Dynamic Mechanical Analysis 

of E-Beam and Thermally Curable 

IPN Thermosets 

by Robert E. Jensen, Giuseppe R. Palmese, and 

Steven H. M&night 

ARL-TR-2698 March 2002 

Approved for public release; distribution is unlimited. 
. 

The findings in this report are not to be construed as an 

official Department of the Army position unless so 

designated by other authorized documents. 

Citation of manufacturer’s or trade names does not 

constitute an official endorsement or approval of the use 

thereof. 

Destroy this report when it is no longer needed. Do not 

return it to the originator. 
Army Research Laboratory 

Aberdeen Proving Ground, MD 21005-5069 

ARL-TR-2698 March 2002 

. 

Dynamic Mechanical Analysis 

of E-Beam and Thermally Curable 

IPN Thermosets 

Robert E. Jensen and Steven H. M&night 

Weapons and Materials Research Directorate, ARL 

Giuseppe R. Palmese 

Drexel University 

. 

Amroved for uublic release: distribution is unlimited. 
Abstract 

Interpenetrating polymer network (IPN) materials possess unique properties that 

can be exploited for multicomponent composite armor systems, including 

toughness and thermal stability. E-beam curing of composites and adhesives 

offers advantages, such as reduced cure shrinkages, over traditional autoclave 

processing by curing multiple resins through the thickness for thick-section 

composites in a single step. Because of the complexity and thickness of 

composite integral armor structures, e-beam curing is an attractive processing 

method for Future Combat System applications, and the development of these 

new resins permits this approach. This research seeks to discern the differences 

in viscoelastic behavior between traditional thermal and e-beam curing in IPN 

resins, as these viscoelastic characteristics may prove important for composite 

integral armor applications. 
Acknowledgments 

Y 

This research was supported in part by an appointment to the Research 

Participation Program at the U.S. Army Research Laboratory (ARL) administered 

by the Oak Ridge Institute for Science and Education through an interagency 

agreement between the U.S. Department of Energy and ARL. The authors also 

wish to acknowledge the ARL Composite Materials Research Materials Center of 

Excellence and the University of Delaware Center for Composite Materials 

whose participation through the joint contract DAALOl-96-2-0048 provided 

additional support. 

. . . 

Ill 
L 

iv 
Contents 

. 

, 

Acknowledgments . . . 

111 

List of Figures vii 

List of Tables ix 

1. 

2. 

3. 

4. 

5. 

Introduction 

Materials 

Experimental 

Results and Discussion 

References 

List of Acronyms and Abbreviations 21 

Distribution List 23 

Report Documentation Page 29 

1 

1 

2 

3 

17 

V 
. 

vi 
List of Figures 

Figure 1. LPN resin constituents. .................................................................................. .2 

Figure 2. Storage modulus plots for thermal-cured IPNs.. ....................................... .4 

Figure 3. Storage modulus plots for e-beam-cured IPNs. ........................................ -4 

Figure 4. Loss modulus plots for thermal-cured IPNs.. ............................................ .5 

Figure 5. Loss modulus plots for e-beam-cured IF’Ns.. ............................................. .5 

Figure 6. Loss tangent plots for thermal-cured IPNs . ................................................ 6 

Figure 7. Loss tangent plots for e-beam-cured IPNs. ................................................ .6 

Figure 8. Loss modulus frequency sweep measurements of e-beam-cured 

1-3 sample.. ................................................................................................................ .9 

Figure 9. Loss tangent frequency sweep measurements of e-beam-cured 

l-3 sample.. .............................................................................................................. .10 

Figure 10. Loss modulus frequency sweep measurements of thermal-cured 

l-3 sample ................................................................................................................. 11 

Figure 11. Loss tangent frequency sweep measurements of thermal-cured 

l-3 sample ................................................................................................................. 11 

Figure 12. Loss modulus E” master curves and KWW fits for therrnal-cured 

samples ...................................................................................................................... 13 

Figure 13. E” master curves and KWW fits for e-beam-cured samples.. ..............14 

Figure 14. Shift factor plots for thermal-cured IPNs . ............................................... 15 

Figure 15. Shift factor plots for e-beam-cured Il?Ns . ................................................ 16 

Vii 
. . . 

VU3 
. 

