Materials Science & Engineering A 570 (2013) 164–171

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Materials Science & Engineering A

0921-50

http://d

n Corr

E-m

journal homepage: www.elsevier.com/locate/msea

Mechanistic aspects of impact initiated reactions in explosively consolidatedmetalþaluminum powder mixtures

B.B. Aydelotte n, N.N. Thadhani

School of Materials Science and Engineering, Georgia Institute of Technology, 771 Ferst Dr. NW, J. Erskine Love Building, Atlanta, GA 30332, USA

a r t i c l e i n f o

Article history:

Received 16 October 2012

Received in revised form

17 January 2013

Accepted 22 January 2013Available online 29 January 2013

Keywords:

Aluminum

Energetic material

Deformation

Impact

Mesoscale

93/$ - see front matter & 2013 Elsevier B.V. A

x.doi.org/10.1016/j.msea.2013.01.054

esponding author. Tel.: þ1 404 894 2888; fax

ail address: bradya@gatech.edu (B.B. Aydelott

a b s t r a c t

The mechanisms influencing impact initiation of reactions in structural energetic materials formed by

explosive compaction of pure Ni, Ta, or W powders mixed with Al powder are investigated in this work.

High speed images of rod-on-anvil impact tests indicate that the energy requirements for initiating

reaction in TaþAl compacts are lower than those in WþAl, while no reaction at all is observed in

NiþAl under the tested conditions. Mesoscale simulations performed on actual microstructures reveal

differences in the deformation behavior of the NiþAl, TaþAl, and WþAl compacts under similar

impact conditions. In the NiþAl and WþAl systems, the impact induced deformation is localized

primarily in the Al constituent, while significant plastic deformation of both constituents is observed in

TaþAl, making it a more reactive system.

& 2013 Elsevier B.V. All rights reserved.

1. Introduction

Structural energetic materials (SEMs) based on metal–polymercomposites [1], thermites [2], and intermetallic-forming mixtures [3–9] have been studied for many applications. SEMs are designed to behigh strength and inert under ordinary conditions. When subject toimpact or shock loading it is possible to initiate rapid, highlyexothermic reactions in these materials. The mechanisms for impactinitiation of such reactions are poorly understood.

It is well known that the plastic deformation of metalsdissipates some energy through formation of dislocations, twins,and other defects; the rest of the energy, often estimated to be90%, is dissipated through adiabatic heating of the material[10,11]. Several researchers [3,7,8] have noted that plastic defor-mation plays a role in the reaction initiation of reactive, powdermixtures. It is believed that deformation-induced, adiabatic heat-ing plays a role in reaction initiation in the material systemsconsidered in this work.

The structural energetic materials considered here are explosivelycompacted, approximately equivolumetric mixtures of NiþAl, WþAl,and TaþAl powders (hereafter NiþAl, TaþAl, or WþAl). Details ofthe compaction process are contained elsewhere [12,13]. Themechanical properties of these explosively compacted mixtures wererecently studied by Wei et al. [13]. The current work focuses on theinfluence of mesoscale microstructure features and constituent

ll rights reserved.

: þ1 404 894 9140.

e).

properties on reaction initiation behavior in these same explosivelycompacted NiþAl, WþAl, and TaþAl mixtures.

Rod-on-anvil impact tests were utilized to study impactinitiated reactions in these systems under conditions of uniaxialstress loading. Direct observation of the time resolved mesoscaledeformation and flow of the constituents is not possible. CTH [14],an Eulerian hydrocode, was used to conduct mesoscale simula-tions of rod-on-anvil impact using NiþAl, WþAl, and TaþAlmicrostructures at 350 m/s. The constitutive properties of the Ni,Ta, W, and Al phases and their interfaces were varied to study theplastic deformation response of the explosively compacted NiþAl,WþAl, and TaþAl, and determine their influence on reactioninitiation. The mesoscale simulations suggest that the extent ofplastic strain in both constituents of these materials stronglyinfluences reaction initiation. It was observed that TaþAl devel-ops more average plastic strain in Ta than either NiþAl or WþAldo in Ni or W respectively. This behavior, driven by the topologyof the harder phase, the mechanical properties of both phases, andthe higher interface strength of the Ta/Al interfaces, and the smallspacing between reactants in TaþAl is responsible for the lowerenergy input necessary to initiate the reaction in the TaþAlsystem. To the best of the authors’ knowledge, this is the firstwork showing the influence of topology and interface strength onthe reaction behavior of intermetallic forming SEM systems.

