Multilayered confinement of iPP/TPOSS and nylon 6/APOSS blends Matthew M. Herbert, Ricardo Andrade, Hatsuo Ishida, João Maia, David A. Schiraldi * Department of Macromolecular Science and Engineering, Case Western Reserve University, Cleveland, OH 44106, United States article info Article history: Received 7 September 2013 Received in revised form 28 October 2013 Accepted 1 November 2013 Available online 8 November 2013 Keywords: POSS Blends Polypropylene Nylon abstract A series of isotactic polypropylene and nylon 6 blends with silsesquioxane (POSS) additives were pro- duced, then layered to nanometer thicknesses to test the effects of confinement upon polymer property modification. POSS is shown to be a poor filler, lacking solubility and favorable interaction with the polymer matrices. It was initially hypothesized that under extreme confinement and orientation, such as in melt-spun fibers, or confined within 2D nanoscale layers, that POSS would undergo forced-assembly into elongated, rebar-like reinforcement structures, or even act as crosslinking molecules for the polymer chains. The current results, however, show POSS existing as large, phase separated aggregates, in order to minimize interactions with the polymer matrix; the aggregates behave as debonded hard particles upon tensile deformation. POSS has been previously shown to enhance the properties of polymer matrices in which the POSS molecules have been grafted to, or copolymerized within the chain, but this is not the case for these POSS blends. In comparison to results from the iPP/DBS/TPOSS system, in which POSS is unable to directly interact with the polymer matrix, and the nylon 6/APOSS system, in which POSS can potentially form hydrogen bonds with the polymer matrix, the results are similar and reveal that POSS blends are largely incompatible with the polymer matrix. Small improvements in blend properties can be made via functionalization of the POSS cage, in order to enhance interactions, but these improvements are quite limited. Even under extreme confinement conditions, where significant deviations from bulk behavior occurs, such as within melt-spun fibers or nanoscale layers, POSS was unable to be forcibly-assemble into reinforcing structures with the polymer matrix. Ó 2013 Elsevier Ltd. All rights reserved. 1. Introduction The reinforcement of polymers attempts to combine the ad- vantageous processing properties of organic materials and the desired characteristics of specific fillers with the end result typi- cally being a volumetric average of contributions from the indi- vidual components [1]. Dramatic improvements in composite systems have been achieved using a relatively low amount of fillers; this can be more easily attained as the filler size is decreased to the micro- and nano-scale, taking advantage of enhanced surface properties and increased interfacial interactions [1]. As filler size is reduced to the nanoscale, however, differences in surface energy become drastically more significant, often leading to aggregation of nanoparticles. The ability of such fillers to disperse within a poly- mer matrix is vital to the final composite properties, leading to many approaches which have been used to increase the quality of such dispersions, including the incorporation of dispersants, grafting onto polymer chains or even copolymerizing fillers directly into the desired polymer chains [1,2]. One of the more widely studied and used families of nanoscale fillers/additives are the silsesquioxanes, also known by the com- mercial name POSS Ò (polyhedral oligomeric silsesquioxane). POSS molecules are hybrid caged molecules related to silicones, with a chemical structure (RSiO 1.5 ) consisting of an inorganic core (SiO 1.5 ) and organic side groups (eR). As a result of this structure, POSS can be incorporated into polymer systems in a number of ways including blending, grafting and copolymerization [3e11]. Many variations of POSS molecules can be synthesized and are illustrated in Fig. 1 , including non-caged structures such as random and ladder structures, caged structures such as the T8, T10 and T12 geometries, and partial caged structures which have one or more open corners [1]. The potential advantages gained from the incorporation of POSS come from its hybrid organic-inorganic nature, wherein the inorganic core provides molecular reinforcement, while its organic periphery allows for covalent or non-covalent interactions with the host polymer. By modifying the type of cage and surrounding organic groups, the properties of POSS can span the gap from inorganic materials to organic substances [12]. Aside from the * Corresponding author. Tel.: þ1 216 368 4243. E-mail addresses: [email protected], [email protected](D.A. Schiraldi). Contents lists available at ScienceDirect Polymer journal homepage: www.elsevier.com/locate/polymer 0032-3861/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.polymer.2013.11.001 Polymer 54 (2013) 6992e7003
0032-3861/$ e see front matter � 2013 Elsevier Ltd.http://dx.doi.org/10.1016/j.polymer.2013.11.001
a b s t r a c t
A series of isotactic polypropylene and nylon 6 blends with silsesquioxane (POSS) additives were pro-duced, then layered to nanometer thicknesses to test the effects of confinement upon polymer propertymodification. POSS is shown to be a poor filler, lacking solubility and favorable interaction with thepolymer matrices. It was initially hypothesized that under extreme confinement and orientation, such asin melt-spun fibers, or confined within 2D nanoscale layers, that POSS would undergo forced-assemblyinto elongated, rebar-like reinforcement structures, or even act as crosslinking molecules for the polymerchains. The current results, however, show POSS existing as large, phase separated aggregates, in order tominimize interactions with the polymer matrix; the aggregates behave as debonded hard particles upontensile deformation. POSS has been previously shown to enhance the properties of polymer matrices inwhich the POSS molecules have been grafted to, or copolymerized within the chain, but this is not thecase for these POSS blends. In comparison to results from the iPP/DBS/TPOSS system, in which POSS isunable to directly interact with the polymer matrix, and the nylon 6/APOSS system, in which POSS canpotentially form hydrogen bonds with the polymer matrix, the results are similar and reveal that POSSblends are largely incompatible with the polymer matrix. Small improvements in blend properties can bemade via functionalization of the POSS cage, in order to enhance interactions, but these improvementsare quite limited.
Even under extreme confinement conditions, where significant deviations from bulk behavior occurs,such as within melt-spun fibers or nanoscale layers, POSS was unable to be forcibly-assemble intoreinforcing structures with the polymer matrix.
� 2013 Elsevier Ltd. All rights reserved.
1. Introduction
The reinforcement of polymers attempts to combine the ad-vantageous processing properties of organic materials and thedesired characteristics of specific fillers with the end result typi-cally being a volumetric average of contributions from the indi-vidual components [1]. Dramatic improvements in compositesystems have been achieved using a relatively lowamount of fillers;this can be more easily attained as the filler size is decreased to themicro- and nano-scale, taking advantage of enhanced surfaceproperties and increased interfacial interactions [1]. As filler size isreduced to the nanoscale, however, differences in surface energybecome drastically more significant, often leading to aggregation ofnanoparticles. The ability of such fillers to disperse within a poly-mer matrix is vital to the final composite properties, leading tomany approaches which have been used to increase the quality ofsuch dispersions, including the incorporation of dispersants,
du (D.A. Schiraldi).
