December 2001

NASA/TM-2001-211264

Effectiveness of Diffusion Barrier Coatings

for Mo-Re Embedded in C/SiC and C/C

David E. Glass

Langley Research Center, Hampton, Virginia

Ravi N. Shenoy

Lockheed Martin Engineering & Sciences, Hampton, Virginia

Zengmei Wang

Chromealloy Gas Turbine Corporation, Orangeburg, New York

Michael C. Halbig

Glenn Research Center, Cleveland, Ohio
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December 2001

NASA/TM-2001-211264

Effectiveness of Diffusion Barrier Coatings

for Mo-Re Embedded in C/SiC and C/C

David E. Glass

Langley Research Center, Hampton, Virginia

Ravi N. Shenoy

Lockheed Martin Engineering & Sciences, Hampton, Virginia

Zengmei Wang

Chromealloy Gas Turbine Corporation, Orangeburg, New York

Michael C. Halbig

Glenn Research Center, Cleveland, Ohio
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either expressed or implied, of such products or manufacturers by the

National Aeronautics and Space Administration, Lockheed Martin

Engineering & Sciences, or Chromealloy Gas Turbine Corporation.
Effectiveness of Diffusion Barrier Coatings For

Mo-Re Embedded in C/SiC and C/C

David E. Glass, NASA Langley Research Center, Hampton, VA 23681

Ravi N. Shenoy, Lockheed Martin Engineering & Sciences, Hampton, VA 23681

Zengmei Wang, Chromealloy Gas Turbine Corporation, Orangeburg, NY 10962-2515

and

Michael C. Halbig, NASA Glenn Research Center, Cleveland, OH 44135

Abstract

Advanced high-temperature cooling applications such as heat-pipe-cooled leading edges

may often require the elevated-temperature capability of carbon/silicon carbide (C/SiC) or

carbon/carbon (C/C) composites in combination with the hermetic capability of metallic

tubes. In some applications, it may be advantageous to co-process metallic tubes with the

composite. Under those conditions, it is important to understand the effect of the harsh

environment of co-processing and additional thermal exposure on the metallic tubes in

contact with carbon and silicon carbide. In this paper, the effects of C/SiC (using pyrolytic

carbon (PyC) and CVI SiC processing steps) and C/C on tubes fabricated from several

different refractory metals were evaluated. Though Mo, Nb, and Re were evaluated in the

present study, the primary effort was directed toward two alloys of Mo-Re, namely, arc cast

Mo-41Re and powder metallurgy Mo-47.5Re. Samples of these refractory metals were

subjected to either the PyC/SiC deposition or embedding in C/C. MoSi2(Ge), R512E, and

TiB2 coatings were included on several of the samples as potential diffusion barriers. Also,

several of these samples were exposed to time at temperature in a vacuum as controls. The

effects of the processing and thermal exposure on the samples were evaluated by

conducting burst tests, microhardness surveys, and scanning electron microscopic

examination (using either secondary electron or back scattered electron imaging and energy

dispersive spectroscopy). The results showed that a layer of brittle Mo-carbide formed on

the substrates of both the uncoated Mo-41Re and the uncoated Mo-47.5Re, subsequent to

the C/C or the PyC/SiC processing. Both the R512E and the MoSi2(Ge) coatings were

effective in preventing not only the diffusion of C into the Mo-Re substrate, but also the

formation of the Mo-carbides. However, Mo-silicide layers formed on the Mo-Re

substrates as a result of the chemical interaction with both the R512E and the MoSi2(Ge)

coatings. This Mo-silicide layer grew with heat treatment time, albeit at a slower rate than

the Mo-carbide layer in the uncoated condition, and thus reduced the effective thickness of

the substrate. The development of lateral and transverse cracks in the R512E coating may

have diminished the efficacy of that coating. For the Mo-47.5Re alloy, burst strength

degraded moderately (relative to the uncoated condition) when the alloy tubes were coated

with either MoSi2(Ge) or TiB2. The degradation in burst strength was more drastic

subsequent to the coating followed by the PyC/SiC deposition, with the degradation from

