Use of Self-Propagating High-Temperature Synthesis 

Reactions in Refractory Alloy Fabrication 

by Laszlo J. Kecskes 

ARL-TR-3465 April 2005 

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Army Research Laboratory 

Aberdeen Proving Ground, MD 21005-5069 

ARL-TR-3465 April 2005 

Use of Self-Propagating High-Temperature Synthesis 

Reactions in Refractory Alloy Fabrication 

Laszlo J. Kecskes 

Weapons and Materials Research Directorate, ARL 

Approved for public release; distribution is unlimited. 
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Use of Self-Propagating High-Temperature Synthesis Reactions in Refractory 

Alloy Fabrication 

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14. ABSTRACT 

The titanium carbide (TiC) self-propagating high-temperature synthesis (SHS) reaction was used in a hot-explosiveconsolidation process to fabricate tungsten and molybdenum-based alloy billets. In this method, a Ti+C mixture is reacted to 

provide a source of heat for the refractory precursor powders. As the TiC reaction proceeds, the exterior of the alloy sample 

heats up rapidly, near or above 2000 °C, while its interior lags behind, heating up more gradually to 1600 °C. At the 

completion of the TiC reaction, the alloy sample’s exterior begins to cool, causing the temperature to equilibrate to an 

isothermal condition. At this time, the sample is compacted to high density by the application of an explosively generated 

pressure wave. Experiments were conducted to determine the relationship between the geometries of the Ti+C mixture and the 

alloy sample. For the Ti+C mixture, a doughnut arrangement was found to provide an optimum heating rate and temperature 

profile to the sample. The effects of the relative amount of the SHS material and alloy sample on the duration of the heating 

cycle, peak interior temperature, and rate of cooling were studied. The properties of the samples and TiC product were 

evaluated by density measurements, scanning electron microscopy, energy-dispersive x-ray spectroscopy, and x-ray diffraction 

analysis. The HEC method and the resultant product structures are described. 

15. SUBJECT TERMS

explosive compaction, refractory alloys, processing, microstructure, tungsten, molybdenum, titanium 

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Laszlo J. Kecskes 

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iii

Contents 

List of Figures iv

1. Introduction 1

2. Experimental Procedures 2

2.1 Heat-Balance Calculations ..............................................................................................2 

2.2 Temperature Measurements ............................................................................................2 

2.3 Compaction Experiments ................................................................................................4 

3. Results and Discussion 5

3.1 Temperature Measurements ............................................................................................5 

3.2 Compaction Characteristics...........................................................................................10 

3.3 Alloy Substructures .......................................................................................................13 

4. Conclusions 16

5. References 17

Distribution List 19 
iv

List of Figures 

Figure 1. Schematic diagram of the fixture used in the temperature measurements. .....................3 

Figure 2. Schematic diagram of the HEC fixture. ..........................................................................4 

Figure 3. Temperature data for W−Ti; nTiC/ns = 3.2. ......................................................................7 

Figure 4. Temperature data for W−Ti; nTiC/ns = 6.3. ......................................................................7 

Figure 5. Temperature data for W-Ti; nTiC/ns = 9.1........................................................................8 

Figure 6. Interior temperature data for W-Ti; nTiC/ns = 3.2, nTiC/ns = 6.3, and 

nTiC/ns = 9.1 is shown in 3-D......................................................................................................8 

Figure 7. Temperature data for Mo−Ti; nTiC/ns = 5.8. ..................................................................10 

Figure 8. Temperature data for W−Ti in the scaled-up fixture; nTiC/ns = 6.0. ..............................11 

Figure 9. Macrophotographs of a typical W−Ti billet showing top, bottom, and 

side views (sectioned billet placed bottom to bottom).............................................................12 

Figure 10. XRD scans of typical HEC billet core regions; W−Ti is shown in 7A 

and Mo−Ti is shown in 7B.......................................................................................................13 

Figure 11. XRD scans of typical HEC billet core regions; W Ti is shown in 7A 

and Mo Ti is shown in 7B........................................................................................................14 

Figure 12. SEMs of polished surfaces of typical HEC billet core regions; W−Ti is 

shown in 8A and Mo−Ti is shown in 8B. ................................................................................15 
1

1. Introduction 

Traditional forming routes of melting, casting, and forging are usually unsuitable for the preparation 

and processing of refractory-metal-based multiphase alloys. Often, when the extent of intermixing 

of the constituents is to be severely restricted, nonconventional powder metallurgy (PM) fabrication 

techniques, such as the use of high explosives for consolidation, may be sought out. PM is also 

more practical since near-net-shape products significantly reduce the need for machining. 

Explosive compaction is attractive because the rapid passage of a shock wave through the powder 

bed generates very high pressures and, therefore, frictional heating and melting occurs between 

adjacent particles. As such, intergrain bonding is produced on a microscopic level. The ease of 

compaction is strongly related to the nature of the powder assemblage, its particle size, size 

distribution, shape, and packing density. Due to the rapid densification rate, dynamically 

consolidated samples often suffer from low densities, poor bonding, and cracking (1). Highdetonation-velocity explosives (6000–8000 m/s) can produce the desired hardness, intergrain 

bonding, and compressive strength, however, with profuse cracking. Cracking may be reduced by 

using low-detonation-velocity explosives (2000–4000 m/s), but at the expense of consolidation 

quality. 

