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Glass transition dependence of ultrahigh strain rate response in amine cured epoxy resins
Polymer 53 (2012) 5917e5923
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Glass transition dependence of ultrahigh strain rate response in amine cured epoxy resins Daniel B. Knorr Jr., Jian H. Yu, Adam D. Richardson, Mark D. Hindenlang, Ian M. McAninch, John J. La Scala, Joseph L. Lenhart* U.S. Army Research Laboratory, Aberdeen Proving Ground, MD 21009, United States
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 July 2012 Received in revised form 25 September 2012 Accepted 29 September 2012 Available online 17 October 2012
Ultrahigh strain rate performance in a series of model amine cured epoxy resins was investigated as a function of the glass transition temperature (Tg) of the cured polymer network, where the network Tg was systematically varied through the monomer stiffness, structure, and size. The high rate response was characterized in terms of a projectile penetration velocity, V50BL(P) (ballistic limit, protection criteria), which describes the projectile velocity with a 50% probability of sample penetration. One factor that dictates the ballistic performance of the epoxy networks, at effective rates of 104e105 s1, is the difference between the measurement temperature and the glass transition temperature of the network. Sub-Tg relaxations did not have a measurable effect on ballistic performance, and neither did the monomer structure and functionality outside of the influence of the resin Tg, while off-stoichiometric (excess amine) formulations improved V50BL(P) slightly with high Tg epoxies. The results have implications in protective materials for military, aerospace, transportation, and construction industries, where high strain rate insults from airborne debris, high rate collisions, and natural events are increasingly considered during product design. Published by Elsevier Ltd.
Keywords: High strain rate Epoxy-amine Glass transition
1. Introduction Epoxy resins are exploited heavily in both commercial and military technologies including fiber reinforced composites, electronics packaging, adhesives, transportation, construction, and infrastructure applications. A combination of processing versatility, chemical diversity, chemical resistance, and environmental stability makes epoxies useful for these broad applications [1]. However, a critical challenge is that these polymers are notoriously brittle [2]. In addition, high rate mechanical performance is increasingly important for light weight protective equipment in military, aerospace, transportation, and construction industries [3,4], yet epoxies have not been designed to function in these extreme mechanical environments. Substantial research has focused on the quasi-static mechanical properties of epoxy resins and the ballistic performance of many industrial formulations has been evaluated. However, these industrial formulations are complex and proprietary concoctions of many additives including: * Corresponding author. U.S. Army Research Laboratory, 4600 Deer Creek Loop, Aberdeen Proving Ground, MD 21005, United States. Tel.: þ1 410 306 1940; fax: þ1 410 306 0676. E-mail address: [email protected] (J.L. Lenhart). 0032-3861/$ e see front matter Published by Elsevier Ltd. http://dx.doi.org/10.1016/j.polymer.2012.09.058
monomer mixtures, chain extenders, viscosity modifiers, toughening agents, degassing agents, and even low volatility plasticizers, making fundamental insight difficult to extract. Timeetemperature superposition is potentially useful for predicting high strain rate behavior, but has only been demonstrated in select cases in the high strain rate regime [5,6]. Therefore, a critical knowledge gap exists in terms of the chemical, physical, and structural factors that govern high rate response in polymer networks. Reinforced epoxy resins are indispensable components in many protective technologies. The complex energy dissipative mechanisms in these composites are key, and take a variety of forms involving tensile failure, elastic deformation and kinetic energy transferred to the moving portion of the composite [7,8]. Zee and Hsieh found that, within a fiber reinforced polymerematrix composite itself, four processes are responsible for energy dissipation: (i) heat, (ii) fiber breakage/deformation, (iii) interfacial delamination and (iv) matrix cracking/deformation [9]. While the first process is common to most dissipative mechanisms, the second two are unique to composites, and have been studied extensively [10e18]. Mechanical response of the polymer matrix alone can be responsible for 20e35% of the total energy dissipation [9]. Despite this potential contribution, few studies have focused on improving high rate dissipative capabilities or on obtaining
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a fundamental understanding of matrix-only performance in ballistic environments [19]. High rate characterization of metals and ceramics is common with split-Hopkinson pressure bar (SHPB) techniques [20,21]. However, measurements of polymers are particularly difficult due to their large strain deformation capability, where inertial effects and complex stress states can convolute the mechanical response [22,23]. For example, deviations from linearity of compressive yield stress as a function of log strain rate are common with SHPB testing of polymers. Typically, yield stress increases from linear behavior at higher strain rates [24,25]. While some groups have attributed this to molecular relaxations in polymers at high strain rates [26,27], friction at the polymerebar interface can equally explain these results. Our own testing has demonstrated that lubricants are ineffective at preventing specimen “barreling” during high rate compression testing. In addition, depending on sample size, strain rate, and stress state, adiabatic heating due to plastic deformation [28] can be substantial (in some cases 50 C or more [29]), making quantification beyond the low strain modulus difficult. Finally, use of timeetemperature superposition to predict ballistic performance in thermosets is not firmly established, although it is suggested by work on polyurea and thermoset polyurethane elastomers, where an apparent rubberye glassy transition was observed moving from low to high strain rates at constant temperature [30e32] as well as in polybutadiene based polyurea [33]. Because of these challenges, ballistic characterization of polymers with high rate projectiles is an appropriate approach for assessing mechanical response. One way to quantify high rate dissipation capability is to measure the V50BL(P) [34], which generally relates to the velocity at which there is a 50 percent probability of penetration of an aluminum foil witness target behind the sample through transfer of fragments onto the witness target, or complete penetration of the sample and witness target by the incoming projectile. V50BL(P) has proved useful in the understanding of the dissipative capability of metals [35] and polymeric composites [36], and measuring V50BL(P) at a constant areal density provides data that are directly comparable [37]. In this work we characterize the ballistic performance in a series of model amine cured epoxy resins as a function of the glass transition temperature (Tg) of the cured polymer network, where the network Tg was systematically varied through the monomer stiffness, structure, and size. The effective strain rates for projectile impact were 104e105 s1, which is 10 times greater than typical SHPB tests and 105 times greater than typical quasi-static rates. Key factors that control the epoxy mechanical response at these high strain rates are identified. In particular, the ballistic performance of the network is controlled by the difference between the measurement temperature and the glass transition temperature of the network, while sub-Tg relaxations had no measurable effect. 2. Experimental 2.1. Materials Materials used in this study are provided in Fig. 1. Diglycidyl ether of bisphenol A (DGEBA) and diglycidyl ether of bisphenol F (DGEBF), i.e., EPON 825 and 862, respectively, were obtained from MillerStephenson. Polypropylene oxide based-Jeffamine diamines in various molecular weights (230, 400, 2000 and 4000 g/mol) and a triamine (T403, w440 g/mol) were provided by Huntsman. Diaminopropane (DAP), 3-aminophenyl sulfone (3DDS), 4-aminophenyl sulfone (4DDS), m-phenylenediamine (mPDA), tris(2-aminoethyl) amine and aniline were purchased from Sigma Aldrich. Cyclic diamine curing agents 4,40 -methylenebis(cyclohexylamine) (PACM) and 2,20 -dimethyl-4,40 -methylenebis(cyclohexylamine) (MPACM),
O
O
H2N
O
DGEBA
O
O
H2N
O
DGEBF
H2N
S
H2N
O
NH2
O
4DDS
Jeffamines H2N
S O
Aniline
NH2
DAP
3DDS
O NH2
H2N
NH2 x
NH2
O
NH2
NH2
MPACM
O
O H2N
NH2
PACM
O
N
TEATA
NH2
H2N
O
x
O
NH2 y NH2 z
T-403
Fig. 1. Epoxy resins diglycidyl ether of bisphenol A (DGEBA) and diglycidyl ether of bisphenol F (DGEBF) and curing agents 4,40 -methylenebis(cyclohexylamine) (PACM), 2,20 -dimethyl-4,40 -methylenebis(cyclohexylamine) (MPACM), polyetheramines (Jeffamines and T-403 where x þ y þ z ¼ 5e6), 4-aminophenyl sulfone (4DDS), 3aminophenyl sulfone (3DDS), m-phenylenediamine (mPDA), tris(2-aminoethyl)amine (TEATA), diaminopropane (DAP) and aniline.