All rights reserved.
grafting onto polymer chains or even copolymerizing fillers directlyinto the desired polymer chains [1,2].
One of the more widely studied and used families of nanoscalefillers/additives are the silsesquioxanes, also known by the com-mercial name POSS� (polyhedral oligomeric silsesquioxane). POSSmolecules are hybrid caged molecules related to silicones, with achemical structure (RSiO1.5) consisting of an inorganic core (SiO1.5)and organic side groups (eR). As a result of this structure, POSS canbe incorporated into polymer systems in a number of waysincluding blending, grafting and copolymerization [3e11]. Manyvariations of POSS molecules can be synthesized and are illustratedin Fig. 1, including non-caged structures such as random and ladderstructures, caged structures such as the T8, T10 and T12 geometries,and partial caged structures which have one or more open corners[1]. The potential advantages gained from the incorporation ofPOSS come from its hybrid organic-inorganic nature, wherein theinorganic core provides molecular reinforcement, while its organicperiphery allows for covalent or non-covalent interactions with thehost polymer. By modifying the type of cage and surroundingorganic groups, the properties of POSS can span the gap frominorganic materials to organic substances [12]. Aside from the
Fig. 1. Variety of POSS geometries including: A) Random; B) Ladder; C) Partial Cage; D) T8; E) T10; F) T12 [1].
M.M. Herbert et al. / Polymer 54 (2013) 6992e7003 6993
synthesis and preparation of POSS molecules, the efficiency of theirphysical properties is primarily dependent on their abilities tohomogeneously disperse and interact with the polymer matrix,which in turn is dependent on the thermodynamic interactionsbetween POSS and polymer. [9,13].
A great deal of research has gone into studying the effects ofPOSS which has been grafted onto [3,8] or copolymerized [4e7]into a polymer chain. Less is understood of the behavior and in-teractions of POSS blended into a polymer system, and how toobtain successful reinforcement or polymer vitrification when us-ing POSS as an additive. Many studies have examined blendingPOSS into different polymer systems, but most have included highconcentrations of POSS upwards of 10 wt% [14e18]. Many of theseprevious studies have also deemed POSS as a structurally ineffec-tive filler, due to the resultant high degree of aggregation, whichsignificantly reduces the properties of the blend. What thesestudies have overlooked is the low critical solubility ratio of POSSwithin these systems. Recent work has explored the solubility pa-rameters of POSS-polymer combinations, in an effort to makepredictions about dispersion and compatibility [19e21]. Thesesolubility parameter methods produce inaccurate results, andignore contributions from hydrogen bonding, as well as contribu-tions from the core silicate group [13]. The optimal solubility ofPOSS additives in specific polymer matrices can be better deter-mined using Hansen solubility parameters, which includes contri-butions from dispersion, polar and hydrogen bonding forces, andmore accurately predict POSS-polymer interactions and POSSdispersion uniformity, compared with group contribution calcula-tions [13].
The enhancement of melt-spun fiber properties of amino-propylisobutyl POSS (APOSS)/nylon 6 blends was recently reported[22]. A “sweet spot” of POSS solubility in the polymer matrix existsat a low 2.5e3.0 wt%, where a significant increase in the Young’smodulus and yield stress is observed. This peak in mechanicalproperties was seen to only occur in highly elongated melt-spun
fibers, and not in bulk injection molded samples. At low APOSSconcentrations, small elongated aggregates were formed in theaxial direction, although molecular dispersion of APOSS was re-ported as well. As the concentration of POSS was increased past thecritical concentration, larger spherical aggregates began to form,resulting in POSS acting primarily as a processing aid with theattendant decrease in the blend’s mechanical properties [22]. It washypothesized that under extreme confinement and orientation,such as in melt-spun fibers, that POSS would undergo forced-assembly into elongated, “rebar-like” reinforcement structures.Experimental evidence was inconclusive, however, as to themechanism of reinforcement.
Based on the results from the Hansen solubility parametersprediction of a high degree of compatibility between filler andmatrix [13], and the previous work with melt-spun APOSS/nylon 6fibers [22], two POSS-polymer combinations were chosen for thepresent work, in which to explore the effects of 2D layered nano-confinement on the resultant structureeproperty relationships ofthe blends. To produce a system under 2D confinement, layer-multiplying coextrusion technology (Fig. 2) was used in whichfilms containing up to thousands of layers can be continuouslyproducedwith individual layer thicknesses ranging from themicro-to the nanoscale [23]. The first system of trisilanolphenyl POSS(TPOSS), blended into isotactic polypropylene (iPP) along with thenucleating agent dibenzylidene sorbitol (DBS), represents a systemwith no direct interactions between POSS and polymer. Instead, thenucleating agent (DBS) is used as a dispersant and compatibilizerfor TPOSS within iPP [24e26]. The second system of APOSS,blended into nylon 6,represents a system in which direct in-teractions between POSS and polymer, via hydrogen bonding, ispossible. Comparison of these two systems under multilayerednano-confinement will be reported herein, in an effort to shedmore light on the self- and forced-assembly behavior of POSS-polymer blends, in which POSS is able to interact directly andindirectly with the polymer matrix.
Fig. 2. Layer-multiplying coextrusion technology (reprinted courtesy of the publisher) [23].
M.M. Herbert et al. / Polymer 54 (2013) 6992e70036994
2. Experimental
2.1. Materials
Isotactic Polypropylene (Pinnacle PP 1703) was obtained in theform of pellets from Entec Polymers, LLC with MFI 3.5 g/10 min(ASTM D1238), density 0.9 g/cm3 (ASTM D1505). Nylon 6 resin(Grilon� FG40-NL) was obtained in the form of pellets from EMS-Grivory with MVR 20 g/10 min (ISO 1133), density 1.14 g/cm3 (ISO1183). Two grades of POSS were obtained from Hybrid Plastics(Hattiesburg, MS). Trisilanolphenyl POSS (TPOSS, SO1458) andAminopropylisobutyl POSS (-POSS, AM0265) were both available inthe formof awhitepowderandwereusedas received.Dibenzylidenesorbitol (DBS), a dispersion aid and nucleation agent, was obtainedfrom Milliken Chemicals (Spartanburg, SC) as Millad 3905, with amolecular weight of 358.4 g/mol and a Tm of 225 �C Fig. 3(AeC)presents the chemical structures of TPOSS, APOSS and DBS,respectively.