TiB2 being more severe than from MoSi2(Ge). MoSi2(Ge) thus appeared to be more

effective than TiB2 as a diffusion barrier coating for Mo-47.5Re. However, none of the

coatings were effective at preventing both C and Si diffusion without some degradation of

the substrate.
2

Table of Contents

Introduction ......................................................................................................................................3

Materials and Coatings .....................................................................................................................5

Metallurgical Characterization...........................................................................................................3

R512E Coating ...........................................................................................................................6

MoSi2(Ge) Coating...................................................................................................................12

AC Mo-41Re......................................................................................................................12

PM Mo-47.5Re ..................................................................................................................14

AC Mo................................................................................................................................18

TiB2 Coating.............................................................................................................................22

AC Mo................................................................................................................................22

PM Mo-47.5Re ..................................................................................................................23

Burst Tests......................................................................................................................................26

Concluding Remarks ......................................................................................................................29

Acknowledgement...........................................................................................................................29

References......................................................................................................................................30

Appendix: Literature Survey ...........................................................................................................34

Refractory Metal Compatibility with Carbon and Silicon Carbide ............................................34

Carbon Diffusion Barriers........................................................................................................37

Platinum, Iridium, and Rhenium..........................................................................................37

Titanium Diboride...............................................................................................................39

Molybdenum Disilicide ......................................................................................................39

Overview...................................................................................................................................40
3

Introduction

Ceramic matrix or carbon/carbon composite (generically termed refractory composite herein)

structures with embedded refractory-metal tubes may have numerous high-temperature applications.

One such application is for light-weight heat-pipe-cooled leading edges with high-heat-flux

capabilities, where refractory-metal tubes embedded in the refractory-composite structure serve as

the containers for liquid-metal heat pipes. Refractory metals are utilized since they exhibit high

melting points and high strengths at elevated temperatures, and can also withstand high processing

temperatures required for manufacturing refractory-composite materials. In addition, they have the

ability to be hermetically sealed enabling leak-free operation.

One of the complications encountered embedding a refractory-metal tube in a refractorycomposite structure is lack of chemical compatibility between the two materials. Many of the

refractory metals form carbides or silicides when in contact with carbon or silicon at elevated

temperatures. The formation of carbides or silicides can lower the strength and ductility of the

refractory metal. In addition, for heat-pipe applications, carbon or silicon may diffuse through the

metal tube (i.e., heat-pipe container) and contaminate the heat pipe.

The Appendix presents a brief overview of literature that exists with respect to laboratorybased physical metallurgy aspects of refractory-metal compatibility with C and SiC, and the use of

various diffusion barrier coatings to minimize the unfavorable interactions between the metals and

the refractory composite. It is nevertheless important to study the effectiveness of diffusion barrier

coatings on refractory metals and alloys of practical concern. Concurrently, it is also pertinent to

evaluate the mechanical integrity of such systems comprising the refractory metal, the diffusion

barrier coating, and the C or SiC environment before and after exposure to targeted hightemperature environments. It is with this objective that the present study was undertaken.

This report discusses an evaluation of the effects of pyrolytic carbon (PyC) followed by SiC

deposition, and C/C on several different refractory metals. The PyC/SiC deposition, representing a

C/SiC densification process, was applied by Honeywell Advanced Composites, Inc. (HACI),

Newark, DE and the C/C processing was performed by Carbon-Carbon Advanced Technologies

(C-CAT), Fort Worth, TX. The refractory alloys evaluated were Mo-Re, Mo, Nb, and Re which

were formed into tubes by various processes. R512E (a proprietary Fe, Si, Cr-rich coating

deposited by Hitemco), TiB2, and MoSi2(Ge) coatings were applied to the refractory metals as

potential diffusion barriers. The methods of evaluating the effectiveness of the coatings and the

effect of the C and SiC on the refractory metals were metallurgical characterization, including

micorhardness and SEM, and burst test. The results of the metallurgical characterization are

discussed first, followed by the results of the burst test.

