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Title: Preparation, morphology, mechanical properties and fracture resistance of nanocomposites comprising montmorillonite and polypropylene Author(s): Keener BD, Hudson SD, ~~ Sirivat A
1-LY,
Moran IW, Phongphour
Output This Record Bibliographic Fields
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Source: MATERIALS RESEARCH INNOVATIONS 9 (4): 96-98 DEC 2005 Document Type: Article
i:/
Language: English Cited References: 45
Times Cited: 0
Author Keywords: polypropylene; maleic-anhydride-modified polypropylene; sodium montmorillonite, silane coupling agent; nanocomposite filler, mechanical properties; environmental crack resistance; crack propagation; melt-state compounding; nanocomposite KeyWords Plus: ENVIRONMENTAL-STRESS CRACKING; POLYMER MELT INTERCALATION; MODIFIED LAYERED SILICATES; HIGHDENSITY POLYETHYLENE; NYLOI\J 6-CLAY HYBRID; CLAY HYBRIDS; COMPOSITES; BEHAVIOR; MODEL Addresses: Keener BD (reprint author), Chulalongkorn Univ, Petr & Petrochem Coli, Conduct & Electroact Res Unit, Bangkok, Thailand Chulalongkorn Univ, Petr & Petrochem Coli, Conduct & Electroact Res Unit, Bangkok, Thailand Case Western Reserve Univ, Dept Macromol Sci, Cleveland, OH 44106 USA E-mail Addresses:
[email protected] Publisher: MATRICE TECHNOLOGY LIMITED, PO BOX 41 11 POULTON FY6 8GD, LACANSHIRE, ENGLAND Subject Category: MATERIALS SCIENCE, MULTIDISCIPLINARY IDS Number: 997HT ISSN: 1432-8917
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48
Preparation, morphology, mechanical properties and fracture resistance of nanocomposites comprising montmorillonite and polypropylene
B.D. Keener, S.D. Hudson, Y. Li, I. Moran
Department of Macromolecular Science, Case Western Reserve University, Cleveland, OH 44106,· USA C. Phoi:tgphour, A. Sirivat*
Conductive and ElectroacNve Research Unit Petrd(eum and Petrochemical College, Chulalongkorn University, Bangkok, Thailand
ABSTij.ACT Nanocomposites comprising montmorillonite and polypropylene have been developed. First, montmorillonite silicate nano-particles were amine-functionalized with a silane coupling agent. Polypropylene was then grafted onto the inorganic surface through reaction of maleic anhydride-modified polypropylene (MAPP) with the amine, as characterized by DRIFT· spectroscopy. X-ray diffraction analysis of treated and untreated montmorillonite indicated that the silane coupling agent and polymer expanded the montmorillonite layer spacing. The degree of intercalation and the ability of the grafted chains to crystallize increased with increasing amount of grafted polymer. These grafted-chain nanopmiicles were melt-blended \Vith unmodified polymer (iPP). TEM analysis of microtomed nanocompositc sections indicated many individually dispersed silicate layers. but most of the layers \\ere found in two- to six-layer stacks. Composites based on the nanofillcr sho\\cd modest increases in tensile modulus and yield stress at 5 wt. cYo clay \Vhile retaining some ductility.
TEM analysis of elongated nanocomposite films demonstrated the ability of the filler particles to hinder the propagation of deformation zones. Environmental stress cracking resistance was also improved, as time-to-failure increased three-fold with addition of 10 wt.% filler.