2.2. Preparation of POSS blends
All components were dried in a vacuum oven at 100 �C for 24 hprior to blending. Two POSS-polymer compositions were weighedand dry-mixed prior to blending. Blend A consisted of isotacticpolypropylene, 3.0 wt% TPOSS and 0.3 wt% DBS. Blend B consistedof nylon 6 blended with 2.5 wt% and 5.0 wt% APOSS. Blend com-positions were selected for optimal solubility in their respectivepolymers as determined reported previously [13]. Mixtures weremelt-bended using a Haake twin screw extruder at 190 �C, for blend
Fig. 3. Molecular structures of A
A, and at 240 �C, for blend B, at a rotor speed of 20 rpm. Extrudatewas collected and pelletized in preparation for multilayerextrusion.
2.3. Film preparation
Prepared blends were fed through a multilayered extrusionprocess, which utilizes a process of forced-assembly throughsequential layer multiplication to fabricate thin alternating layers oftwo or three separate polymers [21]. Multilayered films wereextruded in an (AB)n layered structure, containing TPOSS and DBSfor blend A, and APOSS for blend B, in alternating layers of 256 and1024 total individual layers. For each of the two POSS-polymersystems, neat multilayered control films were also prepared. Inaddition to the blend A system, three multilayered control filmswere produced containing 1024 layers; Neat iPP with no fillers, iPPwith 0.3 wt% DBS in alternating layers and iPP with 3.0 wt% TPOSSand 0.3 wt% DBS in all layers. In addition to the blend B system, amultilayered neat nylon 6 film was produced containing 1024layers. Films were extruded at 255 �C and 245 �C, for blends A andB, respectively, at a screw speed of 10 RPM. Resultant extrudedfilms were cold drawn on a chill roll to thicknesses ranging from 12mil (305 mm) to 1 mil (25 mm), producing individual layer thick-nesses from 1200 nm to 25 nm.
2.4. Biaxial orientation
Square specimens 85 mm � 85 mmwere cut from the extrudedfilms, marked with grid patterns, and biaxially stretched in a
) TPOSS; B) APOSS; C) DBS.
M.M. Herbert et al. / Polymer 54 (2013) 6992e7003 6995
Bruckner Karo IV biaxial stretcher at 162 �C and 210 �C, for A and Bfilm types, respectively, with preheat time fixed at 1 min. Engi-neering strain rates of 400% s�1 were used to simultaneously andequi-biaxially draw the square specimens to draw ratios of 3 � 3and 4 � 4. The uniformity of the drawn specimens was determinedfrom the even deformation of the grid patterns.
2.5. Mechanical properties
In order to determine the mechanical properties of the extrudedmultilayered films, ASTM type V tensile bars were cut and tested onan Instru-Met Corporation 1 kN load cell using Instron pneumaticthin film grips. A strain rate of 11 mm/min was used and theresulting tensile datawas analyzed onMTS Testworks software. Thefilms were characterized in terms of Young’s modulus (MPa) andyield stress (MPa), with respect to individual layer thicknesses.
2.6. Image analysis
The morphology of the POSS-polymer multilayered films wereobserved edge-on by scanning electron microscopy (SEM) usingJOEL JSM-6510 scanning electron microscope. In order to acquireedge-on images, films were embedded between two thin layers ofepoxy, to provide structural stability, and microtomed at �100 �Cusing a Leica EM FC6microtome to provide a smooth homogeneoussurface. Films were sputter coated with a thin layer of gold prior toimaging.
Analysis of grain size was observed using an Olympus BH2polarized optical microscope.
2.7. Thermal analysis
The thermal behavior of the POSS-polymer blends at variouslayer thicknesses was investigated using a DSC200 (Mettler Toledo)differential scanning calorimeter (DSC) under continuous nitrogenpurge, at a flow rate of 60 ml/min. The thermal properties such asmelting temperature (Tm), glass transition temperature (Tg), andenthalpy change (DHm) in melting were determined at a heatingrate of 10 �C/min from 25 to 250 �C.
2.8. Rheological analysis
A Paar Physica 501 rheometer from Anton Paar, operated with a25 mm parallel plate setup was used to measure the rheologicaldynamic properties of the nylon 6 and nylon 6/APOSS blends at 240and 245 �C. Disk-shaped specimens of 25 mm diameter and 1 mmthickness were molded at 240 �C in a compression mold. The
Table 1iPP/DBS/TPOSS tensile data, where B.O. indicates biaxially oriented samples. *DBS 12 and 8and TPOSS control films were unable to be biaxial stretched to 4 � 4 draw ratios.
specimens were subjected to 240 �C for 5 min during molding. Thepolymer compound in the rheometer was first heated and kept atthe desired temperature for 2 min to reach equilibrium. Then,frequency sweep at constant temperature with oscillatory shearfrequency between 1 and 100 rad/s was performed. A linearviscoelastic regime was independently confirmed under theseconditions.
2.9. WAXS analysis
Wide angle X-ray diffraction (WAXS) of the POSS-polymerblends were investigated using a Brukner wide angle X-ray in-strument with sealed tube-ray generator and Cu-Ka radiationproducing an incident wavelength of 1.54 Å. The WAXS patternswere examined with respect to layer thickness in order to observeany changes in crystal structures with increased confinement oflayers.
2.10. Raman spectroscopy
MicroRaman scattering studies were performed at room tem-perature with a Horiba Jobin Yvon LabRam HR800 spectrometerequipped with a charge coupled detector and two grating systems(600 and 1800 lines/mm) A He-Ne laser (l¼ 632.8 nm)was focusedon the sample with an Olympus microscope at an optical power of17mWand a spot size of 1 mm2. Raman shifts were calibratedwith asilicon wafer using the 520 cm�1 line.
3. Results and discussion
3.1. TPOSS/DBS/Ipp blend films
Multilayered films containing an alternating layer structure(AB)n of 256 and 1024 layers, with A layers consisting of 3.0 wt%TPOSS and 0.3 wt% DBS in iPP, and B layers consisting of 0.3 wt%DBS in iPP, were successfully produced using the processingequipment illustrated in Fig. 2. In addition to the alternatinglayered films, TPOSS, DBS and iPP control films with similar layerstructures were extruded containing 3.0 wt% TPOSS and 0.3 wt%DBS in every layer, 0.3 wt% DBS in every layer, and neat iPP withoutany additives, respectively.