2. Experimental setup and results

Approximately equivolumetric mixtures of tantalum, tungsten,or nickel with aluminum powders were explosively compacted

Anvil Anvil

Ni+Al,W+Al,

orTa+Al

Cu

Post Impact

Pre Impact

Cu Rod Rod

Fig. 1. (a) A schematic of the rod-on-anvil impact test. (b) TaþAl impact at 500 m/s. The emission of bright light indicates reaction in TaþAl. Total elapsed time for the 16

frames is 7:5 ms.

Table 1Specific kinetic energy at impact corresponding to the lowest energy for reaction initiation, volume fraction,

surface area per unit volume, and mean free path between Ni, W, or Ta and Al in NiþAl, WþAl, and TaþAl

respectively.

NiþAl WþAl TaþAl

Threshold V No reactiona 502.0 m/sa 365.5 m/sa

Threshold KE No reactiona 8.24�106 J/kga 4.75�106 J/kga

Vv 0.5170.018 Ni 0.5470.02 W 0.6170.02 Ta

Sv 68.571.8 mm�1 56.071.8 mm�1 11772.7 mm�1

l¼ 4ð1�VvÞ=Sv 0.028770.001 mm 0.032570.002 mm 0.013370.0008 mm

a Data from Du et al. [12].

Ni

20 um

20 um

20 um

AI WTaAI

AI

Fig. 2. (a) NiþAl microstructure. (b) TaþAl microstructure. (c) WþAl microstructure.

B.B. Aydelotte, N.N. Thadhani / Materials Science & Engineering A 570 (2013) 164–171 165

to a nominal density of 97–99% theoretical maximum density [12].Small, cylindrical samples 3.070.1 mm in diameter and 270.1 mmthick were machined out of the resulting compacts of NiþAl, WþAl,and TaþAl and mounted on 7.62 mm diameter copper projectiles asshown in Fig. 1a. The projectiles were launched with a single stagegas gun at a hardened steel anvil at various velocities in a vacuumchamber evacuated to 50 mTorr to determine the impact velocity atwhich reaction takes place [12]. Cu projectile mass and vacuum levelswere held constant for all experiments. An IMACON 200, a high-speedgated CCD camera, was used to film the impact and capture transientdeformation. The velocity prior to impact was measured using a laserbeam interruption system. It has been demonstrated that the highlyexothermic reactions of metals [15,16], thermites [1,17,18], andintermetallic forming mixtures [1,19–21] produce brilliant lightwhich has been utilized to detect reaction. Fig. 1b shows an exampleof reaction recorded by the IMACON 200 during an experiment onTaþAl after impact at 500 m/s.

The reaction thresholds of NiþAl, WþAl, and TaþAl arepresented as minimum impact energy leading to reaction in

Table 1. It can be seen that TaþAl has a lower minimum specifickinetic energy (in joules per kilogram of sample mass) and impactvelocity for reaction initiation than WþAl, and NiþAl underwentno reaction at any tested velocity (vimpact r550 m=s).

Samples of NiþAl, WþAl, and TaþAl were sectioned,polished, and photographed to reveal their microstructures asseen in Fig. 2. About 25–50 images were taken of the micro-structures of each sample in a uniform random manner. Volumefraction, Vv, and surface area per unit volume, Sv, were measuredtwice on each image. Mean free path, l, between Ni, W, or Ta andAl, is calculated and tabulated along with Vv and Sv in Table 1with 95% confidence. For the computational studies of deforma-tion behavior discussed below, additional contiguous images ofeach material microstructure were combined into a compositeimage covering an approximately 1 mm square area.

Prior work on composite Ni/Al foils has shown that thereaction ignition energy is proportional to the bilayer spacingfor electrical [19,20], mechanical [19,20], or thermal [22,20]ignition. Mean free path, l (analogous to bilayer spacing),

B.B. Aydelotte, N.N. Thadhani / Materials Science & Engineering A 570 (2013) 164–171166

compares the spacing of reactants in NiþAl, WþAl, and TaþAl.Table 1 illustrates that the TaþAl has both the lowest mean freepath and the lowest minimum specific kinetic energy prior toreaction initiation. However, WþAl and NiþAl do not follow thetrend in this case which suggests that other factors, notablydeformation behavior, influence impact initiation of reaction.Deformation behavior is addressed in detail below.