* Corresponding author:
[email protected]
INTRODUCTION Recent developments of nanocomposites comprising polymer and a layered silicate 4
have displayed exceptional barrier, 1,2 fire retardant, 3 and mechanical properties. At the heart of their success are the nanoscale dispersion and high particle aspect ratio made possible by various synthetic preparations. 5-7 Models and experimen~s suggest that dispersion at this level is governed by thermodynamic mixing s-Jo of the particle· surface treatment with the polymer matrix. In some instances,
trea~ed
particles have been· dispersed more efficiently by first
nuxmg with monomer. with \Vhich the entropy of m1xmg 1s greater, and then polymerizing. 7· 11 Most polymer/silicate nanocomposites involve polar polymers that have a favorable interaction with the clay surface: Polyolefin nanocomposites have been more difficult to achieve, because of their nonpolar nature. Recent strategies produced polypropylene (PP) intercalation and beneficial properties. The treatment of Kurokawa et al. 12 •13 involves several steps: cation exchange with hexadecyl ammonium, intercalation of an amphiphilic monomer, polymerization of that monomer, and finally melt-mixing with a polar functional (maleated) PP. The approach of Kawasumi et al. 14 · 15 is simpler: cation exchange with stearyl ammonium followed by meltmixing with a maleated PP. This strategy, however, uses a large amount of surface treatment (the amount of stearyl amine is half that of the clay) and of functionalized polymer.
Intercalated polyolefin nanocomposites have also been prepared usmg unfunctionalized polymer. 16 In this report, we describe a treatment for covalent attachment of PP to silicate clay layers. 17 The treatment consists of two steps. First, clay is treated with a few percent of an aminosilane coupling agent. The amine-functionalized particles are then reacted with maleated PP through amidization to produce grafted chains. The nanofiller can then be dispersed in unmodified PP by conventional melt mixing procedures. The composites exhibit improved mechanical properties and environmental stress cracking resistance. Environmental stress cracking (ESC) . is an important phenomenon for structural plastics (including 19
polyethylene (PE) and PP) in aggressive-environment applications. 18 • These include gaslines, - electrical insulation, an_d waste piping. ESC is characterized by the diffusion and absorption of an external agent into the local damage zones of semicrystalline materials, where the agent then plasticizes the disentanglement of the intercrystalline tie molecules. 20 ·21 In addition to the physical tests, related nanocomposite morphologies (thin film deformation analyses and ESC surface textures) are also investigated using electron microscopic methods.
EXPERIMENTAL
Materials Nanocomposites were prepared from sodium montmorillonite silicate clay, ammosilane coupling agent, maleic anhydride modified isotactic polypropylene (MAPP), and unmodified isotactic polypropylene (iPP). The amount of each component was designated by three weight fractions: x, of silane in the silane-treated clay; y, of treated clay in the chaintethered nanofiller; and z, of nanofiller in the nanocomposite. Thus, the weight fractions of each component in the final nanocomposite are: z. and
Wpp =
1-z.
\VcJa,. =
(1-x) v- z.
Wsilanc =
x y z.
Wr--JAPP =
(1-y)
Na-montmorillonite was obtained from two sources (Kunipia-F, Kunimine Ind.; and Cloisite-Na, Southern Clay Products) and used as received. For comparison purposes, talcfilled composites were also investigated. Monofunctional alkoxy silanes were obtained from United Chemical Technologies. Two amino silanes were employed: 3-aminopropyl diisopropylethoxysilane (217 g/mol.) and 3-aminopropyl
dimethylethoxysilane
Silane
(161 g/mol.).
agents,
dimethyl-octadecyl
methoxysilane (343 g/mol.) and diphenyl-methyl ethoxysilane (242 g/mol.), without amine were used as controls. I ,3-bis(3-aminopropyl)-I, I ,3 ,3-tetramethyldisiloxane (siloxane dimer) was obtained for characterization purposes. Maleic anhydride (MA)-modified polypropylene, PolyBondTM 3002 (maleic content fMA = 0.2%wt., MFI = 7, ASTM D1238-57T), was supplied by Uniroyal Chemical. The weight average molecular weight can be estimated
(~
180,000)··from the melt flow index
according to the relationship Mw = 290,000 MFr0235 , which was derived from data published in the Polymer Handbook. 22 Assuming that maleation proceeds by a chain scission process,~ 3 the number average molecular weight may be roughly estimated from the degree ofmaleation: Mn ~ (nMA) (MMA)/(fMA), where nMA, MMA, and fMArepresent the number of maleic anhydride units per chain, and the molar mass and weight fraction of MA, respectively. According to this assumption, the polydispersity is approximately 4. Matrix isotactic polypropylene, Montell ProfaxTM 6823, (MFI = O.Sg/1 0 min, 230 °C, 2.16 kg, Mw
~
340,000) was supplied
by General Polymers/Ashland Chemical.