Conventional hot-explosive-consolidation (HEC) has been suggested as a way to reduce structural 

integrity problems (2, 3). In HEC, the powder mixture is sealed in a container, heated in a furnace, 

quickly removed, transferred into the compaction fixture, and consolidated. The primary advantage 

of heating is to decrease the material’s yield strength, thereby increasing its ductility. In recent 

years, several research groups (4–10) have explored HEC. 

To limit the intermixing of components that would occur during long heating times in a furnace, a 

modified HEC method was developed. This technique uses an in-situ titanium carbide (TiC) 

“chemical furnace” to heat the powder sample prior to compaction (10). In-situ heating also 

reduces the complexity of the assembly since the precursors are now heated and consolidated in the 

same fixture. 

The idea of “chemical heating” first appeared in conjunction with hot-working of metals (11). 

Since then, others have also attempted to shock-initiate admixed exothermic mixtures to locally 

enhance sample properties (5, 8, 12). 

The chemical furnace is simply an exothermic powder mixture of reactants that, when allowed to 

react, generates heat yielding high temperatures. Such reactions are a subclass of combustion 

synthesis or self-propagating high-temperature synthesis (SHS) reactions (13–15). In SHS, the 

products are directly synthesized from their elemental precursors, via the propagation of a solidsolid combustion front through a green powder compact (e.g., Ti + C → TiC). Once the process is 

initiated, the heat of reaction sustains the reaction until all of the reactants have been consumed. 

These reactions are characterized by rapid reaction rates (~1 cm/s), temperatures over 3000 °C, and 
2

the evolution of trapped impurities from the reactants. Potential advantages of the combustion 

synthesis process are that it requires little energy input and the processing time is reduced from 

hours to seconds. 

The primary objective of this effort was to examine the feasibility of scaling up the HEC process 

with the use of SHS reactions to fabricate high-density, two-phase tungsten-titanium (W−Ti) and 

molybdenum titanium (Mo−Ti) alloys. This required the examination of several processing 

parameters. Since the SHS reaction heating cycle is inherently short, a critical advantage of this 

technique would be the retention of desirable precursor properties, such as grain size distribution 

and morphology. Then, if possible, successful adjustments of the available heat and peak 

temperature achieved during heating would limit the extent of sintering, dissolution, and melting in 

each component. Therefore, the effects of sample configuration and the molar ratio of sample and 

chemical furnace on the structural evolution of the product were of primary concern. 

The test program was carried out in several steps, starting with baseline calculations to identify the 

amount of TiC needed to heat the alloy samples. After a series of auxiliary temperature 

measurements were performed, to experimentally validate the baseline computations, accurately 

determine the peak temperatures and heating times in the fixture, several W−Ti and Mo−Ti alloy 

billets were fabricated. Finally, the HEC process, designed for 5-cm-diameter, 

1-cm-thick billets, was scaled up to fabricate 7.5-cm-diameter, 1-cm-thick W−Ti billets. 

2. Experimental Procedures 

2.1 Heat-Balance Calculations 

Adiabatic heat-balance computations were performed for each of the sample systems to determine 

the required amount of chemical furnace (i.e., a baseline value). Briefly, in the calculations, the 

exothermic heat of formation of TiC (heat output) was equated to the heat required by the sample 

powder mixture (heat input) to elevate its temperature a known amount. Heat inputs for 

contributions from phase transitions (e.g., α−Ti → β−Ti) were included. However, the calculations 

did not consider heat losses or geometrical constraints in the reaction-compaction fixture. Results 

indicated that a ratio of moles of TiC to moles of sample (nTiC/ns) of 1.0 would cause the 

Ti component to melt. 

2.2 Temperature Measurements 

High-purity (99.9%) 12-μm W (Johnson Matthey, Ward Hill, MA), 99.9%-pure –325-mesh 

(equivalent to –44 μm) Mo (CERAC Inc., Milwaukee, WI), and 99.5%-pure –325-mesh Ti (Micron 

Metals, Gates City, UT) were used for the fabrication of the HEC alloy samples. Also, 99.5%-pure 

–325-mesh Ti (Micron Metals, Gates City, UT) and 99.9%-pure 2-μm graphite (C) (CERAC Inc., 

Milwaukee, WI) were used for the TiC chemical furnace. 
3

The W and Ti or Mo and Ti powder mixtures were ball-milled for at least 6 hr under an argon (Ar) 

atmosphere. The W:Ti ratio was 83:17 atomic percent (a/o) in all of the W−Ti alloys, and the 

Mo:Ti ratio was 82:18 a/o in the Mo−Ti alloy. A Ti:C molar ratio of 1.0 was maintained in all of 

the Ti and C powder mixtures. The Ti and C powders were also dry-mixed and ball-milled under 

Ar atmosphere. The Ti+C powder mixture was pressed at 44 MPa into 100-mm-outer-diameter, 

50-mm-inner-diameter, and 20-mm-thick doughnut-shaped green compacts. 