3. Computational studies of deformation behavior

3.1. Computational setup

CTH simulations of rod-on-anvil impacts (schematically illu-strated in Fig. 1a) were performed in 2D plane strain on eachmicrostructure. The composite microstructure images of NiþAl,TaþAl, and WþAl were filtered and processed in MATLAB andthen converted to a CTH diatom. The simulation setup consists ofa 1 mm square microstructure, similar to, though larger thanthose shown in Fig. 2, carried by a copper projectile into a rigidanvil at 350 m/s. The whole assembly emulates the experimentalconfiguration, though reduced in size because of constraints oncomputational resources. The ratio of projectile to sample widthwas preserved.

The constitutive properties of the copper projectile weredescribed using the Johnson–Cook strength model [23]:

sy ¼ ½AþBEnp� 1þC ln

_E_E0

� �� �½1�Tnm

�:

Recall that _E ¼ffiffiffiffiffiffiffiffiffiffi_Eij _Eij

pand E ¼

R_E dt with repeated indices indicating

summation. A, B and C are fitting parameters. Tn is the homologoustemperature, n is the work hardening constant, E0 is a reference strainrate, and m is the thermal softening parameter. For Ni and Al, theSteinberg–Cochran–Guinan strength model [24]:

sy ¼ s0½1þbðEþEiÞ�n 1þ

s0ps0

P

Z1=3þ

G0TG0ðT�300Þ

� �

was used. Strain is defined in the same manner. Ei is the initial plasticstrain, Z¼ v0=v, P is the pressure, G is the shear modulus, b and n

are work hardening parameters, 0 denotes a derivative with respect tothe subscript, and T is the temperature. For Ta and W, the Zerilli–Armstrong BCC model [25]:

sy ¼Dsiþc2E1=2 expð�c3Tþc4T ln _EÞþkd�1=2

was utilized. c2 through c4 are fitting parameters, Dsi is thecomponent of strength from solute or initial dislocation density, k isthe Hall–Petch parameter, and d is the grain size. Model parameterswere taken from the CTH model database. Following Wei et al. [13],the yield strengths Al, Ni, Ta, and W were adjusted based on Tabor’sRule, sy ¼Hv=3.

The Johnson–Cook Fracture model was used for all the materi-als with values drawn from various sources [26,27]. It has theform

Ef ¼ ½D1þD2 expðD3snÞ� 1þD4 ln_E_E0

� �� �½1þD5Tn

�,

D¼XDEi

Ef:

When D¼1 the material is no longer able to support a shearstress. Ef is the failure strain, D1 through D5 are fitting parameters,sn is the hydrostatic stress, and the other parameters have thesame meaning as above.

The fracture strength of each material was taken to beequivalent to the spall strength and typical values from theliterature were used [28]. Interface strength between the nickel,

tantalum, or tungsten and aluminum was varied between 0 and400 MPa to explore the role of interface strength.

The Mie–Gruniesen equation of state was used to describe thepressure, volume, and temperature response for all of the con-stituents. Information about the implementation can be found byconsulting the CTH documentation [29]. No attempt was madehere to include a reaction model for the formation of intermetallicphases; the focus of this work is analysis of conditions prior toreaction rather than the formation of product phases. All of theconstants used for the equation of state and constitutive modelswere drawn from the CTH libraries.

A uniform mesh resolution of 2 mm by 2 mm was utilized. Themaximum mesh resolution was dictated by the available computa-tional power; mesh resolution studies indicate that the maximumstrain values are not fully converged, however the patterns andrelative amounts of strain described in the following sections arestable and consistent, differing only in the very highest value of strainpredicted. The mesh was linearly graded from a cell size of 2–20 mmoutside the area where the microstructure is deformed to conservecomputational resources. The region of the mesh initially containingthe copper rod was still more coarsely meshed with cells beinggraded up to 100 mm; this was also done to conserve computationalresources.

The TaþAl, NiþAl, and WþAl samples were initially simu-lated using the same microstructure (NiþAl) with zero interfacestrength. The material properties, equations of state, and consti-tutive relations were changed to correspond with NiþAl, WþAl,and TaþAl. This made it possible to separate the effects ofmaterial properties on the deformation behavior from those ofmicrostructure morphology. Simulations were then conducted onthe NiþAl, TaþAl and WþAl microstructures while varyinginterface strengths.

3.2. Computational results

Fig. 3 shows a series of contour plots of plastic strain, Ep, forsimulated impact on AlþAl, NiþAl, WþAl, and TaþAl using thesame NiþAl microstructure. The AlþAl data is included to showthe effects of the Ni, W, or Ta particles on the extent of strainlocalization in the surrounding aluminum matrix. It can be seenthat the deformation is localized into bands and enhanced by thepresence of Ni, Ta or W particles.