List of Tables 

Table 1. Viscoelastic parameters derived from fits of master curves . . . . . . . . . . . . . . . . . . . . . 15 

. 

ix 
X 
1. Introduction 

Polymer composites and adhesively bonded hybrid composite structures (e.g., 

composite-ceramic integral armor) are promising candidates for ultralightweight structural armor for Future Combat Systems (FCS). The 

development of composite resins for low-cost processing with improved 

properties wiLl speed insertion of lightweight materials into FCS platforms. The 

processing and performance shortfalls in suitable resins may be addressed 

through exploitation of new electron beam (e-beam) curable systems including 

recently developed interpenetrating polymer network (IPN) materials. IPNs are 

polymers that are formed from the independent polymerization of two or more 

distinct networks, which results in unique molecular and physical 

microstructures offering improved properties [l]. The IPN materials are 

demonstrated to possess unique properties that can be exploited for 

multicomponent composite armor systems, including toughness and thermal 

stability. E-beam curing of composites and adhesives offers advantages, such as 

reduced cure shrinkages, over traditional autoclave processing by curing 

multiple resins through the thickness of thick-section composites in a single step. 

Because of the complexity and thickness of composite integral armor structures, 

e-beam curing is an attractive processing method for FCS applications and the 

development of these new resins permit this approach. This research seeks to 

discern the differences in viscoelastic behavior between traditional thermal and 

e-beam curing in IF’N resins, as these viscoelastic characteristics may prove 

important for composite integral armor applications. 

2. Materials 

Dalal and Palmese previously developed the IPN formulation and cure routes 

[Z]. The IPN formulation consisted of an epoxy network-diglycidyl ether of 

bisphenol F (DGEBF) (EEW = 170 g/mol) cured with bis(p-aminocyclohexyl) 

methane (EEW = 52.5 g/mol) - coupled to a free radical methacrylate network 

(lb-hexanediol dimethacrylate [HDDMA]). The HDDMA free radical 

crosslinking reaction mechanism was initiated either via e-beam or thermal using 

the organic peroxide initiator 2,5-dimethyl-2,5-di-(t-butylperoxy) hexane. To 

couple the epoxy and methacryl networks together, a partial esterification 

reaction was used to convert an epoxy functionality of triphenylohnethane 

triglycidyl ether (TMTE) (EEW = 162 g/mol) to a methacryl functionality 

following the procedure outlined previously [2]. This 33% methacrylated TMTE, 

along with the other constituents, is illustrated in Figure 1. 
diglycidyl ether of bisphenol F (DGEBF) Epoxy 

1,6-hexanediol dimethacrylate (HDDMA) 

ci 

: 2 -0 --r;-C;H2 

.H2” 2 CH3 

33% methacryiated triphenytolmethane triglycidyt ether (33% TOTE) 

Figure 1. IPN resin constituents. 

3. Experimental 

To study the influence of coupling between the DGEBF and HDDMA networks 

by the 33% TMTE the molar ratio of DGEBF to 33% TMTE was systematically 

varied from 4 mols of DGEBF to 0 mols of 33% TMTE (4-O samples) through 

2 
0 mols of DGEBF to 4 mols of 33% TMTE (O-4 samples). Intermediate molar 

ratios of DGEBF:33% TMTE (3:1,2:2, and 1:3) were also examined. The PACM 

concentration was always maintained at stoichiometry with respect to the 

DGEBF and 33% TMTE epoxy resins. The weight fraction of DGEBF and TMTE 

to HDDMA was maintained at a constant ratio of 0.70 to 0.30. 

Two routes of curing the IPNs were pursued. The thermal cure conditions 

consisted of an initial cure at 70 “C for 3 hr, followed by cure at 140 “C for 2 hr, 

and a final post-cure at 250 “C for 1 hr. The two-stage curing schedule was 

selected to strictly control the formation of the epoxy-amine ternplated network. 

Previous research has shown that the Michael’s addition side-reaction between 

the methacrylates and the PACM is minimal at 70 “C [2]. Therefore, the epoxyamine stoichiometry is maintained. The e-beam cure route consisted of initial 

cure at 70 “C for 3 hr, followed by e-beam cure at a radiation dose of 20 Mrad, 

and a final post-cure at 250 “C for 1 hr. 

Dynamic mechanical analysis (DMA) was performed using a TA Instruments 

2980 DMA in the dual cantilever-bending mode. The samples were tested using 

a 20-mm frame and oscillatory displacement amplitude of 7.5 p. The 

displacement amplitude was verified to ensure linear viscoelastic response by 

deriving corresponding stress-strain curves. The typical sample’s width and 

thickness were approximately 12 mm x 2 mm, respectively. Constant heating 

rate experiments were carried out at 2.0 OC/min, with a frequency of 1 Hz from 

50” to 275 “C. Master curves were constructed from data obtained from multiple 

frequency sweeps, which were measured over three decades of frequency, 

ranging from 0.1 to 30 Hz in 3 “C isothermal steps. The glass transition 

temperatures were taken as the peak maximum of the loss modulus (E”) curves 

measured at 1 Hz. All samples were heated twice in the DMA. After the first 

heat, the samples were allowed to slow cool in the DMA to provide matching 

thermal histories. All DMA results in this report were taken from second heat 

measurements. 