were provided by Air Products. All epoxies and curing agents were used as received without further purification. Formulations were stoichiometric mixtures except where specifically noted. For V50BL(P) and dynamic mechanical analysis measurements, epoxies and curing agents were preheated to 60 C. They were then mixed together and stirred vigorously at 60 C for 5 min and were then poured into molds (6 in. 6 in. 0.25) and degassed under vacuum. All formulations except 3DDS and 4DDS were cured under a nitrogen purge with a cure cycle of 80 C for 2 h, 150 C for 8 h, 200 C for 2 h. Curing agents 3DDS and 4DDS were premelted and mixed with epoxy at 125 C under nitrogen; they were then cured at 125 C for 8 h, and 225 C for 4 h. 2.2. V50BL(P) measurement Ballistic impacts were carried out with a 0.22 caliber gas gun at room temperature (22 C). Relative humidity was between 45 and 60% for these measurements. A 5.56 mm diameter steel ball bearing (Type 302, 0.69 g) was used as a projectile to impact the target. Projectile speed was tracked with a Doppler radar (BR-3502, Infinition Inc.). The polymer target (nominal 6 cm 6 cm 0.64 cm) was sandwiched in a target frame with a circular opening of 5.08 cm in diameter. The projectile impacted on the target surface with zero degrees of obliquity. A witness plate (0.05 mm thick 2024-T3 aluminum foil) was placed 2 inches behind the target and was examined for penetration after each shot. Epoxy ballistic penetration was determined when the projectile or a fragment of the projectile or epoxy target penetrated the witness plate due to ballistic impact. Twelve targets were shot per specimen. The V50BL(P) ballistic performance was calculated by taking the arithmetic mean of the three highest non-penetrating and the three lowest complete penetrating impact velocities on the witness plate. For selected V50BL(P) ballistic events, high speed videography was performed to assess whether or not deflection occurred in the sample during impact. Two high speed cameras (SA1, Photron USA,
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Inc.) were used to generate stereo image pairs of the back face deformation of the target. The cameras were placed behind the target fixture, along the direction of the impacting projectile. The frame rate was set at 100,000 frames per second. The targeted area on the target was painted with a random pattern of black dots for displacement tracking. The stereo images were analyzed with a commercially available photogrammetric software program called ARAMIS (GOM GmbH, Germany, distributed by Trilion Quality Systems in the USA). 2.3. Dynamic mechanical analysis Dynamic mechanical analysis (DMA) was performed using a TA instruments Q800 on samples that were nominally 35 mm 12 mm 3 mm. Displacement was 7.5 mm at 1 Hz. 2.4. Mechanical property measurements Testing to determine the fracture toughness of epoxy resins adhered to ASTM D5045-99. Specimens were single-edge-notch bending (SNEB) geometry machined to 0.635 cm 1.27 cm 12 cm (B W L). The test span was set at 10.16 cm with a cross head speed of 10 mm/min. The pre-crack generated by tapping with a cryo-frozen razorblade, and care was taken to generate instantly propagated cracks, rather than non-propagated cracks [38,39]. Tensile properties were measured by preparing samples machined to the dimensions of Type IV specimens according to ASTM D638-10. The cross head speed was set to 5 mm/min and DIC (digital image correlation) [40] was used to obtain strain values for the duration of the test. For compression properties, cylindrical samples 1.27 cm 1.27 cm were prepared and tested using ASTM D695-10 as a guide. The cross head speed was set to 1.3 mm/min. Vertical displacement and DIC were used to calculate strain values. 3. Results and discussion To vary the resin Tg, samples were composed of either DGEBA, or a more flexible analog, DGEBF, and crosslinked with a variety of diamine curing agents, shown in Fig. 1. To obtain elastomers, high molecular weight Jeffamines and mixtures thereof were used (i.e., D400, D2000 and D4000, where x ¼ 6, 33 and 68, respectively). Room temperature glassy epoxies were obtained by curing with the low molecular weight Jeffamines (D230, and D400), DAP, and the cyclic curing agents PACM, MPACM, 3DDS, 4DDS and mPDA. The Tg, defined as the peak in tan d from DMA measurements [41,42], spans a range from about 58 C to 217 C, depending on the resin monomers, which is well above and below the V50BL(P) measurement temperature, T ¼ 22 C. V50BL(P) results for the resins are plotted as a function of T Tg in Fig. 2; these data are normalized to the value for DGEBA/PACM. As shown, in the high T Tg (rubbery) region, where curing agents D2000 and D4000 were used, V50BL(P) values are quite low, but increase rapidly as the resin Tg approaches the measurement temperature where T Tg ¼ 0. A peak in V50BL(P) values was observed for samples whose Tg is about 25 C above the measurement temperature, i.e., those cured with D400. It is interesting to note that the peak in V50BL(P) does not occur when the measurement temperature is equal to the resin Tg, as one might expect. Local heating induced in the material as a result of impact is one possible cause for this effect [43,44], resulting in a temperature shift for the V50BL(P) maximum by approximately 25 C. If the Tg is defined as the peak in loss modulus (E00 , see Table 1), then Tg values tend to shift down somewhat and the peak in V50BL(P) would occur 10e15 C above T Tg ¼ 0, still indicating the possibility of local adiabatic heating. Adiabatic heating during
Fig. 2. Plot of normalized V50BL(P) as a function of T Tg, where T ¼ 22 C, i.e., the measurement temperature. Thin dotted lines tie excess amine, stoichiometric and excess epoxy data points.
