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Reactive Blending of Functionalized Acrylic Rubbers and Epoxy Resins.

C. DISPENZA [1][*]

J. T. CARTER [2]



A high molecular weight acrylonitrile/butadiene/methacrylic acid (Nipol 1472) rubber is chosen to control processability and mechanical properties of a TGDDM (tetra glycidyl diphenyl methane) based epoxy resin formulation for aerospace composite applications. The physical blend of rubber and epoxy resin, achieved by dissolution of all the components in a common solvent, forms a heterogeneous system after solvent removal and presents coarse phase separation during cure that impairs any practical relevance of this material. A marked improvement of rubber-epoxy miscibility is achieved by reactive blending ('pre-reaction') the epoxy oligomer with the functional groups present in the rubber. The epoxy-rubber adduct' so obtained appears as a homogeneous system at room temperature and also after compounding with the curing agent. Depending on the nature and extent of interactions developed between the rubber and the epoxy resin during 'pre-reaction,' materials with different resin flow characteristics, distinctive m orphologles and mechanical properties after curing were obtained. The effect of 'pre-reaction' on the resin cure reaction kinetics has been also investigated.


Epoxy resins are commonly used as polymer matrices in high performance composites for their chemical, mechanical and thermal resistance characteristics (1, 2). Both unidirectional tapes and woven fabric prepregs are produced using epoxy resins, and they find many applications in aircraft construction, including use as the skins on honeycomb core structures. Such structures, which use woven fabric prepregs, are particularly difficult to manufacture to meet the very high design criteria of the aircraft industry since there are some conflicting processes and property requirements. For example, the skin on the final structure must be void-free to eliminate the possibility of water ingress to the honeycomb structure while the aircraft Is in service. Any air pockets and entrapped solvents In the core or in the prepreg tend to expand/volatilize during heating, thus causing voids to form. Unfortunately, this cannot be prevented by the application of high pressure to the Nomex core, as it can with solid panels, since the core would collapse. It is, therefore, necessary to control the rheological properties of the matrix resin both throughout the prepregging process, to eliminate air pockets from the fabric at this stage, and during the honeycomb structure manufacturing process, to allow all volatiles to escape (3). The resin flow characteristics dictate the quality of the prepreg and of the final laminate. In fact, poor resin flow can lead to uneven distribution of the resin in the laminate and difficult removal of entrapped solvents and air, which would, in turn, cause porosity and structural defects. Conversely. excessive bleeding can cause poor performance of the honeycomb core chamfer. A controlled viscous flow, as the temperature is increased in order to convert the relatively low molecular weight polymer into the crosslinked network, is required for a successful utilization of the composite part.

Although lower cure temperatures generally mean also lower glass transition temperature for the cured resins and poor temperature resistance, tetrafunctional epoxy resins cured with amine hardeners have potential high crosslinking density and moderately high glass transition temperature (about 180[degrees]C) at reasonably low temperature of cure. Epoxy matrix formulations based on tetraglycidyl diamino diphenyl methane (TGDDM) and diamino diphenyl sulfone (DDS) are, therefore, baseline prepregs materials. On the other hand, the experience of the past twenty years reveals a number of shortcomings that TGDDM resins have in common with other epoxy resins: viscosity, which dramatically lowers when temperature is increased during the cure cycle, poor damage resistance due to their high crosslinking density.

The resin flow can be controlled and the mechanical performance can be improved by incorporation of elastomers or plastomers [4-14]. The incorporation of such materials in epoxy systems can be done in a variety of ways, but essentially corresponding to two different approaches: 'blending' and pre-reacting' techniques. The 'blending' technique consists in a physical blend of the rubber with the epoxy resin precursors: all the different components are charged either at the same time or at different stages and eventually lightly warmed up in order to speed up the process of mutual dissolution and/or lower the overall viscosity; the resin is then cured in a single complex hardening process, which can involve also the reactive functionalities of the rubber if present. The rubber is generally soluble in the monomers mixture only at higher temperatures than ambient and precipitates as a separate phase from the resin bulk during the cure, owing to the reduced solubility of the elastomer in the cross-linked resin. Thi s methodology Is effective for many purposes but, sometime, improved miscibility and better control of the final morphology are obtained through a 'pre-reacting' technique [12, 15]. In this case, suitable modifiers, with reactive species capable of forming covalent bonds with the resin functional moieties, are reacted with one of the resin precursors, thus ensuring a stable chemical structure to be formed prior to the cure. The rubber/resin 'adduct' that is so formed can be diluted with more of the same monomers used to form it and/or added with the curing agent to produce the final modified resin formulation at the desired resin/rubber ratio.