Plastic strain data like that in Fig. 3 is volume averaged and re-plotted as a function of time for the Al matrix and the Ni, Ta, andW particles separately in Fig. 4a, for simulations where the sameNiþAl microstructure was used for all materials. The AlþAlmaterial shows a much lower average plastic strain than Al inNiþAl, WþAl, or TaþAl. The higher densities and strengths of theNi, W, and Ta relative to Al cause additional deformation in the Almatrix; the W particles cause the highest relative deformation inAl and the Ni particles the least. Ta was found to plasticallydeform more than Ni and W particles.

Figs. 4b and 5a are the average strain vs. time histories forNiþAl and WþAl respectively; plastic strain in the Al, Ni, and Wis not very sensitive to interface strength. Both materials havesimilar morphologies, as shown in Fig. 2a,c with convex Ni or Wparticles embedded in a continuous Al matrix. Due to the greaterstrength of the W and Ni particles and the connected nature of theAl matrix, the Al in NiþAl and WþAl deforms preferentially inshear bands, allowing the Al to accommodate the majority of thedeformation imposed by the Cu rod during impact as shown inFig. 6a, b.

Fig. 5b shows the strain history of TaþAl, illustrating slightlygreater sensitivity to interface strength and considerably moredeformation in Ta than in the Ni or W. The TaþAl material alsodemonstrates less deformation within the Al component. The

0 1 2 3 40

0.5

1

1.5

2

2.5

ε pla

stic

ε pla

stic

Al+AlAl:Ni+AlAl:W+AlAl:Ta+Al

Ni:Ni+AlW:W+AlTa:Ta+Al

0 1 2 3 40

0.5

1

1.5

2

2.5

Al:0 PaAl:1 KPaAl:100 MPaAl:400 MPa

Ni:0 PaNi:1 KPaNi:100 MPaNi:400 MPa

Time (s) Time (s)

Fig. 4. (a) Ep vs. time for AlþAl, NiþAl, WþAl, and TaþAl. The Al matrix is plotted separately from the Ni, Ta, or W particles. (b) Average plastic strain vs. time for NiþAl

for different interface strengths. Al and Ni are plotted separately.

Strain

8

Al+Al

Ni+Al

W+Al

Ta+Al

10

4

2

6

8

0

Fig. 3. Contour plots of plastic strain in AlþAl, NiþAl, WþAl and TaþAl at 1:7 ms post-impact at 350 m/s. The impact direction is toward the top of the page. Note the

strain localization between the particles of Ni, W, or Ta in the respective microstructures.

0 1 2 3 40

0.5

1

1.5

2

2.5

3

ε pla

stic

ε pla

stic

Al:0 PaAl:1 KPaAl:1 MPaAl:100 MPaAl: 400 MPa

W:0 PaW:1 KPaW:1 MPaW:100 MPaW:400 MPa

0 1 2 3 40

0.5

1

1.5

2

2.5

3

Al:0 PaAl:1 KPaAl:1 MPaAl:100 MPaAl: 400 MPa

Ta:0 PaTa:1 KPaTa:1 MPaTa:100 MPaTa:400 MPa

Time (μs) Time (μs)

Fig. 5. (a) Ep vs. time for WþAl for different interface strengths. Al and W are plotted separately. (b) Ep vs. time for TaþAl for different interface strengths. Al and Ta are

plotted separately.

B.B. Aydelotte, N.N. Thadhani / Materials Science & Engineering A 570 (2013) 164–171 167

Ni

W

Ta

AlStrain

5

4

3

2

1

0

Fig. 6. (a) NiþAl microstructure 1:9 ms post-impact. (b) WþAl microstructure 1:9 ms post-impact. (c) TaþAl microstructure 1:9 ms post-impact. Ep in Al in each material is

depicted.

B.B. Aydelotte, N.N. Thadhani / Materials Science & Engineering A 570 (2013) 164–171168

morphology of this material, shown in Fig. 2b, is quite differentthan that of the NiþAl or WþAl. Ta forms a continuous phase andcontains Al islands within it in 2D. This will be addressed in detailbelow. Ta strains more as it is softer than either the W or the Niand work hardens less than Ni. Furthermore, Al cannot deformaround the Ta as easily as it does around the Ni or W as shown inFig. 6a,b. The Al regions link up and form shear bands duringdeformation, but those shear bands cannot propagate as far dueto the unconnected nature of Al in TaþAl as seen in Fig. 6c.This limits the deformation of the Al and requires Ta to accom-modate more deformation. In contrast, the localized deformationin Al circumvents Ni and W particles in NiþAl and WþAl.