high rate impact events is well established [43,44]. During SplitHopkinson bar testing of DGEBA/D400 samples at 2000 s1 this heating effect was strong enough to push the compression sample into the rubbery regime during testing. However, since adiabatic heating can be a slow process, the magnitude of the temperature rise during the ballistic testing is unclear in Fig. 2. High speed infrared videography will be utilized in future experiments to quantify the heating effect. Another contributing factor for the maximum in V50BL(P) occurring at T Tg < 0 is a change in the ballistic failure mechanism for resins with these intermediate Tg values, which will be discussed in more detail later in this paper. After the V50BL(P) peak, decreasing T Tg (increasing resin Tg) values resulted in a steep drop in V50BL(P) values for the glassy materials back to a value of about 1.0. The V50BL(P) reaches a plateau for T Tg values below about 75 C (resin Tg > 100 C). Sub-Tg relaxations can influence high strain rate behavior [26] and are related to improvement in quasi-static toughness in epoxy composites [45]; however, the effect of these transitions on ballistic performance is unclear. The b-relaxation in DGEBAdiamine thermosets is not dependent on the curing agent [46] and has been shown to primarily involve motion of the hydroxypropyl ether groups [47,48] formed after curing as well as aromatic ring flipping in DGEBA itself [49], though these two processes may be interrelated as a slight increase in the bond angle between rings may be required for flipping to take place [50]. DGEBA and DGEBF can undergo similar relaxation in terms of the hydroxypropyl ether groups, but one would expect that ring
Table 1 Glass transition temperatures of select epoxy/diamine networks defined as either peak in tan d or peak in loss modulus from DMA measurements. Curing agent
DGEBA Tg ( C) tan d
DGEBA Tg ( C) E00
DGEBF Tg ( C) tan d
DGEBF Tg ( C) E00
D4000 D2000 D400 D230 DAP PACM MPACM 3DDS mPDA 4DDS
47 28 55 99 137 178 189 192 198 217
58 44 46 91 128 165 175 185 192 206
48 26 52 84 113 140 155 166 160 190
58 43 43 76 105 130 145 158 153 180
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flipping requiring a bond angle increase would be more difficult in DGEBA than in DGEBF. A feature of Fig. 2 is that there is very little difference in the performance of DGEBA and DGEBF resins. Generally, when cured with the same diamine, DGEBF results in a Tg that is lower. Given that Fig. 2 accounts for this Tg effect in plotting V50BL(P) versus T Tg, we see that the difference in flexibility between DGEBA and DGEBF does not influence the V50BL(P) behavior beyond the influence of Tg. Furthermore, the specific structure of the diamine also does not seem to matter. For example, one could envision that the resin cured with 4DDS could undergo ring spinning as a high rate dissipative mechanism, while that with 3DDS could not, due to its asymmetric structure. However, these materials fall on the same plateau as non-cyclic DAP and D230, indicating that such a mechanism may only have a subtle effect. The correlations between sub-Tg linear relaxation behavior on the non-linear properties of polymers is still debated. For example, early work on thermoplastic polycarbonate showed thermal history had more influence on impact strength than sub-Tg relaxations [51]. More recent investigations on semi-aromatic polyamides, polycarbonate, and crosslinking epoxies showed that secondary relaxations with cooperative character influenced yield behavior [52,53], and presumably the fracture toughness of the resin. In contrast, other researchers have shown that fracture properties of epoxy resins are dependent on the glass transition temperature and crosslink density, rather than intermolecular interactions that dictate sub-Tg cooperative motions [54]. In any case, the impact of sub-Tg relaxation processes in polymers on ballistic performance must be investigated in more detail. Specifically, the effective strain rates at the V50BL(P) projectile velocities in this paper (104e105 s1) may not be fast enough to excite these high frequency relaxations. For example, ring spinning
Fig. 4. Number of cracks >1 cm long in V50BL(P) samples of DGEBA and DGEBF cured with diamines as a function of T Tg. The circled region represents a transition in failure mechanism from shear plug failure to large scale “spalling”.
and hydroxypropyl ether rocking transitions occur on the time scale of 106 and 104 s, respectively [49]. So, these transitions may influence the energy dissipated at higher projectile velocities than those explored in this report, or these relaxations may be too fast to provide a toughening effect at the measurement temperature of 25 C. As such, a broader range of measurement temperatures and projectile velocities are currently being explored to determine when or if these relaxations influence ballistic performance.
Fig. 3. Images of DGEBA samples at V50BL(P) with various curing agents: a) D2000, Tg ¼ 28 C, b) D400, Tg ¼ 55 C, c) D230, Tg ¼ 99 C, d) PACM, Tg ¼ 178 C. Approximate site of “spalling” zone is shown by red circles.
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A variety of additional materials were studied, including DGEBA cured with two triamines: T-403 and TEATA, an aliphatic triamine (Fig. 1). These data are shown as circles in Fig. 2, and, while slightly lower, do not show a substantial deviation from the performance of the resins cured with diamines. Further, aniline, a monoamine chain extender that does not provide crosslinking, was also cured with DGEBA (triangle in Fig. 2), and it also does not show substantial deviation from the performance of the diamines. These samples provide further evidence that the a-relaxation provides a primary dissipation mechanism under high rates. Off-stoichiometric formulations of DGEBA with D400, D230 and PACM were also investigated. As shown in Fig. 2, thin dotted lines
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proceed through points from 15% excess epoxy, through the stoichiometric point, then through the 15% excess amine point. The Tg values of the excess epoxy materials are all considerably lower than the Tg of their stoichiometric counterpart, however the excess amine Tg values were only slightly lower. No evidence that offstoichiometric formulations significantly broadened the Tg was observed. The excess amine formulations improved V50BL(P) for the DGEBA/PACM and DGEBA/D230 sample sets, which are glasses with Tg values substantially higher than the measurement temperature. This Tg effect is likely due to the fact that excess epoxy (2 reactive sites per molecule) has a much higher likelihood of resulting in dangling chain ends or unreacted epoxy, while excess amines
Fig. 5. High speed camera images of the impact event for (a) DGEBA/D2000, (b) DGEBA/D400, and (c) DGEBA/PACM. Relative times are listed in the frames.