This study investigates the possibility of controlling the resin flow characteristics and improve the mechanical performance of a TGDDM/DDS epoxy resin, elsewhere characterized in its cure behavior [16, 17], by the incorporation of high molecular weight rubber that is a random copolymer of butadiene, acrylonitrile and methacrylic acid. This commercial copolymer is solid rubber with a fairly low Tg (-40[degrees]C), fast dissolving in several organic solvents and stable at the temperatures required by the TGDDM/DDS resin cure process.

The influence of the rubber-epoxy blending procedure on the resin cure kinetics, its flow characteristics, morphology and mechanical properties upon curing has been investigated.



The epoxy components were a nominally tetrafunctional epoxy resin, MY9663 supplied by Ciba Geigy, and a nominally difunctional epoxy resin, Epon 828, supplied by Shell Chemicals. The actual composition of the MY9663 resin was determined by a "High Pressure Liquid Chromatography" (HPLC) apparatus, equipped with a Spherisorb S5 ODS2 column and an UV detector, using water and acetonitrile gradients as mobile phase at 45[degrees]C. The chromatogram showed the presence of tetraglycidyl diamino diphenyl methane (TGGDM) monomer, dimer and other high oligomers. The highest peak refers to the monomer (68.4%), the second higher (6.27%) represents the dimer content and the complement to 100 takes into account all the higher oligomeric species. The batch used had an epoxide equivalent weight of 125.5. Epon 828 is the diglycidyl ether of bisphenol A (DGEBA) with an epoxide equivalent weight of 344. The curing agent used was 4,4'-diaminodiphenyl sulfone (DDS) supplied by Ciba Geigy as HT 976.

The rubber was a butadiene-acrylonitrile-methacrylic acid copolymer. 28/68.8/3.2 % w, supplied by Nippon Zenon with the tradename of Nipol 1472.

Three different rubber-epoxy resin systems, which correspond to three different blending techniques, have been investigated:

(i) Epoxy resin pre-reacted with Nipol 1472 in the presence of ETPI catalyst (PWE system): Five parts per hundred resin (5 pph) of Nipol 1472 were previously dissolved in acetone and then charged into a flask containing the TGDDM monomer (63.75 pph) and the catalyst (0.5 pph). The reactor was fitted with stirrer, nitrogen inlet and condenser. The contents were heated to gently reflux for 2 hr and then allowed to cool to room temperature when the other epoxy oligomers and the curing agent were also charged and stirred until a homogeneous mixture is obtained.

(ii) Epoxy resin pre-reacted with Nipol 1472 with no ETPI catalyst present (PNE system): Exactly the same procedure was followed as in (i) except that ETPI has not been used.

(iii) Epoxy resin/Nipol 1472 blend (Blended system): The rubber pre-dissolved in acetone and all the epoxy resin precursors at the same composition as in (i) and (ii) were charged into a flask and stirred at room temperature until a clear mixture was obtained.

After either of the above stages had been completed, 19.25 pph DDS curing agent and 12 pph difunctional epoxy resin were added.

Two other resin systems. which do not contain rubber, were formulated to be used as reference:

(iv) Base epoxy system with a TGDDM/DDS as in the blends, (SCN system);

(v) Base epoxy system with 0.5%w/w of ETPI, (SC system).

Sample Preparation

* Samples for kinetic and rheological evaluation. Kinetic and rheological studies were performed on initially uncured samples. Acetone was removed by heating the sample to 60[degrees]C and maintained at this temperature 30 minutes under partial vacuum and 30 minutes under full vacuum. PWE and PNE appeared as homogeneous mixtures after solvent removal, while the 'Blended' system had a turbid 'milky' appearance. Either fresh samples were used when possible or they were stored at -40[degrees]C in a desiccator until the analysis was performed.