Mechanical tensile testing was performed on samples with totalfilm thicknesses of 12, 8, 4, 2 and 1mil (1mil¼ 25.4 mm), in order toorder to examine the effects of decreasing film and layer thicknessfor each of the film types. Testing was also performed on biaxiallystretched samples (3 � 3 and 4 � 4 draw ratios). Data collectedfrom tensile testing (Table 1) shows a significant increase in the
mil films contained significant processing defects preventing tensile analysis. **DBS
Fig. 4. SEM image of debonded and voiding POSS aggregates.
Fig. 5. DSC melting peaks for TPOSS film series; A) 12 mil and B) 1 mil thicknesses.
M.M. Herbert et al. / Polymer 54 (2013) 6992e70036996
Young’s modulus and yield stress for films containing TPOSS andDBS, as layer thickness decreased below 100 nm (4 mil film thick-ness) in 256 layered films, and below 50 nm (2 mil film thickness)in 1024 layered films. Neat multilayered iPP films exhibit muchlower mechanical properties relative to those containing TPOSSand/or DBS. For TPOSS/DBS containing films, as the as-extrudedfilm thickness was reduced from 12 mil (304.8 mm) to 1 mil(25.4 mm), Young’s modulus rose by an average of 30% for all filmtypes, and by over 100% in the 256 layered films. The yield stressdata followed a similar trend, increasing with decreasing film andlayer thickness. As the layer thickness was reduced below 200 nm,yield stress was observed to increase by 100% for all TPOSS/DBScontaining film types, from approximately 40 MPa to 80e90 MPa.The layered iPP control remained relatively constant at approxi-mately 40 MPa at all thicknesses tested.
Biaxially oriented films, which by nature of the stretching pro-cess, have decreased film and layer thicknesses, exhibited evengreater increases in the tensile properties than the unorientedsamples. Young’s moduli increased by 10e50% upon drawing for allPOSS containing films, while neat iPP and DBS control films did notexhibit such increases. Yield stress increased by approximately 50%for 1024 and 256 layered POSS films, but no increases were seen forthe control films.
The polar nature of TPOSS prevents it from dispersing well inpolypropylene, which previously led to the use of DBS as adispersing agent for TPOSS in iPP [26]. It was demonstrated thatDBS and TPOSS were capable of forming several complex molecularadducts through hydrogen bonding and pep stacking, but no co-valent interactions occurred [26]. It was also revealed that TPOSSprefers to form covalent inter- and intra-molecular bonds betweensilanol groups, which results in large phase separated aggregatesranging from less than 1 mm up to 20 mm in diameter. As evidencedfrom the SEM image in Fig. 4, these aggregates are fully debondedfrom the iPP matrix. Upon tensile deformation, void formationoccurs due to cavitation localized at particleematrix interfaces.Deformation of the polymer matrix around debonded particlesoccurs parallel to the loading direction on both sides of the TPOSSaggregate, leading to plastic deformation through shear yielding[27]. Thermal analysis of the TPOSS/DBS films further confirms thatTPOSS aggregates are non-interacting with the iPP matrix, asindicated by a lack of change in the melting behavior of iPP (Fig. 5)between controls and TPOSS containing films, with respect todecreasing layer thickness. Consistency of the melting tempera-tures indicates that TPOSS and DBS are not covalently interactingwith the iPP matrix. The presence of shoulders at around 160 �C inboth 8mil and 1mil samples, as well as the double peak in the 1miliPP control, can be explained as a disorder-order transition. Thistransition represents partial melting of iPP daughter- and motherlamellae formed through cross-hatching [28].
Secondary strengthening effects of the polymer matrix, such asmicro-void formation around debonded TPOSS aggregates leadingto shear banding and yielding, can provide small enhancementsprimarily to the films’ tensile modulus [29]. We believe that theprimary strengthening mechanism responsible for increasing yieldstress and modulus, with respect to decreasing layer thickness,comes from reduction of the average iPP spherulite size andorientation of the polymer chains. Changes in spherulite size areinduced by additional nucleation sites provided by DBS and TPOSSaggregates. Specifically, the formation of fibrillar networks by DBSfrom the melt, through hydrogen bonding and pep stacking in-teractions, provides additional nucleation sites for iPP crystalgrowth, and is commonly used in industry to produce transparentpolypropylene. Spherulitic size was determined using polarizedoptical microscopy (POM), which showed that films containing
M.M. Herbert et al. / Polymer 54 (2013) 6992e7003 6997
POSS and DBS contain a smaller average spherulitic diameter,compared with neat iPP films. It was previously demonstrated, thatin the presence of TPOSS, the fibrillar concentration formed fromDBS is reduced, due to complex hydrogen bonding interactionsbetween the two molecules [30]. Thus the ability of DBS fibrils toprovide additional nucleation sites is hindered above a criticalTPOSS/DBS ratio [26]. Films containing only DBS have a smalleraverage spherulitic size than films containing both DBS and TPOSS,which in turn have smaller spherulites than neat iPP, as is shown inFig. 6.
In neat polypropylene, which does not contain any nucleatingagents or fillers, larger spherulites result in higher modulus andlower yield stress values [31]. Samples containing smaller spher-ulites deform by necking and cold drawing, followed by fibrillarfracture, leading to an increase in yield stress as observed in thecurrent system, but larger spherulites tend to undergo brittlefracture [31]. Higher modulus values are due to less volume occu-pied by amorphous grain boundaries between spherulites, whichare weaker and more easily deformed than the crystalline phase[31]. This behavior is noticeably different from the data presentedin Table 1, which shows an increase in modulus with respect todecreasing spherulite size, perhaps resulting from the presence ofDBS fibrils, or TPOSS aggregates collected within crystalline regionsforming microvoids upon deformation. For a larger averagespherulite size in thicker films, less amorphous volume leads to ahigher TPOSS concentration within the crystalline regions. Upondeformation, the small inter-aggregate distances in highlyconcentrated areas can lead to the lateral coalescence of micro-voids, reducing the samples ability to withstand stress loads [32].As the average spherulitic size is reduced, higher amorphous vol-umes allow a better dispersion of TPOSS aggregates into amor-phous regions, lowering TPOSS concentrations in crystalline areas.