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(4 reactive sites per molecule) are more likely to propagate (crosslink or chain extend) the polymer network. The DGEBA/D400 set near the peak in V50BL(P), however, showed a decrease in both the excess epoxy and excess amine formulations, where excess epoxy dramatically reduced the V50BL(P) values. These results are in qualitative agreement with the work of Drzal et al., which showed that epoxy formulations with excess amine exhibited higher fracture toughness [55]. Photographic images of characteristic samples are presented in Fig. 3 for the sample exposed to impact velocities that were closest to the V50BL(P) value. Elastomeric samples, Fig. 3 (a), failed by shearing a plug of polymer resin the size of the projectile, and no fracture lines were observed. For samples where the measurement temperature was below Tg, significant fracture occurred, and increased with increasing Tg. A quantification of this is provided in Fig. 4, where the number of radial cracks greater than 1 cm were counted and averaged for all V50BL(P) specimens of a given composition. The number of fracture events increases steadily with decreasing T Tg. Figs. 3 and 4 shed additional insight into why a maximum in V50BL(P) is observed for samples with T Tg z 25 C. Samples with T Tg > 25 C fail by shearing a plug of polymer; therefore, the zone of deformation is fairly small for these samples, resulting in lower ballistic penetration velocities. In contrast, samples with T Tg z 25 C (i.e., DGEBA or DGEBF with D400) fail by extensive “spalling”, where a large cone of material fails in mixed mode fracture behind the site of projectile impact, dissipating large amounts of energy. High glass transition temperature samples exhibited smaller spalling zones, as shown in Fig. 3, resulting in lower V50BL(P) values for the high Tg resins. To investigate the ballistic failure process in more detail, we selected three samples for observations of the ballistic event using a high speed camera: (i) a rubbery, low Tg material (DGEBA cured with D2000), (ii) a glassy, high Tg material (DGEBA cured with PACM), and (iii) the formulation that showed the highest V50BL(P) value with a Tg just above room temperature (DGEBA cured with D400). These samples are highlighted in Figs. 2 and 4. Fig. 5(a), (b) and (c) provide single frame images of the ballistic event from the backside of the sample (i.e., opposite impact) for DGEBA cured with D2000, D400 and PACM, respectively. Full video files can be found in the supplemental information. The white spot on the sample in the initial frame was placed on the sample to indicate the desired location for impact. The displacement detection sensitivity was 0.5 mm, and the targets fractured or spalled before any significant deflection was observed. This means that the kinetic energy of the particle is only being transferred to plastic deformation, cracking, and other failure modes rather than elastic deflection of the sample. Supplementary video related to this article can be found at http://dx.doi.org/10.1016/j.polymer.2012.09.058. As shown in Fig. 5(a) and the video for DGEBA/D2000 in the supplemental information, no cracking occurs, and the projectile merely displaces a portion of material that is observed to be leading the projectile as it exits the sample. The situation is considerably different for the DGEBA/D400 sample (Fig. 5(b)), as the projectile impact results in immediate fracturing local to the impact, with the clear formation of spall on the side in view. Further radial cracking occurs subsequent to the initial impact, and these cracks propagate for a few milliseconds after impact. A similar situation occurs for the DGEBA/PACM sample (Fig. 5(c)), but here the number of cracks is considerably higher than that of the DGEBA/D400 and the spall area behind the impact location is smaller than that of DGEBA/D400. The failure during ballistic impact is a complex process with mixed mode fracture behind the projectile impact site. Specifically, mode I is the dominant fracturing growing radially from the impact
a
b
Fig. 6. Representative stress (s) versus strain (ε) curves under (a) compression and (b) tension for representative samples of DGEBA cured with PACM, D400 or D2000.
site, but a zone of compression behind the projectile and a combination of tensile and shear deformation occur adjacent to the projectile. In order to explore these dissipation processes more systematically, we performed quasi-static compression, tensile and fracture toughness measurements on the same materials (i.e., DGEBA cured with D2000, D400 and PACM) at room temperature. Fig. 6(a) provides representative compression stressestrain curves for the three samples. As shown, while DGEBA/PACM has a higher strain at yield and yield strength, the overall apparent toughness (area under the curve) for DGEBA/D400 is much higher due to its ability to plastically deform after yield without failing. In fact, the DGEBA/D400 epoxy did not break until about 80% strain. In contrast to these, DGEBA/D2000 behaves like a rubber and exhibits high strain at break, but very low apparent toughness. Specific details associated with the compression tests are provided in Table 2. Results for quasi-static tension measurements are provided in Fig. 6(b). Here, we see that DGEBA/PACM and DGEBA/D400 have
Table 2 Compression properties of DGEBA cured with three different diamines. Curing agent
Compressive strength (MPa)
Strain at break
Compressive yield strength (MPa)
Compressive modulus (MPa)
Apparent toughness (MPa)
PACM D400 D2000
278 23 428 8 3.6 0.6
0.43 0.02 0.80 0.01 0.48 0.05
119.9 2.8 67.5 0.8 N/aa
1299 180 1288 32 N/aa
59.9 4.0 95.3 2.5 0.6 0.1
a
D2000 samples essentially yielded immediately at room temperature.
D.B. Knorr Jr. et al. / Polymer 53 (2012) 5917e5923 Table 3 Tension properties of DGEBA cured with three different diamines. Curing agent
Tensile strength (MPa)
Strain at break
Tensile modulus (MPa)
Apparent toughness (MPa)
KIC (MPa m½)
PACM D400 D2000
43.2 7.5 57.2 3.5 0.798 0.17
0.03 0.009 0.032 0.002 0.29 0.07
2180 140 2556 163 N/aa
0.82 0.42 1.81 0.81 0.14 0.07
0.69 0.08 1.48 0.2 N/aa
a D2000 samples essentially yielded immediately at room temperature and tore during the fracture toughness test.
similar moduli, but DGEBA/D400 has a higher tensile strength and again has the ability to yield prior to failure, again providing a higher apparent toughness than the DGEBA/PACM sample. Further, DGEBA/D2000 shows low apparent toughness in tension. Details regarding the tensile measurements can be found in Table 3. Finally, quasi-static mode 1 fracture toughness is consistent with ballistic measurements as DGEBA/PACM and DGEBA/D400 exhibit a room temperature fracture toughness of 0.7 and 1.5 MPa m0.5, respectively (Table 3). The quasi-static mechanical results illustrate that, in tension and in compression, the DGEBA/D400 sample has the ability to dissipate more energy than that of either DGEBA/PACM or DGEBA/D2000 due to its higher apparent toughness. Quasi-static fracture toughness also suggests that DGEBA/D400 would dissipate more energy during cracking than DGEBA/PACM. Given these results, a mixed mode failure mechanism for the ballistic event is a reasonable conclusion. Furthermore, the above result that the V50BL(P) of a material is strongly related to T Tg seems reasonable in that all of the quasistatic mechanical properties are also strong functions of T Tg as well. That is, glassy amine cured epoxies nearer their Tg, like the DGEBA/D400, are generally expected to be tougher and have a higher fracture toughness allowing for improved energy dissipation. 4. Conclusions These results demonstrate that Tg (or more specifically, the difference between the Tg and the temperature at which a ballistic impact occurs) is one dominant factor influencing the ballistic performance of epoxy resins. High frequency relaxation processes such as ring flipping or hydroxypropyl ether motion showed no measurable effect. V50BL(P) shows a peak for materials whose Tg is about 25 C above the measurement temperature, and the high V50BL(P) values are associated with a large zone of “spalling” behind the impact site. This work suggests that timeetemperature superposition might be applied to qualitatively assess ballistic performance and this is a focus of ongoing research. Quasi-static compression, tension and fracture toughness for selected resins was consistent with ballistic performance in the epoxies. The research has important implications for light weight protective materials in a broad range of industries including military, aerospace, transportation, and construction technologies. References [1] McGrath LM, Parnas RS, King SH, Schroeder JL, Fischer DA, Lenhart JL. Polymer 2008;49(4):999e1014. [2] Nielsen LE, Landel RF. Mechanical properties of polymers and composites. New York: Marcel Dekker; 1994. [3] Grossman E, Gouzman I. Nuclear Instruments & Methods in Physics Research Section B-Beam Interactions with Materials and Atoms 2003;208:48e57. [4] National Research Council. Opportunities in protection materials science and technology for future army applications. The National Academies Press; 2011. [5] Zubeldia A, Larranaga M, Remiro P, Mondragon I. Journal of Polymer Science Part B-Polymer Physics 2004;42(21):3920e33. [6] Cardwell BJ, Yee AF. Polymer 1993;34(8):1695e701.