Metallurgical Characterization

R512E, MoSi2(Ge), and TiB2 diffusion barrier coatings were metallurgically characterized on

several different refractory metals for their ability to protect the metallic substrate from detrimental

effects of C and Si diffusion and reactions. Samples were characterized in the uncoated, as-coated,

and as-coated plus heat treated conditions, both with and without simulated C/C or C/SiC

processing. A discussion of the characterization of each coating on various refractory-alloy tubes

follows, with the results for the uncoated condition, wherever possible, serving as the baseline for

comparison. The discussion details the changes in physical and chemical nature of the coatings and

substrate microstructures in so far as these changes may ultimately affect the mechanical properties

of the refractory-metal tubes. The metallurgical characterization was limited to the study of Mo,

Mo-41Re, and Mo-47.5 Re tubes.
4

The appearance of the tubes after the PyC/SiC deposition was quite varied and appeared to be

significantly affected by the presence of a diffusion barrier coating. A photograph of an uncoated

AC Mo-41Re “D-shaped” tube (#17) after PyC/SiC deposition is shown in Fig. 1. The radius of

the curved portion of the tube is 0.3 in, and the tube segment is 1-in. long. The surface of the tube

was very rough, but in general, the PyC/SiC remained attached. This is in stark contrast to a TiB2

coated AC Mo tube with PyC/SiC deposited, shown in Fig. 2, where the surface was very smooth,

but the outer layer spalled. The “D-shaped” tube, again with a 0.3-in. radius, is shown in the

center of Fig. 2, and pieces of PyC/SiC layers that spalled are shown above and below the tube.

Fig. 1 Photograph of uncoated AC Mo-41Re “D-shaped” tube after PyC/SiC deposition (#17).

Fig. 2 Photograph of TiB2 coated AC Mo “D-shaped” tube after PyC/SiC deposition (#24).

The uncoated PM Mo-47.5Re surface appearance after PyC/SiC deposition (Fig. 3) looked

very different from the MoSi2(Ge) coated PM Mo-47.5Re surface (Fig. 4) after PyC/SiC

deposition. There were no large spalled sections in the uncoated condition (see Fig. 3), but a very

rough surface. The surface roughness due to the PyC/SiC on the uncoated PM Mo-47.5Re tube

(Fig. 3) was much larger than on the uncoated AC Mo-41Re tube (Fig. 1). As with the TiB2 coated

AC Mo tube with PyC/SiC deposition, the PyC/SiC on the MoSi2(Ge) coated Mo-47.5Re tube

spalled (see Fig. 4) and left a relatively smooth surface.

From the above photographs, it is clear that the PyC and SiC layers adhere to Mo and Mo-Re

surfaces in significantly different manners depending on whether the tubes are coated or not with a

diffusion barrier. Though the photographs indicate gross levels of adherence of the PyC and SiC
5

layers, they do not provide an indication of coating adequacy or inadequacy. On several coated

samples, a spalled layer was present. The uncoated AC Mo-41Re, though it had a very rough

surface, did not have significant spallation or flakes. Finally, it should be noted that the spalling of

the PyC/SiC would be very different if the tubes were embedded and processed in a fiber preform.

Fig. 3 Photograph of uncoated PM Mo-47.5Re tube after PyC/SiC deposition (#28). (Outer

tube diameter is 0.197 in.)

Fig. 4 Photograph of MoSi2(Ge) coated PM Mo-47.5Re tube after PyC/SiC deposition (#23).

(Outer tube diameter is 0.197 in.)

Materials and Coatings

The variables in this study were the refractory metals and alloys that were formed into thin

wall tube specimens, protective coatings, environment (constituent materials from C/C or PyC/SiC

processing), and the thermal exposure time and temperature. Each of these variables significantly

influenced the metallurgical condition and mechanical properties of the material.