4. Results and Discussion 

Figures 2-7 illustrate the dynamic mechanical response of the thermal and 

e-beam-cured samples at constant frequency (1 Hz) as a function of temperature. 

The overall trends in the dynamic mechanical properties are similar between the 

thermal and e-beam curing routes. In both cases, as the molar ratio of 33% TMTE 

to DGEBF is increased, the Tg is elevated and the glass to rubber transition zone 

is broadened. The trifunctional TMTE increases the crosslink density (pc) of the 

cured IPN. An increase in the width of the distribution of segmental relaxation 

times is consistent with an increased oc in highly crosslinked epoxy networks [3]. 
2.5 - 

2.0 - 

1.5 - 

1.0 - 

0.5 - 

0.0 - 

40 60 80 100 120 140 160 180 200 220 240 260 

Temperature (“C) 

Figure 2. Storage modulus plots for thermal-cured IPNs. 

2.5- 

2.0 - 

1.5 - 

G5 l.Oill 

0.5 - 

0.0 - 

-0-4 

-1-3 

-2-2 

--3-l 

-4-o 

~,.,,,.,.,,,,,.,,,,,.,,,,, 

40 60 80 100 120 140 160 180 200 220 240 260 280 

Temperature (“C) 

Figure 3. Storage modulus plots for e-beam-cured PNs. 

4 
160 , 

BO60- 

-1-z 

- 2 -2 

-3-l 

-4-G 

q %mJD ~~00000 

40 60 80 100 120 140 160 180 200 220 240 260 

Temperature (“C) 

Figure 4. Loss modulus plots for thermal-cured Il?Ns. 

300 , 

250- 

200- 

;;; 150- 

% 

Ll 

ioo50- 

o- 

--0-4 

-1-Z 

-2-2 

-3-I 

-4-a 

i , ., ., ., ., , , , , ’ , ’ , ’ , , , , , , , 

40 60 80 100 120 140 160 180 200 220 240 260 280 

Temperature (“C) 

Figure 5. Loss modulus plots for e-beam-cured Il?Ns. 

5 
0.6 

-a- 0 -1 

-I-: 

1 2 -; 

1 

60 80 100 120 140 160 180 200 220 240 260 

Temperature (“C) 

Figure 6. Loss tangent plots for thermal-cured IPNs. 

0.9 - 

0.8 - 

0.7 - 

0.6 - 

0.5- 

a 0.4 - 

ii 0.3 - 

0.2 - 

0.1 - 

0.0 - 

-- O- 

-l- 

-2- 

-33 

-44 

1 

0 

-0.1 ! 

, 

, 

, 

, 

, 

) 

, , , ) , , I',',' I 'I 1, 

40 60 80 100 120 140 160 180 200 220 240 260 280 

Temperature (“C) 

Figure 7. Loss tangent plots for e-beam-cured LPNs. 

6 
The 33% TMTE also covalently couples the epoxy-amine and free radical 

methacryl networks together, which increases the heterogeneity of molecular 

environments. This increase in T, and breadth of the glass to rubber transition 

with increased coupling between the individual components of an IPN has been 

observed in other systems [4-7’j. Akay and Rollins [4] determined that the 

broadening of the glass transition cannot be explained by a simple increase in 

crosslink density alone, and the coupling between the networks results in more 

complicated viscoelastic interactions. An apparent increase in crosslink density 

can be noticed in the rubbery regions of the storage modulus plots presented in 

Figures 2 and 3. 

While the overall trends in the dynamic mechanical properties between the 

thermal- and e-beam-cured samples are similar, sufficient differences are present 

that merit further investigation. The e-beam-cured samples exhibit a steady rise 

in Ts from 122” to 190 “C, as seen in the E” and tan 6 plots, as the molar ratio of 

the TIUTE coupler is extended through the 4-O to O-4 samples, respectively. The 

Ts for the 4-O (no coupler) thermal-cured samples is nearly identical to the 

e-beam counterpart at 123 “C, but the Ts values for the thermal route plateau at 

145 “C beginning with the 2-2 sample. Fourier transform infrared (FTIR) results 

(not shown) indicate that the e-beam cure method achieves a greater degree of 

methacryl conversion than the thermal route, which results in the increased 

crosslink density and Ts for the e-beam-cured samples [2]. The depressed 

thermal conversion of methacryl functional groups is presumably due to the free 

radical propagation becoming restricted because of diminished molecular 

mobility as the thermal cure is advanced. It is interesting to note that the 

ultimate Ts reached during thermal cure of 145 “C closely matches the median 

thermal cure step of 140 “C. The thermal post-cure step at 250 “C increases the 

degree of methacryl conversion from 75-79% to 84-89%, but does not yield a 

further increase Ts. The e-beam-cured samples have a higher degree of 

methacryl conversion prior to post-curin g (91-94%) and the additional post-cure 

step only increases the methacryl conversion by 2-4%. 