Functiona/ization with Silane Coupling Agent Montmorillonite (clay) was dispersed at typically 1 wi.% in either water or climethylacetamide (DMAc) at 85°C. In either liquid, the clay dispersions \vere cloudy and yellowish in appearance. Based on visual observation of particle settling, \Vater dispersed the
clay better than did DMAc. The coupling agent was dissolved at 2% by volume in either a 95/5 percent mixture of ethanol/water (for reaction with clay dispersed in water) or in DMAc (for reaction with clay dispersed in DMAc) and allowed to mix for half an hour at room temperature. The coupling agent solution was then added to the inorganic dispersion in the appropriate amount and the mixture was allowed to react at 90-1 00
oc for half an hour. The
product was then filtered and dried under vacuum for 24 hours, and then ground into a fine powder with mortar and pestle or a ball-mill.
Amidization (chain grafting) The amine-functional clay powder
w~s
redispersed into DMAc at 1-10 o/owt. and 90
oc. The maleic anhydride-modified polyolef!n was dissolved in hot xylenes (130 °C) at
I- 3
w1.%. Under these conditions, the hydrolyzed-diacid form of the anhydride is not expected. The cloudy DMAc solution was added dropwise to the clear colorless polymer solution. As the amount ofDMAc solution increased, the intensity ofthe cloudiness naturally increased. If the amount of DMA_c relative to xylene \vas sufficiently large, a small further increase in the fraction of DMAc caused the s0lutes (polymer and treated clay) to clump together, leaving the solvent mixture clear and colorless. When additional xylene was added, the solute dispersed again, forming a cloudy yellow dispersion. This behavior indicated that the polymer and clay act as a single solute. (In contrast, when the MAPP was replaced by unmodified polypropylene, the polymer and clay remained separate: when the unmodified polypropylene precipitated, the treated clay remained dispersed in the solvent mixture. The polypropylene precipitate was white.) After all of the DMAc solution was added, the mixture was reacted for 15 minutes at 125
oc. The 'nanofiller solution·
precipitate was filtered, washed in methanoL vacuumdried at 40
was allmved to cool, and the
oc for 24 hours. and ground
into a fine powder with a mortar and pestle. For example, when y = 50%, the resulting
powder was similar, but lighter, in appearance to the original clay. The yield of this reaction was greater than 90%.
Characterization ofparticle treatment X-ray diffraction (XRD) was used to examine the morphology following each step of nanocomposite preparation. Diffractograms, obtained with a Phillips X-ray diffractometer, using monochromated Cu K< radiation, were analyzed using Bragg's Law and the Scherrer equation. Diffuse reflectance (DRIFT) infrared spectra were acquired using a BioRad FTIR at a resolution of 2 cm-1 with 16 coadded scans. These spectra provided direct analysis of the reaction between surface amine and polymer anhydride groups.
Nanocomposite compounding and characterization Composites were melt-mixed using one of four different processing methods. For example, using a twin triangle-screw compounder (Brabender), iPP pellets were first melted at 180 °C, to which the nanofiller powder was added slowly by sprinkling. The material was then sheared at 180-190
oc and 140 rpm for up to
15 minutes. Using a single screw extmder,
drymixed iPP pellets and nanofiller powder were added simultaneously to the extmder hopper, then mixed at 200
oc
oc and 50 rpm. The material residence time in the extruder at 200
was under 3 minutes. Using a table-top twin screw extruder, supplied by DACA
Instruments, powder and pellets were mixed at 200
oc for 3 minutes. Using a two roll mill,
having the front and back cylinders heated to approximately 180 and 200 °C, respectively, polymer pellets were first melted into a rolling sheet_ to which the nanofiller pO\vder \vas added. Mixing was facilitated by folding the melt stream together with the use of a spatula: this process was sustained for five minutes.