* Neat Resin Panels. In order to obtain neat resin panels, appropriate amounts of resin, already stripped of the acetone and moisture present and stored at -40[degrees]C for few hours to freeze, were shaped by means of a picture frame within the plates of a press. The panels in their picture frame were bagged up with a connection to the vacuum line and placed within an autoclave, which was, in turn, placed between the plates of a heated press. Vacuum was applied at the start and then released after the autoclave was charged with 5 atm of [N.sub.2] gas. The following heating cycle was used: heating to 180[degrees]C at 2[degrees]C/min, 2 hour dwell at 180[degrees]C, cooling to room temperature at 2[degrees]C/min. The temperature was measured by means of a thermocouple allocated in the autoclave body close to one of the sample surfaces. After the cure cycle, the solid resin panels were demolded and cut to the appropriate dimensions for the dielectrical, morphological and mechanical characterizations.


* Differential Scanning Calorimetly (DSC). Isothermal cure was first performed placing 5-10 mg of resin within sealed aluminum pans in a Perkin-Elmer DSC-7 apparatus for different lengths of time. After dwelling at the temperature of 180[degrees]C, samples were quenched in liquid nitrogen and then their residual heat of reaction was measured in a dynamic run carried out from 30[degrees]C to 300[degrees]C at 10[degrees]C/min.

* Potentiometric titration. Samples of resin formulations as in (i) and (ii) were withdrawn at different times from the reactor during the pre-reaction, devolatilized in a vacuum oven at 60[degrees]C, dissolved in toluene/isopropanol solution and finally titrated using potassium hydroxide. The change in electrode potential has been followed as a function of the volume of titrating solution. The equivalence point of the reaction has been revealed by a sudden change in the potential values.

* Chemorheology. A Rheometrics System RDA-II fitted with a parallel-plate test cell was used for the rheological measurements. The test sample was confined in the gap between two plates having 1.25 cm radius. The typical gap between the plates was 1.5 mm. The test chamber was always preheated to the test temperature before loading each sample to ensure precise gap settings. Tests were performed in 'time sweep' mode at angular frequency of 10 rad/sec, strain amplitude of 10% and three different temperatures: 140[degrees]C, 160[degrees]C and 180[degrees]C. Complex viscosity ([[eta].sup.*]), storage modulus (G') and loss modulus (G") values as a function of the time were recorded.

* Resin flow evaluation. The same apparatus and sample geometry as above were used. Test samples were preconditioned at room temperature by pressing between two poly(tetrafluoroethylene) cloths to produce flat circular shaped sheets with given uniform thickness. All samples were stored at -30[degrees]C before being loaded in the rheometer. Flow tests were carried out in 'temperature sweep mode' at 10% strain, 10 [rads.sup.-1] angular frequency, in the temperature range 60[degrees]C-180[degrees]C with a heating rate of 2[degrees]C/min.

* Dielectric Spectroscopy. Dielectric tests were carried out under vacuum with an a.c. HP precision LCR meter model 4284A, with a three electrode cell on circular samples of 5 cm diameter [18]. Loss tangent, tan [delta], was measured at the fixed frequency of 1 kHz in the temperature range from -80[degrees]C to 80[degrees]C. The heating rate was about 1.5[degrees]C/min. From the same experiment permittivity data were withdrawn but they have been considered unreliable because of the uneven sample thickness.

* Mechanical Testing. The test pieces were machined from plaques 80 mm square with thicknesses between 4 and 5 mm and stored in air at 23[degrees]C, 50% relative humidity before testing at the same conditions. The toughness of the materials was characterized using two parameters: yield stress in compression and toughness using a linear elastic fracture mechanics analysis. The yield stress in compression was determined using a deformation rate of 1 mm [min.sup.-1].

Fracture strength, [K.sub.1c], and fracture toughness, [G.sub.1c], were measured using a deformation rate of 1 mm [min.sup.-1]. The ductility factor was calculated as a ratio between [K.sub.c] and the yield stress in compression. For all samples, the plastic zone size at the crack tip was of the same order as the crack tip radius (i.e. between 1 and 10 [micro]m). The precision of the LEFM analysis becomes worse as the plastic zone decreases from 10 [micro]m to 1 [micro]m crack tip radius; this has the effect of decreasing the computed values for [K.sub.c], and thus reducing the ductility factor. Since these materials exhibit a range of plastic zone sizes from 4 to 8 [micro]m, it is possible to presume that the true toughness is greater than the measured.