The differences in glassy behavior between the thermal- and e-beam-cured IPNs 

are readily apparent as seen in Figures 2 and 3, respectively. The glassy storage 

moduli of the thermal-cured IPN vary greatly and rise from 1 GPa for the 4-O 

sample to 2.5 GPa for the O-4 sample. The glassy storage modulus of the e-beamcured samples lies near 2.5 GPa for all of the samples regardless of the DGEBF to 

TMTE molar ratio. These results were also reproduced using another DMA [2], 

which adds weight to the argument that the differences in glassy response 

between the curing routes are real, and not instrument artifacts. Mobility in the 

glassy state is related to sub-T, local secondary relaxation mechanisms such as 

the B transition [8]. The secondary B transition in epoxies based upon common 

diglycidylether of bisphenol A (DGEBA) is centered near -60 “C and is caused by 

7 
motion of the hydroxypropylether linkage and phenol ring flips [9]. The 

amplitude of the damping peak caused by the B transition has been found to be 

inversely proportional to the crosslink density at ambient temperatures in 

traditional microstructure networks, which results in a decreased glassy 

modulus [lo-131. A possible explanation proposed for this effect is that a 

crosslink point isolates short segments of network chain, therefore uncoupling 

local motion from other neighboring chain segments [ll]. It has also been 

postulated that decreasing the cross-link density will decrease the glassy 

modulus due to weaker Van der Waals interactions [14]. The results obtained for 

this research point towards the uncommon trend in that the thermal-cured IPNs 

with the lowest pc (4-O sample with 34% methacryl conversion) have the lowest 

magnitudes of glassy modulus. The e-beam-cured IPNs ail have similar 

methacryl conversions of 95-97% and similar glassy moduli. The thermal- and 

e-beam-cured IPNs have identical thermal histories after the post-cure step, so it 

is difficult to gauge the cause of the modulus differences (free volume, degree of 

conversion, IPN microstructure, etc.) without performing a detailed analysis of 

the p transition. On the basis of differences in glassy modulus between the 

e-beam and thermal-cured IPNs, one could also expect interesting differences in 

the sub-Ts nonequilibrium behavior [15]. This could lead to variations in the 

physical-aging rates between the e-beam and thermal-cured IPNs. 

In addition to the differences in glassy modulus between the thermal- and 

e-beam-cured samples, significant changes in the tan 6 peak in the glass 

transition region are also apparent. Figures 6 and 7 illustrate the tan 6 plots as a 

function of DGEBF to TMTE molar ratio for the thermal- and e-beam-cured 

samples, respectively. The e-beam-cured sample displays a continuous shift in 

T, towards higher temperatures with a corresponding broadening of peak width 

as the amount of TMTE is increased. The tan 6 curves for the e-beam samples are 

also relatively symmetric and are absent of any shoulders. As mentioned 

previously, the Ts of the thermal-cured O-4, l-3, and 2-2 samples plateau near 

145 “C due to diffusion considerations. Shoulders in the tan 6 signals of the 

thermal-cured l-3 and O-4 samples also appear that are not evident in the 

corresponding e-beam samples. The observation of a shoulder or second peak in 

the DMA spectra of an IPN is caused by multiple molecular relaxation 

mechanisms of multiple phase domains [5,16]. 

However, it is difficult to ascertain the molecular origin of the relaxation 

mechanism responsible for the tan 6 shoulders present in the thermal-cured l-3 

and O-4 samples from the simple isochronal DMA measurements performed for 

Figure 6. An alternative approach to enhance the shoulder region of the tan 6 

peaks is to sweep the frequency v) of the DMA measurement at isothermal 

temperatures. This method has. been used successfully to separate the 

overlapping B-transition of an epoxy matrix from the low temperature 

cL-transition of an included rubber toughener [17]. The peak-shoulder separation 

8 
achieved during this type of frequency sweep is improved for overlapping 

molecular processes that differ greatly in Arrhenius activation energy (EJ, as 

described by equation (1): 

where 

T = the temperature (in degrees Kelvin) of the peak maximum, and 

To = the reference temperature. 