The auxiliary experiments were carried out with respect to the ratio of moles of TiC exothermic 

material to moles of alloy sample (nTiC/ns) to measure the sample temperature and rate of heating in 

the reaction-compaction fixture during the heating phase of the HEC process. For these 

measurements, the explosive container was not used. Instead, a series of five W−5Re/W−26Re 

thermocouples were inserted through the base of the reaction vessel into the middle of the powder 

sample. Two chromel-alumel thermocouples were also placed in the vessel, one near the electric 

match and another beneath the Ti+C doughnut, to mark the start and end of the TiC reaction. 

Figure 1 depicts the block diagram of the fixture with the lateral arrangement of the thermocouples. 

In these experiments, once the SHS reaction was initiated, the electro motive force (EMF) output of 

the thermocouples was recorded as a function of time. 

Figure 1. Schematic diagram of the fixture used in the temperature measurements. 

Several separate temperature-time measurements were made. In the first series for W−Ti, the 

nTiC/ns values were set to 3.2, 6.3, or 9.1. In the fourth measurement for Mo−Ti, the nTiC/ns was 

Piston

Anvil

Ti+C Mild-Steel Case Ti+C

ZrO2

Sample

Thermocouples Anvil

1 2 3 4 5

Igniter

Grafoil
4

5.9. Another measurement was made in the scale-up fixture. In the scale-up fixture, the overall 

lateral dimensions of the TiC chemical furnace and alloy sample were scaled up by a factor of 

1.5, while the vertical dimensions remained unchanged. The scale-up measurement was with 

W−Ti, using a nTiC/ns of 6.0. 

2.3 Compaction Experiments 

Figure 2 shows the schematic of the HEC reaction vessel. The vessel was designed to contain the 

Ti+C SHS reaction and, at the same time, to serve as the explosive compaction fixture. The lid 

assembly of the vessel was fitted with a container that held the explosive charge. The mild-steel 

vessel contained the alloy powder sample, the Ti+C green compact, zirconia (ZrO2), graphite sheet 

insulation, igniter mixture, electric igniter (not shown), and the explosive driver package. The 

explosive driver package consisted of a set of high-hardness steel anvils, aluminum piston (lid 

assembly), poly vinyl chloride (PVC) container holding powdered Amatol (80:20 NH4NO3:TNT) 

explosive (detonation velocity = 3850 m/s), detasheet booster, and the detonator. 

Figure 2. Schematic diagram of the HEC fixture. 

Booster

Amatol

Detonator

Piston

Igniter

Sample Ti+C

Mild-Steel Case

Ti+C

Anvil

ZrO2

Grafoil

Anvil
5

Because large quantities of volatile impurities are released by the TiC SHS reaction, vent holes were 

drilled into the mild-steel vessel. ZrO2 was used to limit heat losses from the reacting Ti+C 

compact and to protect the piston assembly and explosives from overheating. The graphite sheet 

functioned as an inert barrier, preventing the diffusion of contaminants into the sample. 

Commercial model-rocket igniters were used to initiate the SHS reaction. The electric match was 

remotely activated, which, in turn, initiated the igniter mixture and Ti+C green compact. As the 

TiC reaction proceeded to completion, normally in about 3 s, the alloy powder sample was heated 

(via diffusion) to high temperatures. When the optimum amount of time had elapsed, the piston, 

accelerated by the detonation of the explosive, densified both the sample and the TiC doughnut. 

When the fixture and its contents cooled to ambient temperature, the sample billet was extracted. 

After the extraction of the billet, the typical postconsolidation characterization of the samples 

consisted of macrophotography, sectioning, determination of the bulk density, and metallographic 

preparation of scanning electron microscopy (SEM), energy-dispensing x-ray spectroscopy (EDS), 

and x-ray diffraction analysis (XRD) samples. The compaction uniformity of the billets was 

evaluated by density measurements. The billets were cut into several roughly equal sections, and 

the mass density was measured with a water immersion technique. The measurements revealed that 

the billets were uniformly dense with a variation of about ±0.5%. The billets were further sectioned 

into wafers (parallel or perpendicular to the compaction direction) for SEM, EDS, and XRD. Each 

XRD sample was polished and mounted into sample holders, then step-scanned through a 2θ range 

of 20–120° using a 0.025° increment and a 5-s dwell time. SEM, EDS, and XRD were also used to 

examine the precursor powders. It may be noted that the graphite was not analyzed beyond SEM. 

Three W−Ti billets were produced, using an explosive charge mass to piston mass ratio (c/m) of 

1.7 ± 0.1. The nTiC/ns values were 3.2, 6.4, or 8.3. Note the slight deviation from the nTiC/ns values 

used in the temperature measurements. The c/m value of ~1.7 was found to produce samples with 

optimum (i.e., highest density without excessive cracking) results. For the W−Ti billet that was 

made in the scale up fixture, the c/m was 2.0 and the nTiC/ns was 6.4. Finally, an Mo−Ti billet was 

consolidated with a c/m of 2.1 and nTiC/ns of 5.8. 

3. Results and Discussion 

Preliminary configurational experiments determined that placement of the TiC chemical furnace in 

a concentric geometry is most suitable for uniform consolidation. Compactions using layered (e.g., 

sandwich) and complete envelopment geometries did not achieve simultaneous ignition of separated 

chemical furnace elements, resulting in nonuniform compactions and sample densities. Sample 

billets also tended to contain large radial and laminar macrofractures. 