This trend of increasing yield stress with respect to decreasingspherulite size is plotted in Fig. 7 to demonstrate a phenomenonsimilar to the HallePetch relationship [33]. Microvoids formed inareas of low TPOSS aggregate concentrations, at amorphous grainboundary sites, can act as barriers to chain movement and canreduce mobility of the amorphous phase. It is shown in Fig. 7,compared with neat multilayered iPP, films containing POSS/DBSsee a significant increase in yield stress with reduction in spherulitesize. This trend is even more prominent in the film system con-taining TPOSS in all layers (TPOSS control), relative to TPOSS inalternating layers.
Decreasing film and layer thickness is also responsible forfurther decreases in spherulite size due to the confinement ofspherulites within individual layers. Thus as the layer thickness isreduced below a critical spherulite thickness, the spherulites arereduced in size, until a point where the film becomes highly
0
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0 50 100 150 200 250 300 350
PP
POSS
DBS
256
1024
Spherulite Size (µm) vs. Film Thickness (µm)
FilmThickness (µm)
Ave
rage
Sph
erul
ite
Size
(µm
)
Fig. 6. Plot of decreasing spherulite (grain) size (mm) with respect to film thickness(mm).
amorphous in biaxially oriented samples. Thicker 12 mil filmsshowed an average spherulite size of 17e19 mm,whichwas reducedto around 10e12 mm for 1 mil films (Fig. 6). The results presented inFig. 6 of average grain size vs. film thickness illustrates a trend ofdecreasing spherulite size with decreasing film thickness, due toreduction of individual layer thicknesses below a critical spherulitesize. Based on POM images shown in Fig. 8, biaxially orientedsamples, stretched to 3� 3 and 4 � 4 draw ratios from 12 mil and 8mil thick films, exhibit an increased amount of amorphous regions.As a result of the biaxial orienting process, iPP chains becomeelongated and oriented parallel to the loading directions, dimin-ishing the spherulite size and concentration. The tensile propertiesof the biaxially oriented films, presented in Table 1, show a signif-icant increase in modulus for TPOSS containing films, and in yieldstress for all non-control films. This selective increase in tensileproperties demonstrates a shift in the primary deformationmechanism from spherulitic dominated behavior to behaviordominated by the quantity of TPOSS present in the system; thereexists a critical solubility concentration of TPOSS in iPP [13]. Too lowa concentration of TPOSS, and it has negligible effects on the iPPmatrix, too high a concentration of TPOSS, and it phase separatesout of iPP in the form of large aggregates. These larger and moreconcentrated aggregates can have a large effect on the iPP matrix,where increased void size and small inter-aggregate distances inhighly concentrated areas can lead to the lateral coalescence ofmicrovoids, reducing the samples ability to withstand stress loads[32]. This effect can be observed in the yield stress properties,where films containing an optimal wt% of TPOSS (3.0 wt%) seefurther increases in yield stress, while films containing higherconcentrations of TPOSS, or no TPOSS at all, see either no change ora decrease in the yield stress. Increases in the modulus of biaxiallyoriented films can be primarily attributed to increased orientationof the iPP chains in the loading direction, and a shift in the me-chanical properties from ductile to brittle films. As the chainsbecome increasingly oriented, there is much less chain movementupon deformation, and crazing, rather than shear yielding, be-comes the main deformation mechanism.
To further probe the effect of TPOSS and DBS on the reinforce-ment of multilayered iPP films, wide angle X-ray diffraction (WAXS)spectroscopy was performed on all film samples (Fig. 9). Theseresults clearly indicate that iPP crystallizes into an a-polypropylenecrystal structure for all film types. The a phase is characterized by amonoclinic unit-cell crystal structure (a ¼ 6.66; b ¼ 20.78;c ¼ 6.495�A and b ¼ 99.6�) and is the most common crystal form iniPP [28]. It can be observed that the only peaks present in thespectra belong to the a phase form of iPP, and no additional peaksarise due to the incorporation of TPOSS or DBS into the system.
0
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PP
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Spherulite Size (µm) vs. Yield Stress (MPa)
Yie
ld S
tres
s (M
Pa)
Spherulite Size (µm)
Fig. 7. Plot of increasing yield stress (MPa) with respect to decreasing spherulite(grain) size (mm) illustrating a similar phenomenon to the HallePetch relationship.
Fig. 8. POM images demonstrating the size decrease and elongation of spherulites with respect to decreasing layer thickness. A) 200 mm thick as-extruded film; B) 50 mm thick as-extruded film; C) 20 mm thick biaxially oriented film.
M.M. Herbert et al. / Polymer 54 (2013) 6992e70036998
For as-extruded film samples, a gradual decrease in peak heightis observedwith respect to decreasing film and layer thickness. Thispeak decrease can be attributed to the reduction of spherulite size,and increased volume of amorphous iPP. Biaxially oriented filmsamples show further reduction in the a-iPP spectrum, except forthe 110 peak, in which peak height increases for 3 � 3 drawnbiaxially oriented samples. This peak is representative of the peri-odicity along the chain length, and can be attributed to increasedorientation of the chains in the loading direction after drawing. Thereduction in WAXS peak heights, and thus crystallinity in thesample, was also verified using DSC. In this method, crystallinitycan be calculated from the heat of enthalpy (area under the meltingpeak), using the equation in the inset of Fig. 10. The trend ofdecreasing crystallinity with respect to decreasing layer thickness,shown in Fig. 10, is consistent with our earlier conclusions.
3.2. APOSS/nylon 6 blend films
A second system of multilayered films, consisting of amino-propylisobutyl POSS (APOSS) melt blended into nylon 6, was
Fig. 9. WAXS plots showing the a-PP crystal structure spectrum with respect to; A)decreasing film thickness in as-extruded samples, and B) decreasing film thickness inbiaxially oriented samples.
produced using the processing equipment illustrated in Fig. 2, andcharacterized as a function of POSS concentration and layer thick-ness. Five film sets were prepared, varying the weight percent ofAPOSS around a critical concentration of 2.5 wt%, from 0 to 5 wt%,and varying the overall number of layers within the films; 1024layered films of neat nylon 6 (0 wt%), 2.5 and 5.0 wt% APOSS inalternating layers, and 256 layered films of 2.5 and 5.0 wt% APOSSin alternating layers. For each of the film sets, variation of the layerthickness was accomplished by changing the overall thickness ofthe films. Film thicknesses of 12, 8, 4, 2 and 1 mil (1 mil ¼ 25.4 mm)were extruded. Biaxially oriented samples were also prepared fromthe extruded films by drawing 12mil and 8mil thick films to a 3� 3draw ratio, to produce highly oriented layer thicknesses between34 and 22 nm.