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Polymer journal homepage: www.elsevier.com/locate/polymer
Glass transition dependence of ultrahigh strain rate response in amine cured epoxy resins Daniel B. Knorr Jr., Jian H. Yu, Adam D. Richardson, Mark D. Hindenlang, Ian M. McAninch, John J. La Scala, Joseph L. Lenhart* U.S. Army Research Laboratory, Aberdeen Proving Ground, MD 21009, United States
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 July 2012 Received in revised form 25 September 2012 Accepted 29 September 2012 Available online 17 October 2012
Ultrahigh strain rate performance in a series of model amine cured epoxy resins was investigated as a function of the glass transition temperature (Tg) of the cured polymer network, where the network Tg was systematically varied through the monomer stiffness, structure, and size. The high rate response was characterized in terms of a projectile penetration velocity, V50BL(P) (ballistic limit, protection criteria), which describes the projectile velocity with a 50% probability of sample penetration. One factor that dictates the ballistic performance of the epoxy networks, at effective rates of 104e105 s1, is the difference between the measurement temperature and the glass transition temperature of the network. Sub-Tg relaxations did not have a measurable effect on ballistic performance, and neither did the monomer structure and functionality outside of the influence of the resin Tg, while off-stoichiometric (excess amine) formulations improved V50BL(P) slightly with high Tg epoxies. The results have implications in protective materials for military, aerospace, transportation, and construction industries, where high strain rate insults from airborne debris, high rate collisions, and natural events are increasingly considered during product design. Published by Elsevier Ltd.
Keywords: High strain rate Epoxy-amine Glass transition
1. Introduction Epoxy resins are exploited heavily in both commercial and military technologies including fiber reinforced composites, electronics packaging, adhesives, transportation, construction, and infrastructure applications. A combination of processing versatility, chemical diversity, chemical resistance, and environmental stability makes epoxies useful for these broad applications [1]. However, a critical challenge is that these polymers are notoriously brittle [2]. In addition, high rate mechanical performance is increasingly important for light weight protective equipment in military, aerospace, transportation, and construction industries [3,4], yet epoxies have not been designed to function in these extreme mechanical environments. Substantial research has focused on the quasi-static mechanical properties of epoxy resins and the ballistic performance of many industrial formulations has been evaluated. However, these industrial formulations are complex and proprietary concoctions of many additives including: * Corresponding author. U.S. Army Research Laboratory, 4600 Deer Creek Loop, Aberdeen Proving Ground, MD 21005, United States. Tel.: þ1 410 306 1940; fax: þ1 410 306 0676. E-mail address: [email protected] (J.L. Lenhart). 0032-3861/$ e see front matter Published by Elsevier Ltd. http://dx.doi.org/10.1016/j.polymer.2012.09.058
monomer mixtures, chain extenders, viscosity modifiers, toughening agents, degassing agents, and even low volatility plasticizers, making fundamental insight difficult to extract. Timeetemperature superposition is potentially useful for predicting high strain rate behavior, but has only been demonstrated in select cases in the high strain rate regime [5,6]. Therefore, a critical knowledge gap exists in terms of the chemical, physical, and structural factors that govern high rate response in polymer networks. Reinforced epoxy resins are indispensable components in many protective technologies. The complex energy dissipative mechanisms in these composites are key, and take a variety of forms involving tensile failure, elastic deformation and kinetic energy transferred to the moving portion of the composite [7,8]. Zee and Hsieh found that, within a fiber reinforced polymerematrix composite itself, four processes are responsible for energy dissipation: (i) heat, (ii) fiber breakage/deformation, (iii) interfacial delamination and (iv) matrix cracking/deformation [9]. While the first process is common to most dissipative mechanisms, the second two are unique to composites, and have been studied extensively [10e18]. Mechanical response of the polymer matrix alone can be responsible for 20e35% of the total energy dissipation [9]. Despite this potential contribution, few studies have focused on improving high rate dissipative capabilities or on obtaining
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a fundamental understanding of matrix-only performance in ballistic environments [19]. High rate characterization of metals and ceramics is common with split-Hopkinson pressure bar (SHPB) techniques [20,21]. However, measurements of polymers are particularly difficult due to their large strain deformation capability, where inertial effects and complex stress states can convolute the mechanical response [22,23]. For example, deviations from linearity of compressive yield stress as a function of log strain rate are common with SHPB testing of polymers. Typically, yield stress increases from linear behavior at higher strain rates [24,25]. While some groups have attributed this to molecular relaxations in polymers at high strain rates [26,27], friction at the polymerebar interface can equally explain these results. Our own testing has demonstrated that lubricants are ineffective at preventing specimen “barreling” during high rate compression testing. In addition, depending on sample size, strain rate, and stress state, adiabatic heating due to plastic deformation [28] can be substantial (in some cases 50 C or more [29]), making quantification beyond the low strain modulus difficult. Finally, use of timeetemperature superposition to predict ballistic performance in thermosets is not firmly established, although it is suggested by work on polyurea and thermoset polyurethane elastomers, where an apparent rubberye glassy transition was observed moving from low to high strain rates at constant temperature [30e32] as well as in polybutadiene based polyurea [33]. Because of these challenges, ballistic characterization of polymers with high rate projectiles is an appropriate approach for assessing mechanical response. One way to quantify high rate dissipation capability is to measure the V50BL(P) [34], which generally relates to the velocity at which there is a 50 percent probability of penetration of an aluminum foil witness target behind the sample through transfer of fragments onto the witness target, or complete penetration of the sample and witness target by the incoming projectile. V50BL(P) has proved useful in the understanding of the dissipative capability of metals [35] and polymeric composites [36], and measuring V50BL(P) at a constant areal density provides data that are directly comparable [37]. In this work we characterize the ballistic performance in a series of model amine cured epoxy resins as a function of the glass transition temperature (Tg) of the cured polymer network, where the network Tg was systematically varied through the monomer stiffness, structure, and size. The effective strain rates for projectile impact were 104e105 s1, which is 10 times greater than typical SHPB tests and 105 times greater than typical quasi-static rates. Key factors that control the epoxy mechanical response at these high strain rates are identified. In particular, the ballistic performance of the network is controlled by the difference between the measurement temperature and the glass transition temperature of the network, while sub-Tg relaxations had no measurable effect. 2. Experimental 2.1. Materials Materials used in this study are provided in Fig. 1. Diglycidyl ether of bisphenol A (DGEBA) and diglycidyl ether of bisphenol F (DGEBF), i.e., EPON 825 and 862, respectively, were obtained from MillerStephenson. Polypropylene oxide based-Jeffamine diamines in various molecular weights (230, 400, 2000 and 4000 g/mol) and a triamine (T403, w440 g/mol) were provided by Huntsman. Diaminopropane (DAP), 3-aminophenyl sulfone (3DDS), 4-aminophenyl sulfone (4DDS), m-phenylenediamine (mPDA), tris(2-aminoethyl) amine and aniline were purchased from Sigma Aldrich. Cyclic diamine curing agents 4,40 -methylenebis(cyclohexylamine) (PACM) and 2,20 -dimethyl-4,40 -methylenebis(cyclohexylamine) (MPACM),
O
O
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DGEBA
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O
H2N
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DGEBF
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Jeffamines H2N
S O
Aniline
NH2
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3DDS
O NH2
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NH2 x
NH2
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MPACM
O
O H2N
NH2
PACM
O
N
TEATA
NH2
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x
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T-403
Fig. 1. Epoxy resins diglycidyl ether of bisphenol A (DGEBA) and diglycidyl ether of bisphenol F (DGEBF) and curing agents 4,40 -methylenebis(cyclohexylamine) (PACM), 2,20 -dimethyl-4,40 -methylenebis(cyclohexylamine) (MPACM), polyetheramines (Jeffamines and T-403 where x þ y þ z ¼ 5e6), 4-aminophenyl sulfone (4DDS), 3aminophenyl sulfone (3DDS), m-phenylenediamine (mPDA), tris(2-aminoethyl)amine (TEATA), diaminopropane (DAP) and aniline.