3.1 Temperature Measurements 

The results of the measurements of the precursor-powder-bed temperature during the heating phase 

were essential for the consolidation of the alloy billets. The experiments revealed the peak 
6

temperatures and thermal gradients within the sample powder bed and also indicated the optimum 

time for consolidation. The results also showed that there were significant heat losses in the HEC 

fixture. 

The data for W−Ti corresponding to nTiC/ns values of 3.2, 6.3, and 9.1 are presented in figures 3–5. 

In the lower left-hand corner of figures 3–5, the pair of vertical arrows marks the duration of the 

TiC reaction. (Note, in figures 3 and 5, the second arrows are partially masked by the thermocouple 

trace.) The temperature-time trace at the center of the powder bed (TC no. 3) for each case is shown 

on an expanded scale in figure 6. As expected, with an increasing amount of TiC, the duration of 

the SHS reaction increases. In all three cases, the temperature at the edge of the powder beds rises 

rapidly to reach a broad peak; drops; goes through a plateau; then, due to continued heat losses, 

steadily decreases. As the heat diffuses inward to the sample center, the temperature profile is 

changed. At the center, a more sluggish rise in the temperature and the absence of the short-term 

peak result in a wider plateau and subsequent decline only. Thus, an immediate conclusion from the 

data is that the periphery of the powder bed undergoes a different thermal history than the 

midregions. 

For nTiC/ns = 3.2, the peak temperature is 1500 °C at one edge and is near the melting point of Ti at 

the other (Tm Ti = 1670 °C). The center reaches only 1250 °C, well below Tm Ti. Whereas, for 

nTiC/ns = 6.3, the edge temperature reaches 1900 °C at one end and 2100 °C at the other. In this 

case, the center temperature is higher, reaching 1550 °C, but does not exceed Tm Ti. Finally, for 

nTiC/ns = 9.1, the powder bed reaches temperatures above 2400 °C at its periphery and 1650 °C at its 

center. The heating rates for nTiC/ns values of 3.2 and 6.3 are similar, but, for nTiC/ns = 9.1, it is 

different. Specifically, the temperature rise at the center is more rapid, much like that of the 

periphery. With nTiC/ns values of 6.3 and 9.1, the critical aspect of the data is that at about 60 s after 

the initiation of the TiC reaction, the entire W+Ti powder bed becomes effectively isothermal. For 

nTiC/ns = 3.2, the interior of the powder bed does not reach an isothermal plateau at 60 s and is 

subject to moderate temperature gradients. Only considerably later (over 120 s; not shown) at a 

much lower temperature, during overall cooling, is thermal equilibrium reached. 

Differences of the temperature profiles are primarily due to the amount of TiC furnace and slight 

changes with nTiC/ns of the TiC doughnut/sample geometry. The packing density ratio of the TiC 

furnace and sample is ~3. As a result, the respective volumes of the chemical furnace and the 

sample are roughly the same for nTiC/ns = 3.2. Whereas, the volume of TiC is twice that of the 

sample for nTiC/ns = 6.3 and three times that of the sample for nTiC/ns = 9.1. Therefore, the TiC 

doughnut completely encloses for nTiC/ns = 9.1, and partially encloses for nTiC/ns = 6.3, the ZrO2

insulation above and beneath the sample. Then, as the TiC reacts, the ZrO2 is also heated. The 

hotter insulation, in turn, provides greater heating to the sample, resulting in a faster internal 

temperature rise. However, for nTiC/ns = 3.2, most of the ZrO2 lies outside of the TiC doughnut. 

Without direct heat from the TiC reaction, the ZrO2 acts as a heat sink, lessening the interior 

temperature rise. 
7

Figure 3. Temperature data for W−Ti; nTiC/ns = 3.2. 

Figure 4. Temperature data for W−Ti; nTiC/ns = 6.3.

0

500

1000

1500

2000

2500

0 20 40 60 80 100

Time [s]

Temperature [C]

TC#1

TC#2

TC#3

TC#4

TC#5

Start of Reaction

End of Reaction

0

500

1000

1500

2000

2500

0 20 40 60 80 100

Time [s]

Temperature [C]

TC#1

TC#2

TC#3

TC#4

TC#5

Start of Reaction

End of Reaction
8

Figure 5. Temperature data for W-Ti; nTiC/ns = 9.1. 

Figure 6. Interior temperature data for W-Ti; nTiC/ns = 3.2, nTiC/ns = 6.3, and nTiC/ns = 9.1 is 

shown in 3-D. 

0

500

1000

1500

2000

2500

0 20 40 60 80 100

Time [s]

Temperature [C]

TC#1

TC#2

TC#3

TC#4

TC#5

Start of Reaction

End of Reaction

0

200

400

600

800

1000

1200

1400

1600

1800

0 20 40 60 80 100

Time [s]

Temperature [C]

n(TiC)/n(s)=3.2

n(TiC)/n(s)=6.3

n(TiC)/n(s)=9.1
9

The profile of the center thermocouple in figure 3 (nTiC/ns = 3.2) is also noisier, and it exhibits a 

150–200 °C temperature drop lasting for about 20 s. After this temporary heat loss, the temperature 

continues to rise. The center thermocouple profile in figure 4 (nTiC/ns = 6.3) reflects a similar 

behavior, but the effect is less pronounced. The increased temperature fluctuation and/or decrease 

(heat loss) are probably attributed to a partial failure (cracking and delamination) of the ZrO2

insulation beneath the sample powder bed. The insulation is cooler for the smaller nTiC/ns value; a 

cooler ZrO2 may be more prone to greater thermal gradients that would likely generate higher 

internal stresses and cause cracking. Such cracks that disrupt the integrity of the thermally 

insulating barrier could possibly account for the observed temperature fluctuations and/or drops. 