Tensile testing was performed on all samples, which werecharacterized in terms of their Young’s modulus (MPa) and yieldstrength (MPa). In similar behavior to the TPOSS/DBS/iPP films,tensile data presented in Table 2 shows a gradual increase in tensileproperties with respect to decreasing layer thickness below 100e200 nm. In comparing the mechanical property trends of filmsbetween 12 and 1 mil thicknesses, the elastic modulus shows aminimum increase of around 60e75% for 1024 layered films, fromaround 500 MPa up to 850 MPa, and a maximum increase of 85e160% for 256 layered films, from around 600 MPa up to 1450 MPa.Control films, containing no APOSS, likewise saw a significant in-crease in the elastic modulus by over 100%, from 370 MPa up to775MPa. Biaxially oriented films displayed even further escalationsin the elastic modulus for 1024 layered APOSS containing films byas much as 140e175%, to above 2000MPa, compared to 250%, up to2750 MPa, for the nylon 6 control.
Fig. 10. DSC determination of crystallinity for 256 layered film samples with respect todecreasing layer thickness (nm). Crystallinity was calculated from first heating curvesusing the inset equation.
Table 2Nylon 6/APOSS tensile data, where B.O. indicates biaxially oriented film samples.
305 mm 203 mm 102 mm 50 mm 25 mm B.O. (3 � 3) from305 mm to 34 mm
B.O. (3 � 3) from203 mm to 23 mm
Modulus (MPa) vs. film thickness (mm) and number of layers1024 layersNylon 6 control
M.M. Herbert et al. / Polymer 54 (2013) 6992e7003 6999
Yield stress results exhibit slightly different trends comparing12 to 1 mil film thicknesses, where a maximum increase of 60e70%occurs in APOSS containing 256 layered films, from 40 to 45MPa upto 70e75 MPa. In 1024 layered films, a smaller increase in yieldstress of around 20 and 40% is seen for films containing 2.5 and5.0 wt% APOSS respectively, from around 40 MPa up to 50 MPa.Control films see a smaller increase in yield stress by around 15%,from 41 to 48 MPa. However, there is no clear correlation in yieldstress with respect to APOSS wt%, as seen in the as-extrudedYoung’s modulus properties. Biaxially oriented films exhibit con-trasting trends between 256 and 1024 layered films, where 256layered films see a significant decrease in yield stress by 60 and30%, for films stretched to a 3 � 3 draw ratio from 12 mil and 8 milthick films, respectively. On the other hand, biaxially oriented 1024layered films see an increase in yield stress by 30 and 60%, for 2.5 wt% APOSS stretched from 12 mil and 8 mil thick films, respectively.1024 layered films containing 5 wt% APOSS only see a slight in-crease of less than 8%. Neat nylon 6 control films also see an in-crease in yield stress for biaxially oriented samples by 25 and 50%,for films stretched from 12 mil and 8 mil thick films, respectively.The neat nylon 6 films exhibit the lowest increases in yield stress foras-extruded films, but see the highest increase for biaxially ori-ented films.
It is evident from these results that neat nylon 6 control films,compared with APOSS containing films, undergo equal or greaterincreases in elastic modulus, especially in biaxially oriented sam-ples; individual values remain lower for neat nylon 6 in as-extrudedfilm samples. Neat nylon 6 as-extruded films only see a slight in-crease in yield strength, but biaxially oriented neat nylon 6 samplesexhibit higher values than those of APOSS containing films. Theseresults seem to indicate that the inclusion of APOSS within multi-layered nylon 6 films results in an overall decline of tensile prop-erties with respect to increased orientation, relative to neat nylon 6.The highest tensile properties, for both elastic modulus and yieldstress, were observed for the biaxially oriented nylon 6 controlcontaining no APOSS. Films containing 2.5 and 5.0 wt% APOSSunderwent an increase in elastic modulus with respect todecreasing layer thickness, but to a lesser extent than the nylon 6control. Films containing 2.5 wt% APOSS saw a greater increase inmodulus than films containing 5.0 wt% APOSS, indicating that
higher concentrations of APOSS results in greater losses in theelastic modulus. APOSS containing films saw a higher increase inyield stresses, than did the nylon 6 control for as-extruded films;the highest yield stress values came from biaxially oriented neatnylon 6 control films. These results are most likely resultant fromchanges in the crystal structure, as described in the followingparagraphs, with contribution from phase separated APOSS ag-gregates strengthening films via void formation and crazing, as thenylon 6 matrix elongates around the particles. In considering onlyAPOSS contributions, for thicker as-extruded films, greater con-centrations of APOSS produces higher yield stresses, relative to neatnylon 6. However, for highly elongated as-extruded and biaxiallyoriented APOSS samples, void formation and coalesce leads to adecrease in tensile properties and early fracture, relative to neatnylon 6. Without the presence of APOSS in the matrix, biaxiallyoriented neat nylon 6 undergoes a high degree of chain orientation,allowing a much greater increase in properties.
Similar to the iPP/DBS/TPOSS film samples, and displayed inFig. 11, APOSS prefers to form large, spherical phase separated ag-gregates, minimizing interactions with the polymer matrix. Upontensile deformation, void formation occurs due to cavitationlocalized at particleematrix interfaces. Deformation of the polymermatrix around debonded particles occurs parallel to the loadingdirection on both sides of the APOSS aggregate. In order to producethe thinner as-extruded films, faster uptake speeds are used, whichcan result in greater void formation and elongation during pro-cessing. Thus in highly elongated films, lateral coalescence of voidsaround POSS aggregates can lead to early fracture and weakenedtensile properties, relative to neat nylon 6 [34]. Thermal analysis ofthe APOSS/nylon 6 films further confirms that APOSS aggregatesare non-covalently interacting with the nylon 6 matrix. This wasdetermined by a lack of change in the melting behavior of nylon 6(Fig. 12) between controls and APOSS containing films.