* Morphological characterization. The morphological study was carried out on the fracture surfaces from neat resin panels obtained breaking samples into liquid nitrogen, using a Philips 501 model Scanning Electron Microscope (SEM). In order to have a better resolution of the morphology, surfaces were etched with refluxing THF until all the rubber phase was extracted. All samples were mounted on stubs for SEM analysis and gold coated.


The first attempt of incorporating the solid Nipol 1472 elastomer into the TGDDM-DDS resin system followed the approach of physically blending all the components. In order to promote homogenization, they have been pre-dissolved in a common solvent and stirred until a clear mixture was obtained. Nonetheless, after solvent removal, the obtained blend presented a turbid or milky appearance, thus suggesting that the system is heterogeneous at room temperature. Heating the blend to cure temperature (180[degrees]C) did not transform it into a clear mixture. Furthermore, the buildup of molecular weight in the resin matrix during polymerization resulted in a coarse phase separation, with two phases clearly detectable in the cured panels under visual inspection.

In order to improve compatibilization between the phases, reactive blending of the rubber and the TGDDM epoxy monomer was promoted in a stage prior to cure. In particular, the reaction was carried out at 60[degrees]C using acetone as solvent, in the presence of a catalyst (ETPI) for the PWE system and without catalyst for the PNE formulation. In both cases, homogeneous mixtures were obtained after pre-reaction, which retained their clear and transparent appearance also after solvent removal and compounding with DDS and the difunctional epoxy monomer, this last component having been added to the base epoxy system to reduce the network density.

The pre-reaction process of TGDDM and Nipol 1472 converts some the carboxyl groups of the rubber into ester linkages with the epoxy monomer (19). Residual carboxylic group concentration vs. reaction time was measured via potentiometric titration analysis; results are shown in Fig. 1. The presence of catalytic amounts of ETPI considerably increases the rate of this reaction, leading to almost 100% conversion of carboxyl groups in less than two hours. The reaction also occurs when ETPI Is not present, although it proceeds at a lower rate and goes slowly to completion. In particular, after two hours of pre-reaction, the PNE system shows approximately 30% of carboxyl groups remaining and after five hours a residual 10%.

The influence of both rubber and ETPI presence on the overall cure reaction rate was investigated, together with the effect of pre-reaction. The extents of the overall cure reactions vs. reaction time curves were determined for PWE, PNE and Blended systems isothermally cured at 180[degrees]C by DSC experiments. The advancement of cure as a function of time was calculated through measurements of heat of reaction with the assumption that the heat evolved is proportional to the extent of cure, x.

x = [delta]Hrxn - [delta]Hres/[delta]Hrxn (1)

where [delta]Hrxn is the total heat liberated when the uncured mixture is taken to complete cure and [delta]Hres is the residual heat of reaction of partially cured samples.

Results are reported in Fig. 2. Data relative to the pure epoxy system with and without catalyst (SC and SCN respectively) are also reported for comparison. Conversion curves for the formulations containing ETPI (PWE and SC) show approximately the same behavior, while the Blended system, in which the rubber is simply physically mixed with the resin, has the same reaction rate of the pure epoxy system with no catalyst. Interestingly, the PNE system exhibits a conversion curve that lies In an intermediate position between those of systems with ETPI and those of systems without ETPI. The observed results may imply that ETPI, used as catalyst for the pre-reaction, also promotes the resin cure process, covering other effects that may come from changes in the nature and concentration of the functional groups as a result of the pre-reaction. The Blended system behaves as the pure epoxy resin, because it is likely that no change in nature and concentration of reactive groups occurred during compounding. The intermed iate positioning of PNE curve can be explained by a certain degree of epoxy ring opening after two hours of pre-reaction, also confirmed by potentiometric titration results reported in Fig. 1 and with the related increase of the concentration of the hydroxyl groups, which can have a catalytic effect.

The different cure behavior of PWE, PNE and Blended systems Is also reflected in the values of gel time measured through chemorheological experiments carried out at constant temperature.