Figures 8 and 9 show the isothermal frequency sweep E” and tan 6 plots for the 

e-beam-cured 1-3 sample, respectively. Again the e-beam-cured l-3 sample 

displays a single homogeneous transition with no shoulders. The E” peaks range 

in temperature from approximately 170” to 190 “C as the frequency of the 

measurement is swept from 0.1 to 30 Hz, respectively. The corresponding tan 6 

signal peaks are roughly 20 “C greater than the loss modulus peaks. 

8.2 - 

8.0 - 

7.8 - 

-iii7.6- 

a 

- 7.4- 

Ll 

o) 7.2- 

2 

7.0 - 

6.8 - 

6.6 - 

q 0.1 Hz 

o 0.3 Hz 

A IHt 

v 3Hz 

o IOHz 

+ 30 Hz 

-$A++ 

OAVO++ q %oo+ 

q oAv 

q ',AVo_ 

sb IA0 IA0 2;)o 2iio 

Temperature (“C) 

Figure 8. Loss modulus frequency sweep measurements of e-beam-cured 1-3 sample. 
q 0.1 Hz &!+-I 

o 0.3 Hz 

A IHz 

v 3Hz 

o IOHz 

+ 30 Hz 

Temperature (“C) 

Figure 9. Loss tangent frequency sweep measurements of e-beam-cured 1-3 sample. 

In contrast to the e-beam-cured 1-3 sample, E” shoulders are still evident in the 

multiple frequency-sweep plots of the thermal-cured 1-3 sample (Figure 10) at 

high temperatures trailing the peak maxims (-190”-210 “C), but are not shifted 

from the main loss modulus peaks at any measured frequency. Therefore, the 

molecular mechanism responsible for the loss modulus shoulder in the thermal 

case must have an Arrhenius activation energy nearly identical to the primarily 

observed transition. Furthermore, when the loss modulus curves of Figure 10 

and the tan 6 curves of Figure 11 are directly compared, an interesting 

observation is noted. The prnnary transition present in the E” curves (-135”-155 “C) 

correlate to the leading shoulders in the respective tan 6 curves. The trailing 

shoulders in the E” curves (-190”-210 “C) link to the primary peaks in the tan 6 

curves. The difference in temperature between analogous E” and tan 6 peaks for 

the thermal-cured l-3 samples is nearly 40 “C. This raises more suspicion that 

the observed relaxations near 140”-145 “C for the higher crosslink density 

thermal samples are due to memory retention of the median oven-cure stage. 

Cure memory in the dynamic mechanical glass transition has been observed in 

polyamidoamine-epoxy networks by other researchers [18]. Sanz et al. [19] also 

observed shoulders in the tan 6 signals of stoichiometric epoxy-amine systems 

with temperatures and magnitudes dependent upon the thermal-cure conditions. 

This is a result of a fraction of the network structure attaining a high-crosslink 

density during gelation at the median-cure stage. During the post-cure stage, 

10 
8.2 

8.0 - 

7.8 - 

7.6- 

7.4 - 

7.2 - 

7.0 - 

6.8 - 

6.6 - 

j '% 

00 ++ 

00 ++ 

k %T+ 

--o-O.1 Hz 

-o- 0.3 Hz 

-A-l Hz 

-v- 3 Hz 

&lOHz 

-+-30 Hz 

b 

3 

6.4 ! I I I I I 

50 100 150 200 250 

Temperature (“C) 

q 0.1 Hz 

o 0.3 Hz +f@-% 

A IHz 4%% ++ 

v 3Hz 

o IOHz 

+ 30 Hz 6 

Figure 10. Loss modulus frequency sweep measurements of thermal-cured 1-3 sample. 

0.25 

100 150 200 

Temperature (“C) 

250 

Figure 11. Loss tangent frequency sweep measurements of thermal-cured l-3 sample. 

11 
further crosslinking reactions occur, but are diffusion limited within these 

regions of high-crosslink density from the previous cure stage. The physical 

heterogeneity between these regions of low- and high-crosslink density results in 

the broadness observed in the E” curves during DMA experiments. The 

broadness in the E” curves for the thermal-cured IPNs appears to be caused by 

physical mechanisms. In the case of the thermal-cured 1-3 samples, the E, 

calculated from the primary loss modulus peaks (135”-155 “C range) is 477 kJ/mol 

and 456 kJ/mol from the tan 6 peaks (175”-195 “C range). These values are very 

close to E, values of 498 kJ/mol and 509 kJ/mol calculated for the e-beam l-3 E” 

and tan 6 signals, respectively. 