The variation of G0, G00 and tan d as a function of frequency at 240and 245 �C, the same temperatures at which the blends and filmswere prepared, respectively, for nylon 6 and nylon 6/APOSS isshown in Fig.13. At 240 �C, a small influence of the APOSS content isshown mainly on the viscous behavior (G00). However, at low fre-quencies it is possible to observe the influence of POSS aggregatesin nylon 6/APOSS (5.0%) samples, with the increase of G0 and the
Fig. 11. SEM image of nylon 6/APOSS multilayered film displaying the formation oflarge POSS aggregates.
M.M. Herbert et al. / Polymer 54 (2013) 6992e70037000
formation of a plateau for tan d. With the increase of temperature,the presence of APOSS content is more evident. This is evident fromthe decrease of the loss modulus for nylon 6/APOSS (2.5%) samples,and an increase for nylon 6/APOSS (5.0%) samples. It seems that atlow concentrations (2.5%) there is a better dispersion of POSS, withsome level of nylon 6e APOSS interaction, most likelyweak van derWaals forces. This could lead a decrease in chain entanglement andmore free volume in the melt, resulting in lower viscous behavior.For higher concentrations (>2.5%), APOSS aggregates produce in-creases in G00. The increased storage modulus at 245 �C for 5%APOSS content supports this argument, with a more pronouncedplateau for tan d at low frequencies.
Results from Wide-Angle Xray Spectroscopy (WAXS) show astandard response of the nylon 6/APOSS films to increasing amount
140 190 240 290
1024 control
1024 2.5%
1024 5%
256 2.5%
256 5%
140 160 180 200 220 240 260
1024 control
1024 2.5%
1024 5%
256 2.5%
256 5%
Temperature (°C)
300 µm Thick Films
25 µm Thick Films
Temperature (°C)
A
B
Fig. 12. DSC curves showing minimal changes in the melting peak of nylon 6/APOSSfilms for 300 (A) and 25 mm (B) thick films.
of orientation and faster cooling rates during processing. Under aconstant extrusion rate, lower shear stresses and slower crystalli-zation rates are inherent in the fabrication of thicker films, whileextrusion of thinner films produces a larger degree of orientationand undergoes faster crystallization rates, as a result of the lowervolume and faster uptake speeds. Additionally, the presence ofimpurities such as POSS can act as nucleation sites, increasing thenumber of crystals. These differences lead to the formation of themore stable a-nylon 6 crystal structure in thicker films, and g-nylon6 in thinner, more oriented films [35]. As evident in Fig. 14, as-extruded 12 mil thick films, regardless of the presence of POSS,displayed (200) and mixed (002/202) diffraction peaks, at 14.2 and17.0 nm�1 (d-spacing ¼ 0.440 and 0.370 nm, respectively), whichare characteristic of the a-nylon 6 crystal structure. On the otherhand, as-extruded 1 mil films displayed (100) diffraction peak at14.9 nm�1 (d-spacing ¼ 0.413 nm), with a minor (200/201)diffraction peak, characteristic of the g-nylon 6 crystal structure. Inthese crystal structures, plasticity has been determined to be gov-erned by the crystal slip parallel to the hydrogen bonded sheets, i.e.parallel to the (002) crystalline plane in a-nylon 6 and to the (200)plane in g-nylon 6 [35]. a-Nylon 6 forms intra-sheet hydrogenbonds between fully extended chains in anti-parallel fashion, suchthat the amide andmethylene units lie in the same plane. g-Nylon 6forms inter-sheet hydrogen bonds between pleated chains in aparallel fashion, such that the linkages lie 60� out of plane [35],leading to a-nylon 6 having better stability intra-sheet but a largeramount of inter-sheet breakdown, compared with g-nylon 6. Themore stable a-nylon 6 therefore exhibits a relatively highermodulus and yield stress than g-nylon 6.
Biaxial orientation of the 12 and 8mil films imposes even higherdegrees of orientation and chain extension to the films than duringextrusion, resulting in the development of a mesomorphic b-nylon6 crystal structure. The b-nylon 6 crystal form is described as amesomorphic pseudo-hexagonal structure with a random distri-bution of hydrogen bonds about the chain axis, and is a phase in-termediate between the a-nylon 6 and g-nylon 6 crystalline forms[36]. This evolution of the nylon 6 crystal structure from a to g andfinally b, has been previously shown to occur consistently withincreasing strain on the system. While an increase of strain on thea-nylon 6 crystal structure is responsible for the a to g transition,further strain above 150e180% will cause the crystal structure tobegin to revert back to its initial a structure [37]. As shown inFig. 14, the b-nylon 6 crystal structure is characterized by peaksrepresenting both the a-nylon 6 and g-nylon 6 crystal structures.The primary (200) diffraction peak, at 14.2 nm�1 represents the a-nylon 6 intra-sheet spacing, while the smaller (002/202) a peak and(100) g peak at 17.0 and 14.9 nm�1, represent the a-nylon 6 inter-sheet spacing and the g-nylon 6 intra-sheet spacing, respectively.
The degree of crystallinity in each of the film types, shown inFig. 15, was determined using the heat of enthalpy obtained fromDSC melting peaks. For both neat and APOSS containing as-extruded nylon 6 samples, the degree of crystallinity remainsrelatively constant, fluctuating around 20%. However, for biaxiallyoriented samples, an increase in the amount of crystallinity byaround 5e10% is shown, most likely the result of increased orien-tation in the amorphous nylon 6 chains contributing to the meso-morphic b-nylon 6 crystal structure.
As described previously, this system contains a lack of covalentinteractions between APOSS and nylon 6 molecules, although, aswas the case with the structure-property relationships of APOSSwithin melt-spun nylon 6 fibers, it was assumed that there shouldbe a significant amount of hydrogen bonding occurring [22]. Ramanspectroscopy was used in order to probe and analyze the type ofmolecular interactions taking place between APOSS molecules andnylon 6. The results displayed in Fig. 16 indeed show that a
Fig. 13. Rheological data for nylon 6 and nylon 6/APOSS blends at 240 and 245 �C.
M.M. Herbert et al. / Polymer 54 (2013) 6992e7003 7001
significant shift in hydrogen bonding is taking place between thethicker 12 mil and thinner 1 mil APOSS containing films. Incomparing the Raman spectra of unprocessed neat nylon 6 (a) with2.5 wt% APOSS containing film of 12 mil thickness (b), only veryslight changes in peak heights are observed. Overall, these spectra
represent that of standard a-nylon 6. The 2.5 wt% APOSS film of 1mil thickness (c) however, shows peak shifts indicative of thecrystalline transition from a-nylon 6 to g-nylon 6. The primarycharacteristic of this transition is the rotation of the hydrogenbonded amide groups by 60�, shifting from intra-sheet dominated
Fig. 14. WAXS spectra which correlates to changes in the crystallographic structurefrom a-nylon 6 (A), to g-nylon 6 (B) and finally b-nylon 6 (C). Specific peaks corre-sponding to a and g crystal structures are labeled.