The interface strength sensitivity of the Ta deformationappears to be a result of the different microstructure morphologyof TaþAl. With high interface strengths, the aluminum is con-strained to the Ni, W, or Ta. In NiþAl and WþAl, the constraintdoes not change the strain history much because shear bandsquickly form in the Al and most of the deformation is restricted toAl irrespective of interface strength. In TaþAl, shear stressesacross the Al/Ta interfaces lead to earlier and more extensiveformation of shear bands in the Al which then leads to greaterdeformation of the Ta. For lower strength interfaces, the Al/Tainterfaces fail more readily which accommodates the imposeddeformation: shear bands do not form as quickly nor are they asextensive. The result is that TaþAl has its highest average plasticstrain values for high interface strengths.

Quasistatic and dynamic strength data obtained by Weiet al. [13] and SEM images of fracture surfaces provide clues aboutthe interface strength in these material systems. Wei et al. [13]found that TaþAl had much higher ultimate strengths thanNiþAl or WþAl, although a rule of mixtures approach wouldpredict a different result. Using an FEM analysis they concludedthat the Al/Ni interfaces in NiþAl are of lower strength than the

constituents. They were unable to deduce an Al/Ta interfacestrength for TaþAl [13]. No analysis of interface strength wasperformed on WþAl.

In the present work, interface strength is found to influencethe plastic deformation behavior in TaþAl, hence it is importantto have some sense of the relative strength of the Al/Ta interfaces.Comparison of fracture surfaces from Hopkinson Bar compressionsamples of NiþAl and WþAl, Fig. 7a and b respectively, revealsextensive particle pull out and interfacial fracture, clear evidenceof weak interfaces between Al and the other component.In contrast, SEM images of a TaþAl fracture surface from aHopkinson Bar compression sample shown in Fig. 8a and bprovide evidence that fracture in these samples is intergranularrather than interfacial. This suggests that adhesion between Taand Al particles is much greater than adhesion between Ni and Alor W and Al. This higher level of adhesion promotes higher levelsof plastic deformation in TaþAl as shown in Fig. 5b.

If the same TaþAl microstructure with a 400 MPa interfacestrength is altered by assigning the Ta phase the Al equation ofstate and constitutive properties, and assigning the Al phase theTa equation of state and constitutive properties and the simula-tion is rerun, the results are very similar to those observed forNiþAl or WþAl. The amount of deformation in the Ta decreasesdrastically. The amount of deformation in the Al increases. Thelower strength of the Ta relative to the W also concentrates lessdeformation in the Al. These results shown graphically in Fig. 9aconfirm that the topology of the tantalum plays a significant rolein its deformation.

The same exercise of reversing the microstructures wasperformed on the WþAl and NiþAl microstructures, this timewith zero strength interfaces because the interface strength wasseen to be so low in these materials: the material properties werereversed such that W or Ni became the matrix phase and Al

Fig. 8. Fracture surface from TaþAl impacted in the Hopkinson Bar at 705� and 3840� do not show interfacial fracture.

0 1 2 3 40

0.5

1

1.5

2

2.5

ε pla

stic

ε pla

stic

Al/(Al Matrix)Ta/(Al Matrix)Al/(Ta Matrix)Ta/(Ta Matrix)

0

0.5

1

1.5

2

2.5Al/(Al Matrix)W/(Al Matrix)Al/(W Matrix)W/(W Matrix)

Time (μs)0 1 2 3 4

Time (μs)

Fig. 9. (a) Ep vs. time for TaþAl with 400 MPa interface strength for the experimentally derived and the reversed microstructures. Al and Ta are plotted separately.

(b) Ep vs. time for WþAl with zero interface strength for the experimentally derived and the reversed microstructures. Al and W are plotted separately.

Fig. 7. (a) Fracture surface from NiþAl impacted in the Hopkinson Bar showing interfacial fracture between the phases. (b) Fracture surface from WþAl impacted in the

Hopkinson Bar also showing interfacial fracture between the phases [30].