A full factorial test matrix, attempting to study the interplay of all the various materials,

processes, and heat treatment variables, was not feasible for this study so only selected

combinations of the variables were evaluated. The scope of this investigation in relation to metals

and alloys, protective coatings, environment (C and/or SiC), and heat treatment time and temperature

used in the study, in given in Table 1. Table 1 indicates the conditioning of the tube (as received,

coated with PyC/SiC, embedded in C/C, and exposure for some time at a specific temperature, T(t)).

The refractory metals and alloys used in the metallurgical characterization study included arc

cast (AC)- and powder metallurgy (PM)- derived products. Both “D-shaped” and circular tubes

were evaluated. The different refractory metals evaluated were:

• AC Mo-41Re

• PM Mo-47.5Re

• AC Mo

The AC Mo and Mo-Re tubing was drawn by Thermo Electron Tecomet in Wilmington, MA and

the PM Mo-47.5Re tubes were fabricated by Rhenium Alloys, Elyria, OH.

The following three diffusion barrier coatings were evaluated:

• TiB2

• MoSi2(Ge)

• R512E
6

The TiB2 coating was applied by Applied Technologies Coatings, Fort Worth, TX, using the CVD

process. The MoSi2(Ge) coating was applied by Dr. Robert Rapp at Ohio State University, using

the pack cementation process. MoSi2(Ge) is essentially a MoSi2 coating doped with Ge to improve

the coefficient of thermal expansion (CTE) compatibility when applied to certain refractory metals.

The R512E coating is an Fe, Si, Cr-rich coating, previously utilized on Mo-41Re embedded in C/C1,

and was applied by Hitemco, Bethpage, NY using a slurry process.

The various refractory metal tubes were evaluated in the following conditions:

• uncoated (no diffusion barrier coating)

• coated (TiB2, MoSi2(Ge), or R512E)

• PyC/SiC deposition

• embedded in C/C

• exposed for 200 hours at 2000°F in a vacuum after PyC/SiC deposition (both coated with a

diffusion barrier coating and uncoated)

• exposed for 0.5, 2.5, 8, and 32 hours at 2300°F in a vacuum after R512E coating and

embedding in C/C.

Table 1. Material and Processing Conditions Characterized

 Refractory Metal Tube 

AC Mo-41Re PM Mo-47Re AC Mo

Uncoated

as received

as received + T(t) as received as received

PyC/SiC PyC/SiC PyC/SiC

PyC/SiC + T(t) PyC/SiC + T(t)

C/C + T(t)

Coated

MoSi2(Ge) as received as received as received

PyC/SiC PyC/SiC

PyC/SiC + T(t) PyC/SiC + T(t) PyC/SiC + T(t)

R512E as received

C/C

C/C + T(t)

PyC/SiC

TiB2 as received as received

PyC/SiC PyC/SiC

PyC/SiC + T(t) PyC/SiC + T(t)

R512E Coating

The Hitemco R512E coating on AC Mo-41Re was characterized in several different

conditions by scanning electron microscopy (SEM) and microhardness. The Mo-41Re was

characterized in the following conditions, with the specimen number in parenthesis following the

conditioning:

• Uncoated (#5)

• Uncoated, embedded in C/C, and exposed at 2300°F for 8 hr (#57)

• As coated with R512E (#50)

• Coated with R512E and embedded in C/C (#51)

• Coated with R512E, embedded in C/C, and exposed at 2300°F for 0.5 hr (#52)

• Coated with R512E, embedded in C/C, and exposed at 2300°F for 2.5 hr (#53)
7

• Coated with R512E, embedded in C/C, and exposed at 2300°F for 8 hr (#54)

• Coated with R512E, embedded in C/C, and exposed at 2300°F for 32 hr (#55)

Back scatter electron (BSE) images were obtained together with x-ray elemental maps to

identify morphology and chemistries of the coating and substrate. The BSE images rely on the

atomic number contrast in order to reveal details in a microstructure. Microhardness depth profiles,

using a 25 g load in conjunction with a Knoop indenter, were performed to corroborate the

chemistry information and to estimate the depth of contamination in the substrate. The

contamination in the substrate was due to elemental diffusion consequent to the coating process

and/or the 2300°F heat treatment.