Filler particle size and dispersion level of the composites were analyzed visually and by XRD (as described above) and optical and electron microscopy. Phase contrast optical microscopy was done on 9 cx:m-thick samples cut from a molded sheet with a glass knife rotary microtome. The sections were embedded in a non-swelling index matching fluid. The morphology of a twin-screw compounded nanocomposite was examined at higher resolution with transmission electron microscopy (TEM), using a JEOL 1OOCX, operated at 1OOk V. Nanocomposite samples were cut using a diamond knife ultramicrotome at room temperature, and placed on supporting copper grids for TEM analysis.
Mechanical tests All test samples were compression molded into sheets by pressing them between two Mylar-coated ferrotype plates spaced by 1.5 mm. The. molder
plat~s
were heated to 220
oc
then forced together at 500 psi to melt the composite, after which the pressure was released for 10 seconds. Compression was then resumed at 3000 psi for 90 seconds. released for 15 seconds, and resumed again at 3000 psi for 60 seconds. This was followed by rapid water and air cooling of the molder plates under compression. Mechanical properties of nanocomposites with varymg x/y/z formulations were quantified according to ASTM 693D. Dog-bone specimens were cut out of 1.5 mm thick compression-molded films using an ASTM M-III sized cutter and a press. Elongation was done at room temperature and a constant strain rate of 1.0/min (1 00% per minute) up to failure.
ESC tests Rectangular (20
111111
x 75 mm x 1.5 mm) ESC test samples were milled from
compression molded sheets. Precise 3mm cuts were made by pushing a fresh razor blade into
both. sample sides. To ensure identical cut lengths for multiple specimens, the razor was pushed simultaneously into four test samples, which were secured side-to-side in a clamp. The pieces were then fixed in the creep-test clips and the test arms were placed in a surfactant solution (Igepal C0-630® (GAF), 20% in distilled water). The sample arm was then connected to enough weight to apply 10 MPa of nomimal stress to the sample, and the hour timer was started. ESC tests were done on pure i-PP, talc-filled i-PP and a 2/50/10 nanocomposite. Pictures of the test specimen were taken in intervals of 20 hours using a close-up camera. Lab View software image analysis provided the measurement of the crack opening per test time. ESC fracture surfaces were examined for qualitative failure analysis using a JEOL 840A Scanning Electron Microscope.
RESULTS AND DISCUSSION Nanofiller preparation and characterization The treatment of montmorillonite by the silane coupling agent was investigated by XRD (Figure I). Montmorillonite exhibits two types of peaks: 000 and hkO, representing correlations between and within silicate layers, respectively. The 000 peaks are generally slightly aperiodic, because the spacing between silicate layers is somewhat irregular. The absence of hkO reflections and the unique shape of the hkO reflections (steep on the low angle side and gradually decaying on the high angle side) indicate that the individual 2: I silicate layers are shifted randomly relative to each other. 24 12.3
A is the characteristic interlayer 001
spacing of the clay as received (Figure 1, a).
Fully dried clay (in vacuo at 250°C) exhibits a spacing of 9.8 A. indicating that water is bound to the silicate surface in the as received clay. After silane treatment the 001 spacing increases, depending on the amount of silane applied (Figure I, b and c). Application of
4.0wt.% silane causes an increase of only 1.1
A (from
12.3 to 13.4
A,
a gallery thickness of
3.6 A). Further increase in layer spacing is produced at higher silane concentrations. These changes in d-spacing indicate that the silane coupling agent intercalates into the galleries of the silicate, in addition to likely reaction at the edges of the silicate. The narrow gallery thickness at low concentrations suggests that the silane adopts an inclined conformation. At higher concentrations, the silane extends away from the face of the silicate. Under aqueous acid conditions, the amino group on the silane may also ionize and displace a sodium ion from the interlayer gallery.2s We sought to prevent this reaction by employing neutral conditions with minimal water.26 To exclude the possibility of an ammonium cation, the analogous treatment with two different non-amino silanes was also investigated. These silanes were found to intercalate in the same manner, as indicated by an increase in the 00 I dspacing (Figure I, c). Moreover, intercalation Qf the
coupli~g
agent causes a decrease in the
breadth of the 00 I peak, indicating that silane treatment produces a single periodic structure. The silanes may complex with montmorillonite at two potential reaction sites. First at the edge of the platelets, hydroxyls are exposed and available for reaction (condensation) with an alkoxys~lane, by analogy with other hydroxyl surfaces such as silica
27.