were provided by Air Products. All epoxies and curing agents were used as received without further purification. Formulations were stoichiometric mixtures except where specifically noted. For V50BL(P) and dynamic mechanical analysis measurements, epoxies and curing agents were preheated to 60 C. They were then mixed together and stirred vigorously at 60 C for 5 min and were then poured into molds (6 in. 6 in. 0.25) and degassed under vacuum. All formulations except 3DDS and 4DDS were cured under a nitrogen purge with a cure cycle of 80 C for 2 h, 150 C for 8 h, 200 C for 2 h. Curing agents 3DDS and 4DDS were premelted and mixed with epoxy at 125 C under nitrogen; they were then cured at 125 C for 8 h, and 225 C for 4 h. 2.2. V50BL(P) measurement Ballistic impacts were carried out with a 0.22 caliber gas gun at room temperature (22 C). Relative humidity was between 45 and 60% for these measurements. A 5.56 mm diameter steel ball bearing (Type 302, 0.69 g) was used as a projectile to impact the target. Projectile speed was tracked with a Doppler radar (BR-3502, Infinition Inc.). The polymer target (nominal 6 cm 6 cm 0.64 cm) was sandwiched in a target frame with a circular opening of 5.08 cm in diameter. The projectile impacted on the target surface with zero degrees of obliquity. A witness plate (0.05 mm thick 2024-T3 aluminum foil) was placed 2 inches behind the target and was examined for penetration after each shot. Epoxy ballistic penetration was determined when the projectile or a fragment of the projectile or epoxy target penetrated the witness plate due to ballistic impact. Twelve targets were shot per specimen. The V50BL(P) ballistic performance was calculated by taking the arithmetic mean of the three highest non-penetrating and the three lowest complete penetrating impact velocities on the witness plate. For selected V50BL(P) ballistic events, high speed videography was performed to assess whether or not deflection occurred in the sample during impact. Two high speed cameras (SA1, Photron USA,
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Inc.) were used to generate stereo image pairs of the back face deformation of the target. The cameras were placed behind the target fixture, along the direction of the impacting projectile. The frame rate was set at 100,000 frames per second. The targeted area on the target was painted with a random pattern of black dots for displacement tracking. The stereo images were analyzed with a commercially available photogrammetric software program called ARAMIS (GOM GmbH, Germany, distributed by Trilion Quality Systems in the USA). 2.3. Dynamic mechanical analysis Dynamic mechanical analysis (DMA) was performed using a TA instruments Q800 on samples that were nominally 35 mm 12 mm 3 mm. Displacement was 7.5 mm at 1 Hz. 2.4. Mechanical property measurements Testing to determine the fracture toughness of epoxy resins adhered to ASTM D5045-99. Specimens were single-edge-notch bending (SNEB) geometry machined to 0.635 cm 1.27 cm 12 cm (B W L). The test span was set at 10.16 cm with a cross head speed of 10 mm/min. The pre-crack generated by tapping with a cryo-frozen razorblade, and care was taken to generate instantly propagated cracks, rather than non-propagated cracks [38,39]. Tensile properties were measured by preparing samples machined to the dimensions of Type IV specimens according to ASTM D638-10. The cross head speed was set to 5 mm/min and DIC (digital image correlation) [40] was used to obtain strain values for the duration of the test. For compression properties, cylindrical samples 1.27 cm 1.27 cm were prepared and tested using ASTM D695-10 as a guide. The cross head speed was set to 1.3 mm/min. Vertical displacement and DIC were used to calculate strain values. 3. Results and discussion To vary the resin Tg, samples were composed of either DGEBA, or a more flexible analog, DGEBF, and crosslinked with a variety of diamine curing agents, shown in Fig. 1. To obtain elastomers, high molecular weight Jeffamines and mixtures thereof were used (i.e., D400, D2000 and D4000, where x ¼ 6, 33 and 68, respectively). Room temperature glassy epoxies were obtained by curing with the low molecular weight Jeffamines (D230, and D400), DAP, and the cyclic curing agents PACM, MPACM, 3DDS, 4DDS and mPDA. The Tg, defined as the peak in tan d from DMA measurements [41,42], spans a range from about 58 C to 217 C, depending on the resin monomers, which is well above and below the V50BL(P) measurement temperature, T ¼ 22 C. V50BL(P) results for the resins are plotted as a function of T Tg in Fig. 2; these data are normalized to the value for DGEBA/PACM. As shown, in the high T Tg (rubbery) region, where curing agents D2000 and D4000 were used, V50BL(P) values are quite low, but increase rapidly as the resin Tg approaches the measurement temperature where T Tg ¼ 0. A peak in V50BL(P) values was observed for samples whose Tg is about 25 C above the measurement temperature, i.e., those cured with D400. It is interesting to note that the peak in V50BL(P) does not occur when the measurement temperature is equal to the resin Tg, as one might expect. Local heating induced in the material as a result of impact is one possible cause for this effect [43,44], resulting in a temperature shift for the V50BL(P) maximum by approximately 25 C. If the Tg is defined as the peak in loss modulus (E00 , see Table 1), then Tg values tend to shift down somewhat and the peak in V50BL(P) would occur 10e15 C above T Tg ¼ 0, still indicating the possibility of local adiabatic heating. Adiabatic heating during
Fig. 2. Plot of normalized V50BL(P) as a function of T Tg, where T ¼ 22 C, i.e., the measurement temperature. Thin dotted lines tie excess amine, stoichiometric and excess epoxy data points.