Mo silicides

Mo, Si

Fe, Cr rich

Mo-41Re

Mo, Si

Si rich

a

b

c

d

e

Fig. 5 SEM micrograph of R512E-coated AC Mo-41Re in the as-coated condition (#50).

The R512E coating was composed of at least five layers, as shown in Fig. 5. The layered

structure resulted from the coating process step that involved heating the coated substrate at

approximately 2600°F for 2 hr. X-ray elemental maps of the coating are shown in Fig. 6, where the

light colored dots represent a high concentration of the specified element. Layer ‘a’ in Fig. 5 (~4-

μm thick), immediately in contact with the Mo-Re substrate, was substantially rich in Mo and Si,

and appeared to be composed of a mixture of Mo3Si and Mo5Si3. Layer ‘b’ (~10-μm thick),

adjacent to layer ‘a’, was generally rich in Mo and Si, and had a hardness of 1307±94 KHN.

Layer ‘c’ (~25-μm thick) was substantially enriched in Si, and had a high hardness of 2261±336

KHN. Layer ‘d’ appeared to be similar in chemistry to layer ‘b’. Layer ‘e’, with a hardness of

1424±147 KHN, was predominantly Fe and Cr. The elemental x-ray maps for this layer indicated

that compositional gradients existed with respect to both Fe and Cr. The scatter in the hardness

values for these various layers stems from chemical inhomogeneities, and thus attested to the

complex nature of the layered structure of the coating. 

In contrast to the inherently complex R512E coating layer, the Mo-Re substrate was

predominantly a single phase microstructure with a grain size of ~40 μm and a hardness of 415±26

KHN. The hardness scatter in the substrate is indicative of hardness differences due to grain-tograin differences in orientation of Mo-Re crystals. The presence of inter-dendritic shrinkage

cavities in the general microstructure appeared to be intrinsic to the process of arc casting the alloy

billets prior to their fabrication into Mo-Re tubes. It is pertinent to note that the sharp edges of

these inter-dendritic cavities are potential stress risers, and hence their presence is detrimental to a
8

ductility-based mechanical property such as burst strength. In other words, in the absence of these

shrinkage cavities, the strength of uncoated Mo-Re is likely to be significantly higher.

Mo silicides layer

Mo Re

Fe Cr Si

10 μm

Fig. 6 Elemental x-ray maps of R512E-coated AC Mo-41Re (#50).

Mo carbide 

reaction layer,

 ~ 65-μm thick

Fig. 7 SEM micrograph of uncoated AC Mo-41Re with C/C deposition (#57). Inter-dendritic

shrinkage cavities can be seen in the Mo-Re substrate.

The BSE images of the uncoated Mo-41Re tube embedded in C/C revealed the presence of a

~65-μm thick Mo-carbide region in the substrate adjacent to the C/C interface (see Fig. 7). This

transformed region, with a lamellar morphology, had a hardness value in the range of 1760 to 2430
9

KHN. In order to rationalize the development of this hard region in the uncoated sample, it is

instructive to view the binary Mo-C phase diagram (see Fig. 8) for plausible interactions of C with

Mo in the substrate. The isotherm corresponding to the 2300°F heat treatment temperature typical

of the C/C deposition process is superimposed on the figure. The phase diagram indicates the

presence of a eutectoid reaction at 2192±36°F for compositions beyond 5.7 wt.% C. The

decomposition products on cooling through a eutectoid transformation temperature are typically

associated with a characteristic lamellar morphology. The eutectoid constituent in the Mo-C system

is comprised of a mixture of αMo2C and MoC. The carbides are hard and brittle. In the absence

of a coating, therefore, substantial C enrichment of the substrate can take place rapidly during the

2300°F heat treatment, thus rendering a significant portion of the normally ductile cross-section

brittle through carbide formation. The presence of such a brittle layer can rapidly degrade the burst

strength. Beyond the carbidic region, the hardness profile in the substrate rapidly leveled off to the

Mo-41Re base hardness of 412±30 KHN (which incidentally, was the same as for the R512E

coated condition discussed previously).