The number
of these sites per gram of montmorillonite has been measured to be 3x10-seq/g .28 For a 200 nm square particle, this corresponds to approximately 1000 sites. Thus for the amino silane having molecular weight. 2I7, Xmax(edge)
=
0.65%. Second, sodium cations are present on the
platelet face·(- l.lxi0-3eq/g), where the silane may hydrolyze to the silanol, and may attach to the layer surface by hydrogen bonding, similar to various polar organic compounds.2s A third possibility is that the alkoxy silane may form a covalent Si-0 bond with the layer surface, perhaps at defect sites? 9 Whether covalent linkages can be established on the interlayer surface of montmorillonite has been long contested. Bergeno and Deuel31 were early proponents. who first converted montmorillonite to its acid (protonated) form, and then
treated it with organic molecules. Opponents suggested that the organic molecule was merely adsorbed to the surface. 32 More recently, Choudary et al. 33 treated protonated montmorillonite with 3-aminopropyl triethoxysilane, resulting in interlayer expansion. However, evidence for covalent bonding was not put forth until the Si NMR data of Mercier and Pinnavaia,34 which suggested the
covalent bonding of . 3-mercaptopropyltrimethoxysilane to protonated
montmorillonite. The results also suggested that though the silane was trifunctional, most of the silane was attached by only one Si-0-Si bond. Alternatively, the observed changes in intensity may be attributed to increased coupling, due to the addition of silanol species physically adsorbed to the montomorillonite interlayer. After treatment with the silane coupling agent, the amino functional clay was reacted with anhydride-functional polypropylene. DRIFT spectra (Figure 2) show the· conversion from polypropylene anhydride (bottom plot: anhydride and diacid at 1780 and 1715 cm- 1, respectively35) to amide groups (top plot: amide I at 1640 cm-1) following reaction with the amine-functional clay. Formation of the imide is also possible upon the removal of water, which may occur subsequently during melt compounding. Either amide or imide groups accomplish the desired covalent linkage. The tethering of these chains to the montmorillonite particle brings about a change in morphology, as determined by XRD (Figure 1, d). The new peak at~ 22
A indicates that the
grafted polymer is intercalated and has expanded considerably the silicate galleries. At low and intermediate concentrations of coupling agent and tethered chains, the peak at 13.4
A
remains as an indication that some particles are not intercalated with polymer. These particles either have not reacted with polymer, or they possess only chains that are grafted to the external surfaces and edges. Although these peaks correspond to d-spacings whose ratio is approximately I :2, we do not interpret the higher angle peak as a second order, but rather as a separate (initial) population for two reasons. First, the relative intensities of the two depend
on the concentration of polymer. Second, the higher angle peak coincides with the peak reported following treatment only with silane. Consequently, two morphologies are present, and the chain tethering reaction occurs non-uniformly. Based on the relative intensity of these two peaks, the relative fraction of the two morphologies can be estimated qualitatively, assuming equal structure factor per unit volume. As expected, the fraction of the polymer swollen product increases with increasing concentration of silane coupling agent and of grafted polymer. These changes are also accompanied by differences in the packing of the tethered chains. When the amount of tethered polymer is small (I 0% ), no crystalline polypropylene reflections are detected, indicating that crystallization of the grafted polymer within the narrow galleries ( 16.7
A in
thickness) may be inhibited. When the amount of tethered polymer is greater (50%), the usual· ( form of iPP is present (Figure 1; d). These crystals most likely reside outside of the silicate stack, because their size (73A, as measured by the Sherrer equation) is much larger than the silicate gallery thickness. The ability of these tethered chains to crystallize is desirable, because they can co-crystallize with matrix polymer, and potentially thereby increase. mechanical coupling and prevent fracture between the matrix and the nanofiller pm1icle.