It is well known [20, 21] that during the cure, thermosets change from viscous liquids to rigid glassy solids. Concomitant with the changes in structure, there can be a dramatic change in their rheological properties. By performing dynamic-mechanical rheological experiments at constant temperature, it is possible to relate the viscoelastic property-time profiles to the reaction kinetics of the resin system under cure. In particular, at the molecular level, gel is formed when at least one of the molecules has grown very large and its size reaches dimensions of the order of the macroscopic sample. In a chemorheological experiment, the gel point can be measured as the point in which a distinct change in slope of G'(t) curve is observed.

Tests on PWE, PNE and Blended systems were performed at three different temperatures (140[degrees], 160[degrees] and 180[degrees]C) Always, both storage modulus, G', and loss modulus, G", were observed to smoothly increase during the first stage of the cure until they exhibited an abrupt change in slope, located at different times for samples cured at different temperatures. Because the cure reactions are thermally activated, the effect of higher curing temperatures is an enhancement of the reaction rate and a reduction in the time required for the onset of the crosslinking reactions with consequent drastic increase in viscosity. In Figs. 3-5, G' and G" curves for the three systems cured at the temperature of 140[degrees]C are shown. It can be observed that G'(t) and G"(t) curves for PNE and Blended systems cross at certain curing time, while no crossover point is observed for the PWE system (Fig. 3). Similar behavior was observed at 160[degrees]C, while no crossover point was detected for the three systems when the expe riments were run at 180[degrees]C. Therefore, it seems that the pre-reaction in the presence of ETPI promotes an early development of resin elasticity, which soon dominates the viscous flow. The different behaviors among the systems are less pronounced at 180[degrees]C because the high temperature speeds up the curing reactions, in this case, the rapid formation of a crosslinked epoxy-amino network being the major cause of enhancement in elasticity.

Gel time values for the three systems, measured as the onset of the abrupt change in slope of the G'(t) curves under the various isothermal curing conditions, are reported in Table 1. Gelation for PWE occurs always earlier at all the temperatures considered. According to the above-discussed thermo-kinetic analysis and in good agreement with the more detailed study carried out for SC and SCN systems elsewhere reported [16, 17], ETPi seems to behave like a catalyst for the primary amine-epoxy reaction, this reaction dominating the cure until vitrification occurs.

Resin flow evaluation was undertaken on the three systems under investigation. For each system, Complex Viscosity ([[eta].sup.*]), Dynamic Storage Modulus (G') and Dynamic Lass Modulus (G") as a function of temperature were measured and reported in Figs. 6-8. It can be observed that when the Nipol 1472 Is not fully pre-reacted with the epoxy monomer (PNE and Blended system), it is always G" [greater than] G', and both the elastic and viscous components of the resin viscosity decrease drastically with an increase of temperature, while for PWE, G' [greater than] G" and the resin flow characteristics are maintained during the whole range of temperatures considered. This behavior could be consistent with the presence of a 3D-network after the pre-reaction stage owing to the formation of epoxy bridges between different molecules of rubber, which promote the resin elasticity. Conversely, a sudden rise of G' is clearly detectable in the curves of both the PNE and Blended systems and may be related with the occurrence of phase-separation and coalescence phenomena within the resin bulk [22]. Not surprisingly, this phenomenon is more pronounced and occurs at lower temperature for Blended than for the PNE system.

In order to obtain more detailed information on the structure of the adduct formed between the rubber and the epoxy resin and, consequently, on the nature of interaction between the two phases of these blends, dielectric spectroscopy analysis was performed. It is well known that the dielectric polarization of an heterogeneous material also includes interfacial polarization phenomena, which, in turn, are related to the extent of interactions between the polar groups of each component at the interface (18). Beside the influence of the molecular structure on the dielectric behavior, i.e. the polarity of its chain segments, its electric conductivity, and the effect of morphology, such as the degree of dispersion and the extent of the interface, must be taken into account. In this study, the loss tangent, tan [delta], as a function of temperature was measured in the temperature range of -80[degrees]C + 150[degrees]C, in order to evaluate the dissipation process related to the main transition of the rubbery phase. The analysis was undertaken only on PWE and PNE neat resin systems, as the observed occurrence of macroscopic phase separation phenomena in the cured Blended system prevented the attainment of representative samples. In Fig. 9, plots of tan [delta] as a function of temperature at the frequency of 1 kHz are reported. The higher dumping amplitude and the lower temperature in correspondence to the maximum of the curve observed for the PNE system are in accordance with the expected lower degree of interlock between the two phases in this case and the possible formation of separated rubbery phase domains.

