are seen to</p><p>decrease with increasing VA content. It is well known</p><p>that mechanical properties of polymers depend strongly</p><p>on the macro structures developed during polymer pro-</p><p>cessing. The surface roughness, crystalline structure, and</p><p>chain orientation are typical macroscopic properties devel-</p><p>oped during polymer processing. Since the lms were pre-</p><p>pared by compression molding using the same mold, the</p><p>surface roughness of the lms is very similar (Fig. 6)</p><p>even though their crystallinity levels vary. Birefringence</p><p>measurements show that chain orientation is not devel-</p><p>oped in unoriented EVA lms by compression molding.</p><p>This result was also conrmed in our differential scanning</p><p>calorimetry (DSC) data shown in Fig. 6. EVA with high</p><p>VA content has lower relative crystallinity (35 J/g) than</p><p>EVA with low VA content (120 J/g). The decrease of the</p><p>Youngs modulus as a function of the VA content seems</p><p>to be related to the increasing difculties of forming</p><p>FIG. 5. Macro- and nano-mechanical properties of EVA as a function</p><p>of VA content: (A) modulus (B) normalized hardness.</p><p>FIG. 6. Surface roughness and crystallinity of EVA as a function of</p><p>VA content.</p><p>280 POLYMER ENGINEERING AND SCIENCE-2008 DOI 10.1002/pen</p></li><li><p>macroscopic crystalline structures with increasing VA</p><p>content. On the other hand, the effect of macroscopic struc-</p><p>ture on the nano-modulus is reduced because the intentions</p><p>are performed at a scale much smaller than the dimensions</p><p>of macroscopic structures. In other words, the deformation</p><p>due to a nano-indentation does not affect the deformation</p><p>of the macroscopic structure. Figure 5B shows the effect</p><p>of VA content on hardness. Both the macro- and the</p><p>nano-hardness decrease with increasing VA content. It is</p><p>not possible to compare the hardness values directly,</p><p>since the measurement of macro-hardness is a function of</p><p>the method used for the hardness measurement. In a</p><p>Shore hardness measurement a value of 0 is obtained if</p><p>the indenter penetrates the EVA completely, and a read-</p><p>ing of 100 is obtained if no penetration occurs. On the</p><p>other hand, nano-hardness values are related to the meas-</p><p>ured pressure of a nano tip indenter (Eq. 3). To comparethe measurements, we made a relative comparison of</p><p>hardness values by normalizing the hardness data based</p><p>on the hardness value of EVA with 3 wt% VA content.</p><p>Our results show that the decrease in macro-hardness due</p><p>to increasing VA content is much less than that observed</p><p>for the nano-hardness measurements. Again, we believe</p><p>that this result is due to the fact that the nano-indentation</p><p>is performed at a nano scale while the macro-hardness</p><p>measurement is an average of the hardness on the macro</p><p>scale.</p><p>In addition to the effect of the crystalline structure on</p><p>mechanical properties, the chain orientation is another</p><p>important parameter in polymer processing that governs</p><p>mechanical properties. EVA lms with 3 wt% VA were</p><p>stretched at different draw ratios. In Fig. 7, the macro-</p><p>and nano-moduli for the oriented EVA lms are shown.</p><p>We observe that the macro-modulus increases strongly</p><p>with increasing draw ratio but that the nano-modulus is</p><p>almost independent of the draw ratio. This behavior can</p><p>be explained in terms of the structural changes occurring</p><p>during the draw process. In Fig. 8, relative crystallinity</p><p>and birefringence data are shown as a function of the</p><p>draw ratio. Although EVA is a rubber like material,</p><p>stress induced crystallization does not take place during</p><p>the stretching. However, molecular orientation was intro-</p><p>duced by uniaxial drawing for draw ratios between two</p><p>and three. Since EVA lms are uniaxially stretched, both</p><p>the crystalline and amorphous chains gradually align</p><p>their axis along the machine direction. The birefringence</p><p>results are macroscopic scale denitions...</p></li></ul>

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