2300 ¡F

Fig. 8 Mo-C binary phase diagram.

Hardness depth profiles of the Mo-41Re substrate for several different conditions are shown

in Fig. 9. Note that the origin of the hardness profile is at the substrate/C/C interface (for the

uncoated sample) or at the substrate/R512E interface (for the coated samples). The uncoated

sample showed a carbide region over the first ~65 μm depth (see Fig. 7). The thickness of the Mocarbide region, as determined from the BSE image, was fully corroborated by the plateau region of

the depth profile in Fig. 9. This plateau region had high hardness relative to the substrate, and can

therefore be expected to be less ductile than the Mo-Re substrate. All of the coated samples, on the

other hand, indicated no significant increase in hardness of the substrate (with measurements

starting below the Mo-silicide layer). It is thus apparent that the R512E coating did significantly

prevent C from diffusing into the Mo-Re substrate, but at the cost of a high hardness Mo-silicide

layer formation.

The hardness of the substrate varied slightly with coating and exposure time as shown in Fig.

10. The Mo-41Re in the uncoated condition, with a substrate hardness of 474±36 KHN, had a

microstructure similar to that of the Mo-Re substrate in the R512E-coated sample. The inherently

lower hardness of the coated sample compared to that for the uncoated condition suggests that the

coating process (together with the firing at 2600°F) brings about a thermally induced softening of
10

the Mo-Re substrate as shown in Fig. 10. The moderate but measurable increase in the substrate

hardness for the coated and exposed samples (with respect to that for the as-coated condition)

suggests small levels of possible Si contamination.

100

1000

0 20 40 60 80 100 120 140

#50:as-coated

#51:as-coated+C/C as densified

#52:as-coated+C/C+1/2hr@2300¡F

#53:as-coated+C/C+2.5hrs@2300¡F

#54:as-coated+C/C+8hrs @2300¡F

#55:as-coated+C/C+32hrs@2300¡F

Microhardness,

 KHN 25gf

Distance into Mo-41Re substrate, x, μm

Uncoated, embedded in

C/C, heated 8 hrs at 2300°F

Mo-carbide

region

x

Mo carbide

reaction layer,

 ~ 65-μm thick

x

Fig. 9 Microhardness depth profile of coated and uncoated Mo-41Re substrate.

Increasing the time of the heat treatment affected both the coating characteristics and the MoRe substrate. The samples coated with R512E, unlike the uncoated sample embedded in C/C,

showed no carbidic region. The presence of the R512E coating thus appeared to prevent C from

diffusing into the Mo-Re substrate. For all the heat treatment times, the R512E coating on the

2300°F exposed samples embedded in C/C showed qualitatively the same layered structure as for

the as coated condition discussed above. However, the thickness of the Mo-silicide layer

(corresponding to Mo3Si and Mo5Si3) increased with increasing exposure time, as shown in Fig.

11. Correspondingly, there was a decrease in the Si-enriched layer. This clearly suggests that with

increasing duration of the 2300°F exposure, diffusion gradients were set up such as to cause

depletion of Si from the Si-rich layer and a concomitant growth of the Mo-silicide layer on the

substrate. The growth of this brittle silicide layer was at the expense of the Mo-Re substrate.

In contrast to a thickness of ~65 μm of the carbide layer that developed in the uncoated

sample during 8 hr of heat treatment at 2300°F, only ~5.5 μm of the Mo-silicide layer developed

during the same duration for the coated sample. The hardness depth profiles in the substrate

indicated no significant concentration gradients due to Si diffusion. However, the moderate but
11

measurable increase in the substrate hardness with respect to that for the as-coated condition

suggests (see Fig. 10) small levels of possible Si contamination.