A goal of this research was to investigate the possibility of using DMA to 

correlate the observed densities and cure shrinkage of the IPN adhesives to free 

volume. Gerard et al. [20] have used this approach to relate the crosslink density 

and chain flexibility of model epoxy-amine networks to free volume. Akay and 

Rollins [4] also examined the free-volume and glass-transition broadening of 

polyurethane/poly (methyl methacrylate) IPNs from viscoelastic standpoint. For 

this research, the viscoelastic interpretations were derived from a similar 

perspective using master curves of the loss modulus, which were constructed in 

the frequency (0) domain in the glass transition region (Ts + 30 “C) following the 

time-temperature superposition principle (tTsp) [3]. Prior to shifting the 

modulus isotherms horizontally along the frequency axis, a vertical shift factor 

correction of T,/T was applied, where To represents the reference temperature. 

The horizontal shift factors (aT > Tg) were then fitted to the empirical and free 

volume derived Williams-Landel/Ferry (WLF) equation (2). The WLF equation 

is used to describe the temperature dependence of the distribution of relaxation 

times (2) in viscoelastic materials. The Ci and C2 constants of the WLF equation 

were determined using a nonlinear least squares fit. The fractional free volume 

at the glass transition (fs) and coefficient of thermal expansion of the fractional 

free volume (o1f) were then solved using the Cl and C2 constants and setting 

empirical Doolittle constant (B) equal to 1[20,21]. The Cl and C2 constants of the 

WLF equation can also be used to solve for E,. The breadth of the distribution of 

relaxation times was quantified by fitting the master curves to the KohlrauschWilliams/Watts (KWW) equation (5) and determining the coupling parameters 

69 PI- 7x62 P P ammeters were found using the methodology outlined by Weiss 

et al. [23,24]. 

12 

loga, = logf = 

-cd-T,) 

1; c, +&To) - 

c, 2- f 

2.3fg 

; c,=L. 

a/ (3) 
E, = -2.303RT2 d loga, 

dT ’ (4) 

&)=exp[-($]+o <p51. (5) 

Figures 12 and 13 show the E” master curves and KWW fits for the thermal and 

e-beam-cured samples, respectively. For each sample, the isotherms successfully 

shifted to form smooth and continuous master curves. The KWW equation also 

provided reasonable descriptions of the master curves, with the p parameters 

summarized in Table 1. For each cure route, it can be seen that the breadth of the 

E” master curves with increased Th4TE concentration. This increase in coupling 

between networks and subsequent increase in crosshnk density is reflected as a 

decrease in the p parameters (0.237 for the e-beam 4-O sample to 0.113 for the 

e-beam O-4 sample and 0.202 for the thermal 4-O sample to 0.113 to the O-4 

thermal sample). The reference relaxation time used for the KWW fits is also 

provided. 

25 

-6 -4 -2 0 2 4 6 6 

log coaT 

Figure 12. Loss modulus E” master curves and KWW fits for thermal-cured samples. 

13 
J 

250- 

225- 

200- 

xi- 175- 

g 150- 

f‘ 125- 

. 

c loom h 75- 

50- 

25- 

-6 -4 -2 0 2 4 6 8 

log oa, 

Figure 13. E” master curves and KWW fits for e-beam-cured samples. 

However, for tTsp to be gauged a success, the shift factor plots and WLF 

parameters must also be considered. Table 1 also summarizes TP E, CI, Cs fs, 

and of. Figures 14 and 15 show the shift factor plots for the thermal and e-beamcured samples, respectively. The values calculated for the thermal-cured O-4, 

1-3, and 2-2 samples must certainly be incorrect as multiple relaxation processes 

were clearly observed in the E” and tan 6 signals [3]. This can also be seen in the 

shift factor plots (Figure 14), as the curves for these samples were nearly linear 

and deviated from WLF behavior. The thermal-cured 2-2 sample fits the WLF 

equation so poorly that none of the parameters could be reasonably determined 

or approximated. The WLF parameters obtained for the thermal-cured samples 

with little inter-network coupling, 4-O and 3-1, seemed reasonable. 