Nanocomposite compounding An x-ray diffractogram of a 2/50/50 nanocomposite material (melt processed in the twinscrew microcompounder) is shown in Figure I, e. TI1e characteristic peaks of the ( phase of iPP are observed between 2\ values of 13 and 23 o, as expected. The weak peak at a d-spacing of approximately 13.4 A indicates that the un-intercalated morphology is present to some degree: nevertheless. the size of these clay aggregates is reduced. as indicated by the increased breadth of this peak. At lowest angle, the 001 peak associated with intercalated polymer has a mucl_1 greater intensity relative to the peak at 13.4A. The 13.4A peak is present
for all compositions containing a small amount of coupling agent (e.g., x = 2% ). When the amount of iPP is increased (i.e. decreased z), the relative intensity of the lowest angle peak decreases, suggesting a greater degree of exfoliation. The dispersion of the nanofiller particles was examined with respect to clay type, filler content, and processing method. Each of these variables was significant. The most uniform dispersions were obtained with fillers made with montmorillonite from Southern Clay, having significant amounts of tethered polymer (y
=
50%), and being compounded by
either the two-roll mill or the twin-screw microcompounder. Such samples are visually translucent and exhibit a fine dispersion of clay particles. (In contrast, samples having large aggregates of pmiicles (a few microns in size) are opaque.) TEM examination of a microtomed section of a (well-dispersed) 2/50/50 nanocomposite (Figure 3) indicates the presence of single silicate layers as well··as polymer-swollen stacks of several silicate layers, consistent with the low angle peaks observed by XRD (Figure L e). Smaller and fewer aggregates are observed at lower filler concentrations (e.g., z < 10).
Mechanical tests In typical composites, significant improvements in elastic modulus ( 1.5 times increase) and tensile strength (1.2X increase) are only achieved at high filler amounts (usually 30 to 50 percent by weight), which greatly impairs the ductility of the matrix polymer. The early failure of highly filled composites can be attributed to a variety of circumstances.36-38 Heavy inclusion of filler results in a smaller portion of the composite cross-sectional area occupied by the matrix, leading to a decreased potential for deformation. Early matrix failure is also caused by the development of cracks around filler particles and void-created aggregates following local particle/matrix debonding.36 In nanocomposites. smaller pmiicles and improved interactions limit
f~1ilure
by the aggregation or crack formation.
Figure 4 a shows the elastic moduli values for different composite formulations. An improvement in modulus (1.25X) was achieved at low filler content (1.0%) for all composites. This modulus increase is high compared to traditional filled polypropylenes: talc-filled PP achieved a 1.19 fold increase in modulus only at 30% filler loadings.38 This superior enhancement of modulus at low filler content has been reported for other nanocomposite systems. 4 The rapid modulus increase at 1.0% filler was followed by a steady increase with filler addition in all systems except those treated with siloxane dimer instead of silane coupling agent. The moduli of the dimer-composites returned to the matrix values at higher filler additions, which is below the lower limit defined by the law of mixtures for traditional composites with lamellar particles. Because the other formulations· (even composites of .. ·untreated c.laY and i-PP only, i.e. 0/1 00/z) do not exhibit this decrease, it is postulated that the . dimer has a detrimental effect on the modulus at higher filler levels. The dimer-composite modulus at 2-10 % is nearly the same as the pure matrix. a behavior· similar to traditional dispersion-filled (spherical or low aspect ratio) composites. 36 This ch~nge from a lamellarparticle to dispersion-filled mode of reinforcement is possibly the result of dimer-induced particle aggregation. As aggregation persists, the highaspect ratio morphology of the nanoparticle is lost and the reinforcing effect of the particles diminishes. The highest modulus value recorded for all composites at 10% filler loading was with the 2/50/10 (Figure 4 a). A more highly-filled nanocomposite (2/50/50, 25 percent inorganic by weight) achieved a two-fold increase in modulus, which was also greater than that reported for traditional mica-filled (1.8 fold increase at 45 weight%) and talc-filled (1.7 fold increase at 45 w·eight %) composites?) Composite tensile stress values also indicated trends in the different formulations (Figure 4 b). Yield strength increased slightly (1.15X) for all formulations at low filler