200

240

280

320

360

400

440

480

Microhardness , KHN 25gf

uncoated

R512-coated

R512-coated + C-C as-densified

coated + C-C + 0.5hr@2300¡F

coated + C-C + 2.5hrs@2300¡F

coated + C-C + 8hrs@2300¡F

coated + C-C + 32hrs@2300¡F

diffusion of small levels of Si

softened condition

as-received condition

520 

Fig. 10 Microhardness of AC Mo-41Re substrate.

0.0

5.0

10.0

15.0

Mo silicide

thickness,

μm

0.0 8.0 16.0 24.0 32.0

Time, hr

Mo silicides {

Si rich

Fig. 11 Mo-silicide layer thickness in AC R512E-coated Mo-41Re as a function of time at

2300°F .

An SEM micrograph of an R512E-coated Mo-41Re sample embedded in C/C and heated for

0.5 hr at 2300°F in a vacuum is shown in Fig. 12. It is readily seen that both lateral and transverse

cracks developed in both the R512E coating and the Mo-silicide layer. These cracks tended to

grow with increasing time of heat treatment. Thus, eventually the coating became ineffective in
12

preventing C diffusion into the substrate. In addition, the silicide layer also tended to debond, thus

increasing the thermal resistance to heat flow.

Mo-41Re substrate

R512E coating with lateral

and transverse cracks

Fig. 12 SEM micrograph of R512E-coated AC Mo-41Re embedded in C/C and heated for 0.5

hr at 2300°F (#52).

MoSi2(Ge) Coating

A Ge-doped MoSi2 coating was applied to AC Mo-41Re, AC Mo, and PM Mo-49.5Re by

Dr. Robert Rapp of Ohio State University. The MoSi2(Ge) coating was applied by a pack

cemetation process. After the coating was applied, PyC/SiC was deposited on some of the tubes

by HACI. An analysis of the Mo and Mo-Re alloys, with and without the coating, follows.

AC Mo-41Re

AC Mo-41 Re tubes with and without the MoSi2(Ge) coating were characterized in the

following conditions:

• Uncoated (#5)

• Uncoated with PyC/SiC deposition (#17)

• Uncoated with PyC/SiC deposition and exposed at 2000°F for 200 hr (#31)

• MoSi2(Ge)-coated (#13)

• MoSi2(Ge)-coated with PyC/SiC deposition (#15)

• MoSi2(Ge)-coated with PyC/SiC deposition and exposed at 2000°F for 200 hr (#16)

On the uncoated sample with PyC/SiC deposition, a Mo-carbide reaction layer formed

adjacent to the Mo-Re substrate. It is thus clear that in the absence of a protective barrier coating,

the PyC/SiC processing was associated with undesirable substrate/PyC/SiC reaction products. The

Mo-carbide layer was brittle and thus represented a loss of effective load-bearing cross section for

the substrate. Further, the thickness of the Mo-carbide layer increased from an initial 5 μm to 20

μm after the 2000°F/200 hr heat treatment (see Fig. 13). In contrast, on the coated sample (see Fig.

14) the MoSi2(Ge) coating appeared to be effective in preventing carbon from the PyC/SiC layer

from diffusing into, and reacting with, the Mo-Re substrate. As a result, no Mo-carbide layer was

observed. On the other hand, Si from the MoSi2(Ge) coating appeared to react with Mo in the MoRe substrate forming a Mo-silicide layer. Further, the thickness of this reaction layer increased
13

with heat treatment time. As a result of the Mo-silicide layer, and perhaps because of its CTE

incompatibility with the substrate, extensive cracks developed at the coating/substrate interface in the

2000°F/200 hr heat treated sample.