The most interesting viscoelastic behavior is observed in the e-beam-cured PN 

where no directly observable multiple relaxation peaks or shoulders were found 

in the E” and tan 6 signals. The shift factor plots for the e-beam-curing process 

(Figure 15) are similar to the thermal-cured counterparts in that as the molar ratio 

of TMTE coupler to DGEBF is increased, the temperature sensitivity of log ar 

decreases. This result is consistent with the results of Ogata et al. [25], where a 

decrease in the temperature and frequency dependency of the glass to rubber 

relaxation for phenol-formaldehyde novolac epoxy networks was found as the 

crosslink density was increased. Significant departures from typical epoxy-amine 

network properties are found in the values of f, and cxf calculated for the e-beamcured IPNs. The uncoupled 4-O sample has values of fs and af that are in the range 

14 
Table 1. Viscoelastic parameters derived from fits of master curves. 

Sample p x101 rx102 T, E&2 

(“c) @J/mol) @J~~ol) 

Cl 

(2) 

fgx1W cqx104 

(s) (“c -1) 

0.38 190 

0.48 181 

0.75 158 

1.4 134 

2.9 122 

E-beam 

472.2 484.4 

497.0 623.6 

627.5 642.3 

597.7 no.2 

631.9 902.8 

47.1 399.1 0.92 

19.8 125.3 2.2 

17.1 94.7 2.5 

13.3 59.4 3.3 

11.0 36.4 4.0 

2.3 

1.8 

2.7 

5.5 

10.6 

1 POT 1 2.02 1.1 1 123 1 623.1 1 699.8 1 14.8 1 63.5 1 2.9 1 4.6 1 

Thermal 

0.35 I 145 I 440.4 I 443.9 I 22.9 I 172.6 1.9 I 1.1 

1.4 1 141 1 477.3 1 614.0 1 11.9 1 63.6 1 3.7 1 5.7 1 

0.5 145 501.4 c c c c c 

2.5 135 659.6 709.4 14.8 66.5 2.9 4.4 

aArrhenius. 

%KF. 

‘Multiple imnsitions observed in loss modulus spectra. 

6- 0 o-4 

0 l-3 

* 2-2 

v 3-l 

0 4-o 

-S! , , , , , , , , , , , , , 

-30 -20 -10 0 10 20 30 

T - T, 

Figure 14. Shift factor plots for thermal-cured IPNs. 

15 
6 

-t 

A 2-2 

-6. , , , , , , , , , , , , , 

-30 -20 -10 0 10 20 30 

T - T, 

Figure 15. Shift factor plots for e-beam-cured IPNs. 

of values for conventional thermosets, as determined by Gerard et al. [20]. As 

the molar ratio of TMTE to DGEBF is increased, the calculated values offs and of 

decrease significantly. The most coupled sample (O-4) has an af value in the 

rubbery region, which is an order of magnitude lower than the typical values 

reported by Gerard. This inter-network coupling profoundly increases the 

overall crosslink density of the IPN. One must also keep in mind that the 

viscoelastic response of the highly coupled e-beam IPN samples also deviates 

considerably from ideal WLF behavior, so the assumption of a single relaxation 

mechanism may not be valid. While it is uncertain whether or not multiple 

relaxation mechanisms from a heterogeneously crosslinked network structure 

are active in the case of the e-beam-cured IPNs, the trends in the viscoelastic data 

point towards a more homogeneously crosslinked network in comparison to the 

thermal-cure counterpart. 

16 
5. References 

1. Sperling, L. H. Interpenetrating Polymer Networks and Related Materials. 

New York: Plenum Publishing, 1981. 

2. Dalal, U. P., and G. R Palmese. “Formation of In-Situ Sequential 

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4. Akay, M., and S. N. Rollins. “Transition Broadening and WLF Relationship 

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and Bimaleimide: Properties of the Neat Resin and Composites.” JournaZ of 

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7. Han, J. L., and K. Y. Li. “Graft Interpenetrating Polymer Networks of 

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8. Fitz, B., S. Andjelic, and J. Mijovic. “Reorientational Dynamics and 

Intermolecular Cooperativity of Reactive Polymers. 1. Model Epoxy-Amine 

Systems.” Macromolecules, vol. 30, pp. 5227-5238,1999. 

9. Bershtein, V. A., N. N. Peschanskaya, J. L. Halary, and L. Monnerie. “The 

Sub-Tg Relaxations in Pure and Antiplasticized Model Epoxy Networks as 

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10. Cukierman, S., J. L. Halary, and L. Monnerie. “Dynamic Mechanical 

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17 
12. Georjon, O., G. Schwach, J. F. Gerard, and J. Galy. “Molecular Mobility in 

Polycyanurate Networks Investigated by Viscoelastic Measurements and 

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H. Sautereau. “Influence of Chain Flexibility and Crosslink Density on 

Mechanical Properties of Epoxy/Amine Networks.” Polymer Engineering and 

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2997,1996. 

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Elastomer-Modified Epoxy-Methacrylate Sequential Interpenetrating 

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Dynamic Mechanical Behavior of DGEBA/DDM Stoichiometric Epoxy 

Systems on the Conditions of Curing Process.” Journal of Appkd Polymer 

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18 
24. Weiss, G. H., M. Dishon, A. M. Long, J. T. Bendler, A. A. Jones, P. T. 