B.B. Aydelotte, N.N. Thadhani / Materials Science & Engineering A 570 (2013) 164–171 169

because the reinforcement phase. The results are very similar tothose seen in TaþAl as shown in Figs. 9b and 10. When W and Niare the matrix, there is substantially more deformation in thosephases. The reversed microstructure simulations show that hav-ing the stronger phase as the matrix increases the amount ofdeformation in that phase. This again shows the importance ofthe microstructure topology of these materials in influencingtheir deformation behavior. When a hard phase is made into thematrix, the topological connectivity is increased and that of thealuminum is decreased. Increasing the connectivity of a phaseleads to additional deformation in that phase. These results also

suggest that the deformation of these materials is influenced to agreater extent by the topology of constituents than by the inter-face density of the constituents.

3.3. Analysis of simulation data and its implication for reaction

response

Mesoscale simulations predict that TaþAl has the highestaverage plastic strain values in the non-Al component whencompared with NiþAl and WþAl. Energy dissipated in plasticwork will be dissipated principally as heat as discussed earlier;

0 1 2 3 40

0.5

1

1.5

2

2.5

ε pla

stic

Al/(Al Matrix)Ni/(Al Matrix)Al/(Ni Matrix)Ni/(Ni Matrix)

Time (μs)

Fig. 10. (a) Ep vs. time for NiþAl with 0 interface strength for the experimentally

derived and the reversed microstructures. Al and Ni are plotted separately.

B.B. Aydelotte, N.N. Thadhani / Materials Science & Engineering A 570 (2013) 164–171170

both phases need to undergo plastic deformation and associatedheating to react. The Ta in TaþAl plastically deforms more thanthe W in WþAl or the Ni in NiþAl and thus has a higherdeformation induced temperature increase. Having extensivedeformation in both phases causes the TaþAl to be more reactivethan NiþAl or WþAl. It was shown that this deformation relieson the mechanical properties of Ta and Al as well as the topologyof the TaþAl microstructure. When WþAl and NiþAl werereversed, they showed more deformation in the hard phase whichsuggests that in that configuration they would be more reactiveas well.

This behavior points to the importance of the topology of thehard phase. In general, it would be necessary to do 3D serialsectioning to determine topological properties in an arbitrarymicrostructure [31,32]. However, in this case, it is known that theTaþAl, NiþAl and WþAl materials were fabricated with 99.8%pure 325 mesh powders of Ta, Ni, W, and Al [12]. Images of thepowder feedstock can be seen in Wei et al. [13]. The Ni and Wparticles are strong, essentially convex, simply connected powderparticles. The Ta particles are spongy and appear to have non-zeroconnectivities. During explosive compaction Al could only flowaround the Ni or W particles, which deformed to a lesser extentbut their connectivity did not change. The Al was able to flow intothe Ta holes and pores. This led to the microstructures shown inFig. 2 with convex Ni and W particles surrounded by a matrix ofAl. The resulting TaþAl microstructure is one in which Ta and Alare probably interpenetrating. It is likely not the case that the Al isentirely composed of isolated islands as perhaps implied by Fig. 2.However, it is clearly the case that the connectivity of the Ta ismuch greater than that of the W or Ni. The 2D simulations, whilenot fully capturing the 3D topology of the material, do correctlycapture the influence of a more topologically connected hardphase on the deformation behavior of NiþAl, TaþAl, and WþAl.Taken with the experimental results, they also suggest that thedeformation and subsequent reaction behavior are more stronglyinfluenced by the topology than the interface density in thisimpact scenario.

The simulation and experimental results of the present worksuggest that the system that is most reactive, TaþAl, combines

higher topological connectivity in the reinforcement phase, duc-tility in both constituents, reduced tendency to work harden inthe reinforcement phase, high average density, good interfaceadhesion, and minimum separation between reactants.

4. Summary

Rod-on-anvil impact experiments conducted on explosivelycompacted equivolumetric mixtures of NiþAl, WþAl, and TaþAlindicate that TaþAl requires the lowest specific kinetic energy toinitiate reaction. Simulations conducted with CTH using realmicrostructures and stereological measurements on the NiþAl,WþAl, and TaþAl suggest that the observed reaction behavior isa result of higher topological connectivity in the reinforcementphase, ductility in both constituents, reduced tendency to workharden in the reinforcement phase, high average density, goodinterface adhesion, and minimum spatial separation betweenreactants.

Acknowledgments

This research was funded by ONR/MURI Grant no. N00014-07-1-0740. The authors acknowledge the support provided by theDOD SMART scholarship program for B. Aydelotte.

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