The data in figures 3–6 reveal that the nTiC/ns significantly affects the heating rates and peak 

temperatures attained within the W+Ti powder bed. Even with the smallest nTiC/ns, the temperature 

at the sample edge nearly reaches Tm Ti. A threefold increase of nTiC/ns raises the edge temperature 

by 1000 °C. The center is less sensitive; the interior temperature varies only by ~400 °C. Ideally, 

for best compaction results, the nTiC/ns should be adjusted such that as much as possible of the bulk 

of the alloy powder bed is near, but does not exceed the Tm Ti, and excessive overheating of the 

periphery is minimized. However, unless a sufficiently large nTiC/ns is used, the entire W+Ti sample 

will not become isothermal during the early part of the heating cycle. The nTiC/ns of ~6.0 satisfies 

both of these criteria. 

Aside from a more sluggish rise in the interior of the powder bed, the Mo-Ti system exhibited 

the same temperature response as the W-Ti system (figure 7). Note that the delayed temperature 

rise of TC no. 3 occurs in the interior of Mo+Ti powder bed. Such differences in the rate of 

temperature change in the precursor powder bed are primarily influenced by kinetic factors such 

as the thermal diffusivity of the powder mixture. Because Mo has a lower thermal diffusivity 

than W (16), all other factors in the fixture being constant, the thermal diffusivity of the Mo+Ti 

mixture will also be less. With a lower thermal diffusivity, the rate of change of the temperature 

in the Mo+Ti bed would be expected to be more sluggish. 

Precompaction temperatures were also measured in the scale-up fixture. The lateral dimension of 

the W-Ti alloy powder bed was scaled from 50 to 75 mm. Within the scale-up fixture, the 

dimensional scaling of the internal components was identical to that of the standard fixture. 

However, because of the larger diameter of the powder bed, all five thermocouples fit along a single 

radius. Figure 8 indicates the relative positions along the radius. The thermocouple data reveal that 

the peak temperatures are similar to those obtained previously. However, the rate of heating in the 

powder bed is proportionately slower and the position of the isothermal plateau is shifted from 60 to 

140 s. The delay time does not scale linearly with the lateral fixture dimensions, but, rather, scales 

quadratically, i.e., with the volume. A peak temperature of 1800 °C was measured at the edge, with 

none of the interior temperatures exceeding 1500 °C. At the edge, the temperature rises to its 

maximum value, then declines gradually through the isothermal plateau. The rise in the interior is 

more moderate, monotonically reaching the isothermal plateau. Beyond the plateau region, the 

powder bed slowly cools to ambient temperature. 
10

Figure 7. Temperature data for Mo−Ti; nTiC/ns = 5.8. 

The extended delay time alters the relative ease of operations of the HEC process. Primarily, to 

protect the explosives from overheating or melting. The longer delay times needed to reach the 

isothermal condition will require a remotely operable explosive handling and transfer system, much 

like that used by Gorobtsov and Roman (3). More importantly, as illustrated with the increase from 

50- to 75-mm-diameter billets, further lateral scaling with much longer delay times will dictate an 

alternative method of encapsulation of the sample to prevent exposure to air or other contamination. 

3.2 Compaction Characteristics 

During compaction of the contents of the HEC fixture, two main factors had to be considered. 

These were the yield strengths of the internal components and their respective temperature histories 

prior to compaction. From the bowed cross-sectional appearance of components, corresponding to 

nonuniform deformation, it was deduced that the compressive wave was not entirely planar (i.e., 

uniaxial) during compaction (17). Additionally, from the shape of the deformed alloy billet near its 

edges, it was inferred that the compaction pressure was at least comparable to, if not higher than, the 

high-temperature yield strengths of β−Ti and W or Mo. 

Differences in yield strengths of the piston, steel anvils, ZrO2, TiC, and the W−Ti billet resulted in a 

limited, radially inward deformation of the more plastic components at the expense of others. Only 

the local distortion of the TiC and ZrO2 disks, in direct contact with the sample powder bed, had a 

marked effect on the shape of the compacted alloy billet. 

0

500

1000

1500

2000

2500

0 20 40 60 80 100

Time [s]

Temperature [C]

TC#1

TC#2

TC#3

TC#4

TC#5

Start of Reaction

End of Reaction
11

Figure 8. Temperature data for W−Ti in the scaled-up fixture; nTiC/ns = 6.0. 

The ZrO2 disk above the sample powder bed remained relatively cool. Therefore, it did not entirely 

sinter after consolidation. The upper one-third of the disk disintegrated into a fine powder; the 

lower two-thirds of the disk was sintered but with crazing cracks. Because of excessive overheating 

by the TiC reaction during and after consolidation, the ZrO2 disk below the powder bed was more 

sintered and contained fewer cracks, but its outer periphery was completely eroded. 