content (1.0%). Again, this yield strength increase in traditional systems is usually only achieved with much higher filler loadings: a 40% talc-filled PP showed a 1.06 fold increase,38 a 20% treated-micafilled PP showed a 1.06 fold increase, 40 and a 40% mica-filled PP gave a 1.2 fold increase. 38 Similar improvements at low filler fraction has been reported for other nanocomposite systems. 4 With increasing filler addition, the tensile strength steadily decreased for all composites except for those with a large fraction of silane-grafted chains (2/50/1 0). Nevertheless, the yield stress of the composites was at least as high as the yield stress ofthe matrix. Figure 5 shows the elongation at break of: the nanocomposite formulations. The silanecoupled chain-grafted nanocomposite (e.g., 2/80/1 0) showed much higher ductility retention (Lb = 320%) at high filler content comp?red to composites lacking one or more treatments. These results indicate that both the sifane coupling agent and anhydride polymer are beneficial to properties .. At 320% strain, the specimens were in the strain-hardening region of the stress/strain curve (ductile failure). This is at least as much ductility (possibly greater) as reported for traditional filled PP, which shows a brittle-to-ductile transition at 1015% filler.
38 41 42 • •
Fmthermore, at filler loadings req~ired to achieve comparable mechanical
property (E, {y) enhancements, the PP matrix is severely embrittled: the ductilities of 40% talc-filled and micafilled PP were reduced to O.os· and 0.02 times their original values, respectively. 38 While a marked difference existed between nanocomposites having silane-treated (x > 0) and untreated (x
= 0) fillers, the effect of silane concentration (x value) on the
mechanical properties was not significant. In terms of elastic modulus, yield strength, and ultimate elongation, there was no discernable difference between formulations with varying x values (x
=
2, 4, and 8). This suggests that each formulation achicwd the same particle
treatment, possibly indicating only external face and edge attachment. However, this was not experimentally verified.
ESC tests
Figure 6 shows in-situ pictures of crack progression in a sample of pure i-PP (Figure 6 ac) and of a 2/50110 nanocomposite (Figure 6 d-f). Immediately following placement of the test pieces under tension, a craze deformation zone was formed at the notch tip (Figure 6 a and d). This region has been observed in similar tests, 20.4 3 .44 where it was shown that the size ~
I .
of the craze region was determined by the stress level. The initial deformation caused a slight opening of the crack, followed by a lengthy period in which the crack opened at a very slow rate. During this 'slow' period (Figure 6 b and e), the formation and growth of secondary crazes 30 degrees above and· below the .. original cut (see arrows) was prevalent, as also observed elsewhere. 21 Eventually the crazes propagated across the length of the cross-section from both sides, forming a diffuse zone of deformation between the two side notches (Figure 6 c and f). Once this point was reached, rapid crack grow1h dominated the material, and failure occurred shortly thereafter. The major difference between the pure i-PP and nanocomposite samples is the time it takes to reach fully diffuse crazing. Compared to i-PP, the nanocomposite has a greater resistance to craze formation/propagation and as a result, much less deformation at a given time. Figure 7 is a plot of crack opening as a function of time, which shows a three-fold increase in time-to-failure for the 2/50110 nanocomposite compared to pure i-PP. Talc-filled PP test samples also show·ed slightly increased time-to-failure (5% talc averaged 393 hours,
30c% talc averaged 274 hours) compared to pure iPP, but not as great an increase as the nanocomposite sample. The resistance in the nanocomposites is most evident at shorter test
times, where initial crack propagation is much slower than i-PP (the time to reach a 1 mm crack opening is about five times longer for the nanocomposite ). The SEM photographs in Figure 8 confirm a brittle-like failure for both test materials in the ESC conditions. The fracture surfaces appear smooth, with no macroscopic ductile pull-out,
with
slight
local
microdeformation
(whitening)
as
observed
by
other
researchers.2 1.43 Since all samples were tested at the same the nominal stress and have the same crack morphology, the ESC tests were performed at equal stress intensity. The increased ESC resistance of the nanocomposite may arise from several potential factors. First, since each particle contains many polymer chain grafts, which incorporate in different crystals within the bulk, the filler may provide additional tie molecules to resist the ESCinduced disentanglement of crystallites. The nanoparticles may also act as nucleating agents and therefore influence crystallization· conditions. While this influence probably relates to the number of intercrystalline links, no such relationship has been established either theoretically or experimentally. 44 Besides providing intercrystalline links, a nanofiller particle may also act as a barrier to Igepal diffusion, w·hich is necessary for failure acceleration. Following the rapid initial craze formation, the rate of craze propagation is dependent on the diffusion of the ESC agent. 2 1.43 •45 The presence of the nanoparticles will increase the tortuous path of the ESC agent, impairing diffusion. This reduction of diffusivity, however, depends on the particle morphology (size, aspect ratio, and degree of layer exfoliation) and the nature of the interphase (barrier or channel-like). Igepal diffusion in these materials was not experimentally quantified.