high rate impact events is well established [43,44]. During SplitHopkinson bar testing of DGEBA/D400 samples at 2000 s1 this heating effect was strong enough to push the compression sample into the rubbery regime during testing. However, since adiabatic heating can be a slow process, the magnitude of the temperature rise during the ballistic testing is unclear in Fig. 2. High speed infrared videography will be utilized in future experiments to quantify the heating effect. Another contributing factor for the maximum in V50BL(P) occurring at T Tg < 0 is a change in the ballistic failure mechanism for resins with these intermediate Tg values, which will be discussed in more detail later in this paper. After the V50BL(P) peak, decreasing T Tg (increasing resin Tg) values resulted in a steep drop in V50BL(P) values for the glassy materials back to a value of about 1.0. The V50BL(P) reaches a plateau for T Tg values below about 75 C (resin Tg > 100 C). Sub-Tg relaxations can influence high strain rate behavior [26] and are related to improvement in quasi-static toughness in epoxy composites [45]; however, the effect of these transitions on ballistic performance is unclear. The b-relaxation in DGEBAdiamine thermosets is not dependent on the curing agent [46] and has been shown to primarily involve motion of the hydroxypropyl ether groups [47,48] formed after curing as well as aromatic ring flipping in DGEBA itself [49], though these two processes may be interrelated as a slight increase in the bond angle between rings may be required for flipping to take place [50]. DGEBA and DGEBF can undergo similar relaxation in terms of the hydroxypropyl ether groups, but one would expect that ring
Table 1 Glass transition temperatures of select epoxy/diamine networks defined as either peak in tan d or peak in loss modulus from DMA measurements. Curing agent
DGEBA Tg ( C) tan d
DGEBA Tg ( C) E00
DGEBF Tg ( C) tan d
DGEBF Tg ( C) E00
D4000 D2000 D400 D230 DAP PACM MPACM 3DDS mPDA 4DDS
47 28 55 99 137 178 189 192 198 217
58 44 46 91 128 165 175 185 192 206
48 26 52 84 113 140 155 166 160 190
58 43 43 76 105 130 145 158 153 180
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flipping requiring a bond angle increase would be more difficult in DGEBA than in DGEBF. A feature of Fig. 2 is that there is very little difference in the performance of DGEBA and DGEBF resins. Generally, when cured with the same diamine, DGEBF results in a Tg that is lower. Given that Fig. 2 accounts for this Tg effect in plotting V50BL(P) versus T Tg, we see that the difference in flexibility between DGEBA and DGEBF does not influence the V50BL(P) behavior beyond the influence of Tg. Furthermore, the specific structure of the diamine also does not seem to matter. For example, one could envision that the resin cured with 4DDS could undergo ring spinning as a high rate dissipative mechanism, while that with 3DDS could not, due to its asymmetric structure. However, these materials fall on the same plateau as non-cyclic DAP and D230, indicating that such a mechanism may only have a subtle effect. The correlations between sub-Tg linear relaxation behavior on the non-linear properties of polymers is still debated. For example, early work on thermoplastic polycarbonate showed thermal history had more influence on impact strength than sub-Tg relaxations [51]. More recent investigations on semi-aromatic polyamides, polycarbonate, and crosslinking epoxies showed that secondary relaxations with cooperative character influenced yield behavior [52,53], and presumably the fracture toughness of the resin. In contrast, other researchers have shown that fracture properties of epoxy resins are dependent on the glass transition temperature and crosslink density, rather than intermolecular interactions that dictate sub-Tg cooperative motions [54]. In any case, the impact of sub-Tg relaxation processes in polymers on ballistic performance must be investigated in more detail. Specifically, the effective strain rates at the V50BL(P) projectile velocities in this paper (104e105 s1) may not be fast enough to excite these high frequency relaxations. For example, ring spinning
Fig. 4. Number of cracks >1 cm long in V50BL(P) samples of DGEBA and DGEBF cured with diamines as a function of T Tg. The circled region represents a transition in failure mechanism from shear plug failure to large scale “spalling”.
and hydroxypropyl ether rocking transitions occur on the time scale of 106 and 104 s, respectively [49]. So, these transitions may influence the energy dissipated at higher projectile velocities than those explored in this report, or these relaxations may be too fast to provide a toughening effect at the measurement temperature of 25 C. As such, a broader range of measurement temperatures and projectile velocities are currently being explored to determine when or if these relaxations influence ballistic performance.
Fig. 3. Images of DGEBA samples at V50BL(P) with various curing agents: a) D2000, Tg ¼ 28 C, b) D400, Tg ¼ 55 C, c) D230, Tg ¼ 99 C, d) PACM, Tg ¼ 178 C. Approximate site of “spalling” zone is shown by red circles.
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A variety of additional materials were studied, including DGEBA cured with two triamines: T-403 and TEATA, an aliphatic triamine (Fig. 1). These data are shown as circles in Fig. 2, and, while slightly lower, do not show a substantial deviation from the performance of the resins cured with diamines. Further, aniline, a monoamine chain extender that does not provide crosslinking, was also cured with DGEBA (triangle in Fig. 2), and it also does not show substantial deviation from the performance of the diamines. These samples provide further evidence that the a-relaxation provides a primary dissipation mechanism under high rates. Off-stoichiometric formulations of DGEBA with D400, D230 and PACM were also investigated. As shown in Fig. 2, thin dotted lines
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proceed through points from 15% excess epoxy, through the stoichiometric point, then through the 15% excess amine point. The Tg values of the excess epoxy materials are all considerably lower than the Tg of their stoichiometric counterpart, however the excess amine Tg values were only slightly lower. No evidence that offstoichiometric formulations significantly broadened the Tg was observed. The excess amine formulations improved V50BL(P) for the DGEBA/PACM and DGEBA/D230 sample sets, which are glasses with Tg values substantially higher than the measurement temperature. This Tg effect is likely due to the fact that excess epoxy (2 reactive sites per molecule) has a much higher likelihood of resulting in dangling chain ends or unreacted epoxy, while excess amines
Fig. 5. High speed camera images of the impact event for (a) DGEBA/D2000, (b) DGEBA/D400, and (c) DGEBA/PACM. Relative times are listed in the frames.