The recovered TiC doughnut was highly friable and was usually compacted to about 

90% theoretical density (TD) only. However, at the time of consolidation, the doughnut was 

sufficiently plastic to deform, filling the fixture cavity. Its outer wall was straight, matching the 

contour of the fixture; its inner wall was irregular, matching the contour of the inner core (consisting 

of the anvils, ZrO2 disks, and sample billet). The TiC formed an annular shear lip below the billet, 

partially filling the space vacated by the eroded lower ZrO2 disk. This flow was incomplete, and a 

slight gap remained between the smaller ZrO2 disk and the TiC lip. The gap was filled by the 

collapsing sample powder bed (18) upon consolidation. 

The disk-shaped HEC billets were ~9 mm thick with a diameter of 50 mm. Figure 9 presents a 

macrophotograph showing the top, bottom, and side views of a typical W−Ti billet. The cut halves 

of the disk are placed with their bottom surfaces facing inward to form the side view. The top 

surface of the billet is mostly flat, while the bottom surface is slightly convex with a ridge forming 

at its edge. This small ridge is about 6.4 mm wide and 3.2 mm high and readily fractures off. The 

texture of the billet’s side ranges from very fine to coarse striations and creases. 

0

500

1000

1500

2000

2500

0 40 80 120 160 200 240

Time [s]

Temperature [C]

TC#1; r/Ro=0.90

TC#2; r/Ro=0.60

TC#3; r/Ro=0.30

TC#4; r/Ro=0.17

TC#5; r/Ro=0.00

Start of Reaction

End of Reaction
12

Figure 9. Macrophotographs of a typical W−Ti billet showing top, bottom, and side views (sectioned billet 

placed bottom to bottom). 

The billets could be divided into an inner core and a periphery. Typically, the core comprised a 

disk with a diameter of ~35–40 mm, while the periphery comprised the outer 10- to 15-mm 

annulus of the billet. The periphery could be further divided into a 5-mm-wide outer edge and an 

intermediate transitional region. The edge usually contained minor annular and radial cracks. No 

evidence of horizontal hairline cracks was found in the core of the billets. Occasionally, some of 

the cores contained a random network of vertical hairline cracks. After sectioning such cores, the 

cracks were seen to occur along diagonals, originating at the top or bottom billet surface and 

terminating at the sides. Most likely, these cracks were caused by the reflection of the compressive 

wave from the free surfaces of the fixture during consolidation. 

Results of the experiments with varying nTiC/ns indicated that the density of W−Ti decreases with 

nTiC/ns. A density of 99.1% TD was achieved at nTiC/ns = 3.2, while a density of 98.6% TD was 

obtained with nTiC/ns = 6.4. Lastly, a density of only 96.3% TD was reached for nTiC/ns = 8.3. It is 

speculated that, during consolidation at higher nTiC/ns values, the outer zone of the sample becomes 

more plastic, causing exit routes to be closed, thereby trapping the interior porosity. As expected, 

due to the lower Tm of Mo and greater thermal softening, densification was better for the Mo−Ti 

sample; specifically, a density of 99.1% TD was obtained using nTiC/ns = 5.8. 

In a preliminary attempt to examine the feasibility of scale up, the standard weldings that held the 

fixture parts together failed during detonation, causing the loss of containment. The W−Ti billet 

was fully dense (98.2% TD) but was recovered in several fragments. For the second attempt, the 

welds were reinforced and strengthened to avoid detonation damage. This test resulted in a whole 

Top View Bottom View Side View
13

disk-shaped billet that was fully dense (99.1% TD), with a slightly convex bottom surface. While 

successful, primarily because of weight, size, and heating time limitations, further lateral scaling of 

the HEC fixture will necessitate the use of machine-operated assembly, delivery, and recovery 

operations. 

3.3 Alloy Substructures 

SEM of the W powder revealed polyhedral particles, with a coarse surface texture, ranging in size 

from 5 to 20 μm. The Mo was polyhedral and faceted, ranging in size from 10 to 30 μm. The Ti 

was irregular, having faceted to elongated shapes, with a size range of 10–40 μm. 

Baseline XRD measurements of the precursor powder lattice parameters were in agreement with 

nominal values for Mo, W, and Ti. XRD scans in figures 10 and 11 show that the cores of the 

alloys have a two-phase structure. In figure 10 for W−Ti with nTiC/ns = 6.4, the W diffraction 

peaks are clearly seen. Likewise, in figure 11 for Mo−Ti with nTiC/ns = 5.8, the diffraction peaks 

for Mo are easily identified. A second phase in each alloy appears as a shoulder (with an inflection 

point) on the low-angle side of each main peak. In Mo−Ti, the shoulders are more magnified and 

much wider, extending to lower angles. Unlike W−Ti, a few distinct peaks also appear on the 

shoulders. 

Figure 10. XRD scans of typical HEC billet core regions; W−Ti is shown in 7A and Mo−Ti is shown in 7B. 

100

1000

10000

100000

20 30 40 50 60 70 80 90 100 110 120

Scattering Angle [2Theta]

Intensity [Counts]

W(110)

TiW(110)

W(200)

W(211)

W(220)

W(310)

W(222)

TiW(200)
14

Figure 11. XRD scans of typical HEC billet core regions; W Ti is shown in 7A and Mo Ti is shown in 7B. 