Inglefield, and A. Bandis. “Improved Computational Methods for the 

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Properties of Phenol-Formaldehyde Novolac Cured Epoxy Resins.” Journal of 

Applied Polymer Science, vol. 48, pp. 583-601,1993. 

19 
INTENTIONALLYLEFTBLANK. 

20 
List of Acronyms and Abbreviations 

aT 

B 

B parameter 

j3 transition 

“C 

Cl 

c2 

DGEBA 

DGEBF 

DMA 

E 

E” 

E, 

e-beam 

EEW 

f!z 

FTIR 

GPa 

HDDMA 

IPN 

mm 

MPa 

PACM 

T 

Tg 

TO 

tan6 

WLF isothermal shift factor 

Doolittle equation parameter 

KWW equation fitting parameter 

localized glassy polymeric sub-T, relaxation 

degrees Celcius 

WLF fitting parameter 

WLF fitting parameter 

diglycidyl ether of bisphenol A 

diglycidyl ether of bisphenol F epoxy resin 

dynamic mechanical analysis 

dynamic storage modulus 

dynamic loss modulus 

Arrhenius activation energy 

electron beam 

epoxy equivalent weight 

fractional free volume at T = Tg 

Fourier transform infrared spectroscopy 

giga Pascals 

1,6-hexanediol dimethacrylate 

interpenetrating polymer network 

Kohlrausch-William/ Watts equation 

millimeters 

mega Pascals 

bis (p-aminocyclohexyl) methane 

temperature 

glass transition temperature 

WLF equation reference temperature 

E”/E 

21 
tTsp 

TMTE 

WLF 

4-O 

3-1 

2-2 

l-3 

O-4 

af 

0 

PC 

T 

time-temperature superposition principle 

triphenylolmethane triglycidyl ether 

Williams-Landel/Ferry equation 

4 mols of DGEBF epoxy to 0 mols of 33% TMTE 

3 mols of DGEBF epoxy to 1 mols of 33% TMTE 

2 mols of DGEBF epoxy to 2 mols of 33% TMTE 

1 mols of DGEBF epoxy to 3 mols of 33% TMTE 

0 mols of DGEBF epoxy to 4 mols of 33% TMTE 

coefficient of thermal expansion of the fractional free volume 

micrometers 

angular frequency 

crosslink density 

relaxation time constant of KWW equation 

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28 
Analysis of E-Beam and Thermally Curable LPN Thermosets 

Robert E. Jensen, Giuseppe R. Palmese,* and Steven H. M&night 

U.S. Army Research Laboratory REPORT NUMBER 

A’I-I’Nz AMSRL-WM-MA ARL-‘l-R-2698 

Aberdeen Proving Ground, MD 210055069 

9. SPONSORlNGlMONlTORlNG AGENCY NAMES(S) AND ADDRESS 10.SPONSORING/MONITORlNG 

AGENCY REPORT NUMBER 

11. SUPPLEMENTARY NOTES 

*Drexel University, Philadelphia, PA 19104 

12a. DlSTRlBUTlON/AVAlLABlLlTY STATEMENT 12b. DISTRIBUTION CODE 

Approved for public release; distribution is unlimited. 

13. ABSTRACT(hfaxlmum200 words) 

Interpenetrating polymer network (LPN) materials possess unique properties that can be exploited for multicomponenl 

composite armor systems, including toughness and thermal stability. E-beam curing of composites and adhesives offers 

advantages, such as reduced cure shrinkages, over traditional autoclave processing by curing multiple resins through the 

thickness for thick-section composites in a single step. Because of the complexity and thickness of composite integral 

armor structures, e-beam curing is an attractive processing method for Future Combat System applications, and the 

development of these new resins permits this approach. This research seeks to discern the differences in viscoelastic 

behavior between traditional thermal and e-beam curing in LPN resins, as these viscoelastic characteristics may prove 

important for composite integral armor applications. 

14. SUBJECT TERMS 15. NUMBER OF PAGES 

IPN, DMA, viscoelasticity, epoxy, methacryl 34 

16. PRICE CODE 

17. SECURITY CLASSIFICATION 16. SECURITY CLASSIFICATION 

OF REPORT 

19. SECURITY CLASSIFICATION 20. LIMITATION OF ABSTRACT 

OF THIS PAGE OF ABSTRACT 

UNCLASSIFIED UNCLASSIFIED UNCLASSIFIED UL 

NSN 7640-01-280-5500 

29 Standard Form 298 (Rev. 2-89) 

Prescribed by ANSI Std. 23418 298-l 02 
l 

INTENTIONALLYLEFTBLANK. 

30