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(4 reactive sites per molecule) are more likely to propagate (crosslink or chain extend) the polymer network. The DGEBA/D400 set near the peak in V50BL(P), however, showed a decrease in both the excess epoxy and excess amine formulations, where excess epoxy dramatically reduced the V50BL(P) values. These results are in qualitative agreement with the work of Drzal et al., which showed that epoxy formulations with excess amine exhibited higher fracture toughness [55]. Photographic images of characteristic samples are presented in Fig. 3 for the sample exposed to impact velocities that were closest to the V50BL(P) value. Elastomeric samples, Fig. 3 (a), failed by shearing a plug of polymer resin the size of the projectile, and no fracture lines were observed. For samples where the measurement temperature was below Tg, significant fracture occurred, and increased with increasing Tg. A quantification of this is provided in Fig. 4, where the number of radial cracks greater than 1 cm were counted and averaged for all V50BL(P) specimens of a given composition. The number of fracture events increases steadily with decreasing T Tg. Figs. 3 and 4 shed additional insight into why a maximum in V50BL(P) is observed for samples with T Tg z 25 C. Samples with T Tg > 25 C fail by shearing a plug of polymer; therefore, the zone of deformation is fairly small for these samples, resulting in lower ballistic penetration velocities. In contrast, samples with T Tg z 25 C (i.e., DGEBA or DGEBF with D400) fail by extensive “spalling”, where a large cone of material fails in mixed mode fracture behind the site of projectile impact, dissipating large amounts of energy. High glass transition temperature samples exhibited smaller spalling zones, as shown in Fig. 3, resulting in lower V50BL(P) values for the high Tg resins. To investigate the ballistic failure process in more detail, we selected three samples for observations of the ballistic event using a high speed camera: (i) a rubbery, low Tg material (DGEBA cured with D2000), (ii) a glassy, high Tg material (DGEBA cured with PACM), and (iii) the formulation that showed the highest V50BL(P) value with a Tg just above room temperature (DGEBA cured with D400). These samples are highlighted in Figs. 2 and 4. Fig. 5(a), (b) and (c) provide single frame images of the ballistic event from the backside of the sample (i.e., opposite impact) for DGEBA cured with D2000, D400 and PACM, respectively. Full video files can be found in the supplemental information. The white spot on the sample in the initial frame was placed on the sample to indicate the desired location for impact. The displacement detection sensitivity was 0.5 mm, and the targets fractured or spalled before any significant deflection was observed. This means that the kinetic energy of the particle is only being transferred to plastic deformation, cracking, and other failure modes rather than elastic deflection of the sample. Supplementary video related to this article can be found at http://dx.doi.org/10.1016/j.polymer.2012.09.058. As shown in Fig. 5(a) and the video for DGEBA/D2000 in the supplemental information, no cracking occurs, and the projectile merely displaces a portion of material that is observed to be leading the projectile as it exits the sample. The situation is considerably different for the DGEBA/D400 sample (Fig. 5(b)), as the projectile impact results in immediate fracturing local to the impact, with the clear formation of spall on the side in view. Further radial cracking occurs subsequent to the initial impact, and these cracks propagate for a few milliseconds after impact. A similar situation occurs for the DGEBA/PACM sample (Fig. 5(c)), but here the number of cracks is considerably higher than that of the DGEBA/D400 and the spall area behind the impact location is smaller than that of DGEBA/D400. The failure during ballistic impact is a complex process with mixed mode fracture behind the projectile impact site. Specifically, mode I is the dominant fracturing growing radially from the impact
a
b
Fig. 6. Representative stress (s) versus strain (ε) curves under (a) compression and (b) tension for representative samples of DGEBA cured with PACM, D400 or D2000.
site, but a zone of compression behind the projectile and a combination of tensile and shear deformation occur adjacent to the projectile. In order to explore these dissipation processes more systematically, we performed quasi-static compression, tensile and fracture toughness measurements on the same materials (i.e., DGEBA cured with D2000, D400 and PACM) at room temperature. Fig. 6(a) provides representative compression stressestrain curves for the three samples. As shown, while DGEBA/PACM has a higher strain at yield and yield strength, the overall apparent toughness (area under the curve) for DGEBA/D400 is much higher due to its ability to plastically deform after yield without failing. In fact, the DGEBA/D400 epoxy did not break until about 80% strain. In contrast to these, DGEBA/D2000 behaves like a rubber and exhibits high strain at break, but very low apparent toughness. Specific details associated with the compression tests are provided in Table 2. Results for quasi-static tension measurements are provided in Fig. 6(b). Here, we see that DGEBA/PACM and DGEBA/D400 have
Table 2 Compression properties of DGEBA cured with three different diamines. Curing agent
Compressive strength (MPa)
Strain at break
Compressive yield strength (MPa)
Compressive modulus (MPa)
Apparent toughness (MPa)
PACM D400 D2000
278 23 428 8 3.6 0.6
0.43 0.02 0.80 0.01 0.48 0.05
119.9 2.8 67.5 0.8 N/aa
1299 180 1288 32 N/aa
59.9 4.0 95.3 2.5 0.6 0.1
a
D2000 samples essentially yielded immediately at room temperature.
D.B. Knorr Jr. et al. / Polymer 53 (2012) 5917e5923 Table 3 Tension properties of DGEBA cured with three different diamines. Curing agent
Tensile strength (MPa)
Strain at break
Tensile modulus (MPa)
Apparent toughness (MPa)
KIC (MPa m½)
PACM D400 D2000
43.2 7.5 57.2 3.5 0.798 0.17
0.03 0.009 0.032 0.002 0.29 0.07
2180 140 2556 163 N/aa
0.82 0.42 1.81 0.81 0.14 0.07
0.69 0.08 1.48 0.2 N/aa
a D2000 samples essentially yielded immediately at room temperature and tore during the fracture toughness test.
similar moduli, but DGEBA/D400 has a higher tensile strength and again has the ability to yield prior to failure, again providing a higher apparent toughness than the DGEBA/PACM sample. Further, DGEBA/D2000 shows low apparent toughness in tension. Details regarding the tensile measurements can be found in Table 3. Finally, quasi-static mode 1 fracture toughness is consistent with ballistic measurements as DGEBA/PACM and DGEBA/D400 exhibit a room temperature fracture toughness of 0.7 and 1.5 MPa m0.5, respectively (Table 3). The quasi-static mechanical results illustrate that, in tension and in compression, the DGEBA/D400 sample has the ability to dissipate more energy than that of either DGEBA/PACM or DGEBA/D2000 due to its higher apparent toughness. Quasi-static fracture toughness also suggests that DGEBA/D400 would dissipate more energy during cracking than DGEBA/PACM. Given these results, a mixed mode failure mechanism for the ballistic event is a reasonable conclusion. Furthermore, the above result that the V50BL(P) of a material is strongly related to T Tg seems reasonable in that all of the quasistatic mechanical properties are also strong functions of T Tg as well. That is, glassy amine cured epoxies nearer their Tg, like the DGEBA/D400, are generally expected to be tougher and have a higher fracture toughness allowing for improved energy dissipation. 4. Conclusions These results demonstrate that Tg (or more specifically, the difference between the Tg and the temperature at which a ballistic impact occurs) is one dominant factor influencing the ballistic performance of epoxy resins. High frequency relaxation processes such as ring flipping or hydroxypropyl ether motion showed no measurable effect. V50BL(P) shows a peak for materials whose Tg is about 25 C above the measurement temperature, and the high V50BL(P) values are associated with a large zone of “spalling” behind the impact site. This work suggests that timeetemperature superposition might be applied to qualitatively assess ballistic performance and this is a focus of ongoing research. Quasi-static compression, tension and fracture toughness for selected resins was consistent with ballistic performance in the epoxies. The research has important implications for light weight protective materials in a broad range of industries including military, aerospace, transportation, and construction technologies. References [1] McGrath LM, Parnas RS, King SH, Schroeder JL, Fischer DA, Lenhart JL. Polymer 2008;49(4):999e1014. [2] Nielsen LE, Landel RF. Mechanical properties of polymers and composites. New York: Marcel Dekker; 1994. [3] Grossman E, Gouzman I. Nuclear Instruments & Methods in Physics Research Section B-Beam Interactions with Materials and Atoms 2003;208:48e57. [4] National Research Council. Opportunities in protection materials science and technology for future army applications. The National Academies Press; 2011. [5] Zubeldia A, Larranaga M, Remiro P, Mondragon I. Journal of Polymer Science Part B-Polymer Physics 2004;42(21):3920e33. [6] Cardwell BJ, Yee AF. Polymer 1993;34(8):1695e701.
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