SEM of transverse (perpendicular) and longitudinal (parallel) cross sections (with respect to the 

compaction axis) revealed a preferentially oriented two-phase structure. The distribution of the 

primary and secondary phases is isotropic in the transverse section; however, both phases tend to be 

elongated along a transverse plane in the longitudinal section. Consistent with compaction wherein 

grains deform plastically without exceeding the breaking stress in the powder body, this orientation 

effect arose from the uniaxial collapse of the initially isotropic powder bed. 

Figure 12 illustrates typical axial cross sections from the alloy core regions. W−Ti with 

nTiC/ns = 6.4 is shown in figure 12, and Mo−Ti with nTiC/ns = 5.8 is shown in figure 12. 

Observation of the W−Ti core reveals loose W grain aggregates (light gray) bounded by a partially 

connected Ti-rich matrix (dark gray). The partial discontinuity of the matrix comes from the 

particle size difference of the precursors, the relative melting points, and the temperature before 

and at the time of compaction. The W grains have retained their original size of 10–20 μm. The 

dimensions of the Ti-rich areas, likewise, correspond to the size of the Ti precursor. EDS indicates 

that the W in the aggregates is nearly pure and the fairly uniform Ti-rich phase has a composition 

of ~25W−75Ti a/o. The second phase observed in the XRD measurements is this Ti-rich β−Ti/W 

solid solution (19). 

100

1000

10000

100000

20 30 40 50 60 70 80 90 100 110 120

Scattering Angle [2Theta]

Intensity [Counts]

TiMo(110)

TiMo(200) TiMo(211)

TiMo(220) TiMo(310)

Mo(222)

Mo(310)

Mo(220)

Mo(110) Mo(200) Mo(211)
15

Figure 12. SEMs of polished surfaces of typical HEC billet core regions; W−Ti is shown in 8A and Mo−Ti is 

shown in 8B. 

Unlike W−Ti, the Mo grain aggregates (light gray) in the Mo−Ti alloy are more interconnected, 

surrounding mostly isolated Ti-rich matrix regions (medium to dark gray). The Mo grains have a 

size of 10–30 μm. The dimensions of the Ti-rich matrix regions correspond to those of the Ti 

precursor. EDS indicates that the Mo in the aggregates is mostly pure. As evident in the figure, 

the Ti-rich matrix phase is segregated and nonuniform. The composition of the medium-gray areas 

ranges from 25Mo−75Ti to 15Mo−85Ti a/o. In contrast, the composition of the dark-gray areas is 

mostly pure Ti. The low-angle shoulders, noted in the XRD scans, correspond to the Ti-rich 

β−Ti/Mo solid solution (19). The increased width and height of the shoulders support the observed 

compositional nonuniformity. Finally, some of the distinct shoulder peaks at 2θ of 34.9, 38.7, 

52.2, 63, and 70.1° are probably associated with the mostly pure Ti regions. 

The presence of the secondary matrix phase obscures the evolution of the alloy microstructures. 

The added complexity derives from the intermixing of the constituents at elevated temperatures. In 

the W−Ti alloy, because of the moderately fast diffusion of W in Ti (20) and the fact that the 

sample is consolidated after a sufficiently long time delay, it was expected that the composition of 

the β−Ti/W solid solution would equilibrate during heating. That is, the concentration gradient 

introduced by a solid-state dissolution of the W grain outer layers in Ti would be quickly reduced 

by fast diffusion. In the absence of a composition gradient (i.e., that introduced by slow diffusion), 

the rate-controlling step of the process is the dissolution of W during heating. The relative 

compositional uniformity of the matrix in the W−Ti alloy core is in agreement with this 

assessment. 

Within the context of the HEC system thermodynamics, the Mo−Ti matrix ought to be 

compositionally uniform similar to that found in the W−Ti alloy. However, despite of a higher 

diffusion rate of Mo in Ti (about three times than that of W in Ti) and the strong similarities 

W-Ti Mo-Ti
16

between the Mo−Ti and W−Ti phase diagrams (19), the Mo−Ti matrix composition is not uniform. 

Due to the lower thermal diffusivity of Mo, the 60-s compaction delay time is most likely 

insufficient to allow adequate equilibration of the β−Ti/Mo matrix. The more sluggish temperature 

rise reduces the effectiveness of heating, and a shorter time at temperature, in turn, prevents the 

homogenization of the matrix. 

4. Conclusions 

A technique whereby high-density, monolithic, single- or multi-phase alloy billets can be fabricated 

has been demonstrated. HEC was conceived primarily for use with difficult-to-consolidate 

refractory alloys when conventional methods are not applicable. In this technique, the precursors 

are first heated by an external exothermic combustion synthesis reaction. When such a reaction 

takes place in a properly insulated fixture, the powdered precursors can be quickly heated to very 

high temperatures. Compaction to full density is accomplished by the pressure wave generated 

during the detonation of an explosive. The thermal diffusion length of the heating heat pulse only 

limits the lateral dimensions of the billets. 

The fundamental appeal of this technique hinges on the shortened processing cycles that are 

characteristic of this process. Concurrent with the application of dynamic pressures, rapid heating 

and cooling permit the formation and retention of nonequilibrium structures that are otherwise 

unobtainable with conventional forming methods. It was shown with W−Ti and Mo−Ti alloys that 

adjustment of the heating rate and peak temperature during heating, as well as the thermal 

management of heat losses upon cooling, could affect the compositional intermixing and 

heterogeneity within the sample billets. 