Conclusions A nanofiller for polyolefins that increases enviro1m1ental stress cracking (ESC) resistance was prepared and characterized. The filler consists of inorganic silicate
montmorillonite nanoparticles that were first amine-functionalized with a silane coupling agent, and then tethered with maleated polymer chains. The reaction products were determined by DRIFT spectroscopy and by XRD analysis. XRD showed intercalation of both coupling agent and polymer and demonstrated the ability of the grafted polymer to crystallize. These particles were finely dispersed into unmodified iPP, using appropriate compounding methods. The mechanical properties and facture resistance of the resulting nanocomposites were investigated. The best mechanical properties (modulus, strength and ductility) were obtained from formulations having both .silane treatment and functional polymer (silanecoupled chains); nanocomposites formed without silane and/or functional polymer had a distinct deficiem:y in one or more mecha!lical properties. Improved ESC resistance was exhibited by the nanocomposite. The increased resistance of the nanocomposites was mo.st evident at shorter test times, where initial crack propagation was ·five times slower in nanocomposite than in i-PP. Nanofiller particles may contribute to ESC resistance by introducing more effective tie chains or reducing diffusion of the stress-cracking agent Igepal.
Acknowledgements
The financial support ofNSF grant EEC-9108700 and PRF grant 31333-G7 is gratefully acknowledged.
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Figure Captions Figure 1. X-ray diffractograms ofvarious samples. a.) Sodium montmorillonite as received; b.) montmorillonite as in (a) treated with x = 4 wt.% aminopropyl diisopropyl ethoxy silane; c.) montmorillonite as in (a) treated with x = 10 wt.% diphenyl methyl ethoxy silane; d.) Na-montmorillonite treated with x = 2 wt.% amino silane andy= 50 w1.% MAPP; and e.) nanocomposite comprising iPP and montmorillonite with x = 2% amino silane, y =50 wt.% MAPP, and z =50 wt.%. See· text for definition of quantities x, y, and z. Figure 2. DRIFT spectroscopy. Upper curve: anhydride polypropylene. Lower curve: Nanofiller comprising montmorillonite, amino silane (x = 2 wt.%), and MAPP (y =50 w1.%). Figure 3. The morphology of a 2/50/50 nanocomposite, showing the presence of both single Clay layers as well as swollen stacks of several clay layers. Figure 4. Tensile properties ofnanocomposites as a function of filler content for different nanocomposite formulations. a.) Elastic moduli; b.) tensile strength. Figure 5. Nanocomposite elongation at break as a function of filler content for different nanocomposite fonnulations. Figure 6.Jn situ photographs of ESC test samples. a- c.) i-PP sample at the times indicated. d
f.) 2/50/10 nanocomposite sample at the times indicated. Magnification calibration can be clone using the metal tab at the left, which is 3.38 mm wide.
Figure 7. The time-to-failure plot for ESC test samples, pure isotactic PP and 2/50110
nanocomposite; crack opening v. time. Figure 8. SEM photographs of fracture surfaces for a.) the pure i-PP sample and b.) the 2/50110 nanocomposite.
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