0
5
10
15
20
25
30
35
0 200 400 600 800 1000 1200 1400
DSC: % Crystallinity from Specific Heat of Fusion
Layer Thickness
%Crystallinity
Fig. 15. DSC determination of crystallinity for 256 layered film samples containing2.5 wt% APOSS, with respect to decreasing layer thickness (nm). Crystallinity wascalculated from first heating curves using the inset equation shown in Fig. 10.
M.M. Herbert et al. / Polymer 54 (2013) 6992e70037002
hydrogen bonding to inter-sheet hydrogen bonding. This rotationof the amide group relative to the incident radiation is believed tobe responsible for the peak disappearance at 3300 cm�1. Theappearance of the 3150 cm�1 peak has been previously reported asa combination of the amide I (CO stretching) and amide II (NH in-plane bending) vibrations, located at 1610 cm�1 and 1460 cm�1,respectively [38], although this peak has also been reported as anovertone of the amide II band [39]. Reduction of the 2900 cm�1
peak, representing CH2 stretching, and broadening of the amide II
Fig. 16. Raman spectroscopy of 256 layered films illustrating the differences in
peaks around 1460 cm�1, has been demonstrated to be a result ofthe crystallographic change from a-nylon 6 to g-nylon 6 [37,40].Finally, a shift of the amide I peak, from 1634 cm�1 to 1610 cm�1 isseen, which suggests a change from nylon 6 intra-sheet hydrogenbonding in a-nylon 6, to inter-sheet hydrogen bonding, betweennylon 6 sheets in g-nylon 6 [41]. It is also possible that this peakrepresents primary amide NH2 bending, which would indicate thatPOSS does in fact interact with the g-nylon 6 crystal form [39].However, as explained previously, the influence that POSS has onthe blend is negligible. Additionally, the shifting of this peak tolower frequencies is indicative of increasing stress within thecomposite.
4. Conclusions
From the results presented on the characterization andstructure-property relationships of both the iPP/DBS/TPOSS andnylon 6/APOSS blends, it is evident that POSS is a poor filler, lackingsolubility and favorable interaction with the polymer matrices. Itwas initially hypothesized that under extreme confinement andorientation, such as in melt-spun fibers, or confined within 2Dnanoscale layers, that POSS would undergo forced-assembly intoelongated, rebar-like reinforcement structures, or even act ascrosslinking molecules for the polymer chains. The current results,however, show POSS existing as large, phase separated aggregates,in order to minimize interactions with the polymer matrix. Addi-tionally, the aggregates behave as debonded hard particles upontensile deformation. POSS has been previously shown to enhancethe properties of polymer matrices in which the POSS moleculeshave been grafted to, or copolymerized within the chain, but this isnot the case for POSS blends. In comparison to results from the iPP/DBS/TPOSS system, inwhich POSS is unable to directly interact withthe polymer matrix, and the nylon 6/APOSS system, in which POSScan potentially form hydrogen bonds with the polymer matrix, theresults are similar and reveal that POSS blends are largely incom-patible with the polymer matrix. Small improvements can be madevia functionalization of the POSS cage, in order to enhance in-teractions, but these improvements are relatively negligible.
Looking into the structureeproperty relationships of thesesystems, mechanical properties exhibit increasing values withrespect to decreasing film and layer thickness. However, WAXS andPOM revealed that these enhancements can be attributed tochanges in the crystal size and structure, via confinement of theindividual layers, rather than interactions between POSS and
hydrogen bonding with respect to film thickness, relative to neat nylon 6.
M.M. Herbert et al. / Polymer 54 (2013) 6992e7003 7003
polymer; reduction of spherulite size in iPP/DBS/TPOSS systems,and strain induced change in the crystal structure of nylon 6 withincreasing chain orientation in nylon 6/APOSS systems. Additionalinfluences on the mechanical properties can arise as a result ofcontributions from POSS aggregates upon sample deformation,responsible for voiding, shear banding and crazing, depending onthe sample thickness and degree of orientation. Further confirma-tion of non-interaction between POSS and polymer is seen via a lackof change in thermal DSC data, however this simply rules out anyunexpected covalent interactions.
Rheological investigations on the POSS-polymer systemsrevealed that under elevated temperatures, POSS solubility anddispersion is slightly enhanced. This enhancement is a result ofbreakup of larger POSS aggregates, allowing for better diffusionthrough the matrix; improvements are seen only for lower (2.5%)POSS loadings. In agreement with work by Jana [26], POSS ag-gregates are shown to act as a processing aid, reducing the vis-cosity of the blends via decreasing chain entanglement andincreasing free volume, allowing greater orientation of thepolymer.
Raman spectroscopy was used to investigate any contributionsto the structure-property relationships as a result of hydrogenbonding between nylon 6 and APOSS; the resultant spectra, whileshowing significant shifts in nylon 6 hydrogen bonding peaks,indicated a lack of POSS-polymer hydrogen bonding interactions.Instead, the results correlate with WAXS results, showing thatunder increased confinement and orientation, nylon 6 undergoeschanges in the crystallographic structure, evolving from the a- to g-and finally b-nylon 6 forms. From these results, it can be concludedthat improved mechanical properties are due to changes in thenylon 6 crystal structure with respect to increasing strain andorientation, and not the result of interactions between POSS andpolymer.
Overall, POSS has been shown to be a poor choice of filler inimproving the properties of POSS-polymer blends. Without directchemical linkage to the polymer matrix, the low solubility of POSSresults in the formation of large, phase separated aggregates. Thepresence of interacting groups on POSS molecules, such as aminegroups on APOSS and nylon 6, have not led to demonstratedimprovement over the non-interacting iPP/DBS/TPOSS system.Even under extreme confinement conditions, where significantdeviations from bulk behavior occurs, such as within melt-spunfibers or nanoscale layers, POSS was unable to be forcibly-assemble into reinforcing structures with the polymer matrix.
Acknowledgment
Financial support from the NSF Center for Layered PolymericSystems, grant number DMR 0423914, is gratefully acknowledged.POSS� is a registered trademark of Hybrid Plastics, Inc.
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