The W−Ti alloy billet core is of high density and has a two-phase structure of W particles 

embedded in a partially continuous Ti-rich β−Ti/W solid-solution matrix. Experiments with the 

Mo−Ti system resulted in similar overall alloy characteristics. Within the interior, unlike the 

matrix in W−Ti, the matrix in Mo−Ti is nonuniform. This nonuniformity arises from the sluggish 

kinetics of the heating phase, causing the dissolution of Mo to be incomplete. Phase uniformity 

may be improved by HEC fixture modifications to increase the amount of ZrO2 insulation or the 

type of chemical furnace material such that the duration of the isothermal plateau is lengthened. 
17

5. References 

1. Linse, V. D. Dynamic Compaction of Metal and Ceramic Powders. Sagamore Army 

Material Research Conference. Innovation in Materials Processing; Bruggeman, G., 

Weiss, V., Eds.; Plenum Press: New York, NY, 1983, 30, 381–404. 

2. Cross, A. Try Hot Explosive-Compacting for Sintered Powder Products. Iron Age 1959, 

184 (26), 48–50. 

3. Gorobtsov, V. G.; Roman, O. V. Hot Explosive Pressing of Powders. International Journal 

of Powder Metall. and Powder Tech. 1975, 11 (1), 55–60. 

4. Wang, S. L.; Meyers, M. A.; Szeket, A. Warm Shock Consolidation of IN 718 Powder. 

Journal of Material Science 1988, 23 (5), 1786–1804. 

5. Ferreira, A.; Meyers, M. A.; Thadhani, N. N.; Chang, S. N.; Kough, J. R. Dynamic 

Compaction of Titanium Aluminides by Explosively Generated Shock Waves: Experimental 

and Materials Systems. Metall. Trans. A 1991, 22 (3), 685–695. 

6. Taniguchi, T.; Kondo, K. Hot Shock Compaction of α−Alumina Powder. Adv. of Ceramic 

Materials 1988, 3 (4), 399–402. 

7. Yoshida, M.; Yoshioka, Y.; Kimura, Y.; Hirabayashi, H.; Sawaoka, A. B.; Pruemmer, R. A. 

Explosive Hot Consolidation of Alumina Powder. Influence on Structure. Aerospace, 

Refractory and Advanced Materials Advances in Powder Metallurgy; Metal Powder 

Industries Federation: Princeton, NJ, 1991, no. 6, pp 199–209. 

8. Shang, S. S.; Hokamoto, K.; Meyers, M. A. Hot Dynamic Consolidation of Hard Ceramics. 

Journal of Material Science 1992, 27 (20), 5470–5476. 

9. Peikrishvili, A.; Japaridze, L.; Chikhradze, N.; Chagelishvili, E. Possibilities of Obtaining 

Combined Samples from Tungsten-Based Alloys by High-Temperature Shock Wave 

Treatment. Metallurgical and Materials Applications of Shock-Wave and High-Strain-Rate 

Phenomena; Murr, L. E., Staudhammer, K. P., Meyers, M. A., Eds., Elsevier: Amsterdam, 

The Netherlands, 1995, chap. 14, pp 99–108. 

10. Kecskes, L. J.; Hall, I. W. Hot Explosive Consolidation of W−Ti Alloys. ARL-TR-669; 

U.S. Army Research Laboratory: Aberdeen Proving Ground, MD, 1995. 

11. Rhinehart, J. S.; Pearson, J. Explosive Working of Metals; MacMillan Company: 

New York, NY, 1963; chap. 9, p 257. 

12. Akashi, T; Sawaoka, A. B. High Density Compacts. U.S. Patent 4,655,830, 1987. 
18

13. Munir, Z. A. Synthesis of High-Temperature Materials by Self-Propagating Combustion 

Methods. Ceramics Bulletin 1988, 67 (2), 342–349. 

14. Materials Processing by Self-Propagating High-Temperature Synthesis (SHS). 

DARPA/ARMY SHS Symposium Proceedings, MTL-SP-87-3, Gabriel, K. A., Wax, S. G., 

McCauley, J. W., Eds.; U.S. Government Printing Office: Washington, DC, 1985. 

15. Combustion and Plasma Synthesis of High-Temperature Materials. Munir, Z. A., 

Holt, J. B., Eds.; VCH Publishers, Inc.: New York, NY, 1990. 

16. Handbook of Chemistry and Physics. Weast, R. C., Ed.; Chemical Rubber Company Press: 

Cleveland, OH, 54th ed, 1973. 

17. Gourdin, W. H. Dynamic Consolidation of Metal Powders. Prog. Material Science 1986, 

30 (1), 39–80. 

18. Page, N. W.; Killen, P. D. Powder Flow During Dynamic and Quasistatic Compaction. 

Powder Metal. 1987, 30 (4), 233–239. 

19. Binary Alloy Phase Diagrams. Massalski, T. B., Okamoto, H., Subramanian, P. R., Kacprzak, 

L., Eds.; William W. Scott: New York, NY, 1990. 

20. Askill, J. Tracer Diffusion Data for Metals, Alloys, and Simple Oxides; Plenum Data 

Corporation: New York, NY, 1970, pp 52–53. 
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