Morphological analysis was carried out by SEM on the fracture surfaces of the PWE and PNE neat resin samples. As it was not possible to directly resolve their morphologies, the analysis was performed on etched samples, using THF as solvent for the rubber. In Fig. 10, micrographs of the fracture surface of PWE and PNE samples are reported. For PWE (Fig. 10a), an essentially homogeneous morphology Is observed, with only few and small voids present, possibly caused by traces of moisture and residual volatiles that could not escape during cure because of the high viscosity of the resin. This morphology is coherent with the high degree of intermolecular links between rubber macromolecules promoted by the pre-reaction in the presence of ETPI. Conversely, small and uniformly distributed cavities (about 1 micron diameter) are left by extraction of the rubber phase in the PNE sample (Fig. 10b), as a result of the fact that the rubber is not fully pre-reacted with the epoxies. For a measurement of the overall area of the cavities for surface unit, the rubber does not seem to be all phase separated.

Table 2 reports yield stress in compression, fracture strength, [K.sub.1C], fracture toughness, [G.sub.1C], and ductility factor, DF, values measured from a three-point bend test in an opening mode on single-edge notched samples and a calculated modulus according to the following equation [23, 24]:

[E.sub.CALCULATED] = [(1 - v).sup.2]([[K.sup.2].sub.C]/[G.sub.C]). (2)

Comparing the results of the two materials, we see that the system pre-reacted with ETPI has higher compressive yield stress, indicating greater stiffness and higher [K.sub.c], both these two parameters leading to greater values of toughness, as shown by the ductility factor.


The incorporation of a solid, high Mw, acrylonitrile/butadiene/methacrylic acid rubber in an epoxy resin system was performed following two different approaches: 'blending' the rubber with all the component of the resin system (Blended system) or pre-reacting it with one of the epoxy monomers, either in the presence of a proper catalyst (PWE system) or not (PNE system), adding only afterward the curing agent and any other component of the resin formulation.

The chosen incorporation technique induces changes in the chemical composition of the systems that affect their cure behavior, as indicated by calorimetric and chemorheological studies. In fact, the catalyst used to speed up the pre-reaction between the carboxylic functionalities of the rubber, and the epoxy moieties of the resin also catalyze the cure reactions (PWE system), as it does the formation of secondary hydroxyl groups from the epoxy ring opening reaction. Furthermore, the complete reaction of the rubber carboxyl groups with epoxy moieties leads to the formation of a three-dimensional epoxy-rubber network, swollen by the un-reacted monomers, which enables control of the resin viscous flow in a wide range of temperatures and prevent coarse phase separation phenomena upon cure, as observed for the Blended system. The superior performance of a highly interpenetrated rubber-epoxy network (PWE) is also reflected in a better mechanical testing response, if compared with the PNE system, which presents a l ower degree of chemical interlock between the two phases and a certain extent of rubber phase separation.

(*.)Corresponding author. Email:

(1.) Dipartimento di Ingegneria Chimica dei Processi e die Materiali University of Palermo, Vaile delle Scienze, 90128 Palermo, Italy

(2.) ICI Technology Wilton Centre, PO Box 90, Middlesbrough, TS90 8JE, UK


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Table 1

Gel Times Measured From Onset of
the G'(t) Slope Abrupt Change.

Temperature PWE PNE Blended

 180 21 27 24
 160 51 54 53
 l40 135 l40 137
Table 2

Mechanical Response of Neat Resin

Net resin system PWE PNE

Yield stress (MPa) 158 147
Kc (MPa [m.sup.-3/2]) 1.398 1.004
Gc (KJ [m.sup.-2]) 0.614 0.375
Calc Mod. (GPa) 3.16 2.7
DF (mm) 0.13 0.07
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Publication:Polymer Engineering and Science
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Geographic Code:1USA
Date:Sep 1, 2001
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