Mo-Carbide layer, ~ 20-μm thick

Si-rich layer, ~2.5-μm thick

Mo-Re substrate

Fig. 13 SEM micrograph of uncoated AC Mo-41Re with PyC/SiC deposition and heated at

2000°F for 200 hr (#31).

Si depleted layer, ~ 10-μm thick

Si rich layer Crack

Coating

Substrate

Fig. 14 SEM micrograph of MoSi2(Ge)-coated AC Mo-41Re with PyC/SiC deposition and

heated at 2000°F for 200 hr (#16).
14

PM Mo-47.5Re

PM Mo-47.5Re samples with and without MoSi2(Ge) coatings were characterized in the

following conditions.

• Uncoated (#3)

• Uncoated with PyC/SiC deposition (#28)

• Uncoated with PyC/SiC deposition and heat treated in a vacuum at 2000°F for 200 hours

(#29)

• Coated with 17 μm MoSi2(Ge) (#2)

• Coated with 17 μm MoSi2(Ge) with PyC/SiC deposition (#23)

• Coated with 17 μm MoSi2(Ge) with PyC/SiC deposition and heat treated in a vacuum at

2000°F for 200 hours (#34)

100

1000

0 20 40 60 80 100 120 140

Hardness, KHN25g

Distance, mμ

Mo-carbides region

Mo-47.5Re+PyC/SiC+heat treated @2000°F

Mo-47.5Re+MoSi2

Mo-47.5Re+MoSi2+PyC/SiC+heat treated@2000°F

Fig. 15 Hardness depth profiles as a function of depth into the substrate for various

conditions of PM Mo-47.5Re.

In the uncoated condition, the hardness of the Mo-47.5Re substrate was 445±35 KHN.

Microhardness depth profiles in the substrates for various conditions are presented in Fig. 15.

With reference to Fig. 15, it is clear that without a diffusion barrier coating, the 2000°F/200 hr heat

treatment produced a 50-60-μm thick, hard, and brittle Mo-carbide layer between the substrate and

the PyC/SiC layer. The x-ray elemental maps for this condition are presented in Fig. 16. With

respect to the SiKα map therein, it is evident that ~14-μm thickness of the Mo-Re substrate is

affected by Si diffusion. Of this, the first ~6-μm thickness appears to be comprised of a Mo-
15

silicide layer with a distinct mottled appearance (see the corresponding BSE image), and the

remaining 8-μm thickness a Mo-Re-Si solid solution. The presence of the original PyC layer is

evidenced by the vertical strip of high carbon intensity on the extreme left side of the CKα map. The

ubiquitous presence of carbon in the CKα map also attests to substantial carbon diffusion in the

Mo-Re substrate as a result of the heat treatment. 

In contrast to the uncoated condition, the MoSi2(Ge) barrier coating effectively prevented the

Mo-carbide formation. The substrate hardness for the various conditions are plotted as a bar graph

in Fig. 17. Even when the sample was MoSi2(Ge) coated with no PyC/SiC deposition, the nominal

substrate hardness was substantially higher than the uncoated condition. The elemental x-ray maps

for this sample are presented in Fig. 18. The Si x-ray map therein shows a ~10-μm thick Sidepleted region within the original MoSi2(Ge) layer at the coating/substrate interface; no perceivable

Si diffusion into the substrate is apparent. The substantial increase in the substrate hardness for

this condition is therefore likely to be due to rather copious amounts of the χ (MoRe) phase,

brought about by the high-temperature effects associated with the coating process. The presence of

the Si-depleted layer within the MoSi2(Ge) layer also implies that MoSi2(Ge) is unstable when it is

in contact with the Mo-47.5Re substrate at the coating temperatures. In addition, cracks parallel to

the coating/substrate interface were apparent within the coating. 

Mo

Si

BSE image

10 μm C

SiC

 layer

Mo-carbides

(~60 μm)

Substrate

PyC

Mo-silicide

Mo-Re-Si

Fig. 16 Elemental x-ray maps for uncoated PM Mo-47.5Re with PyC/SiC deposition and heat

treated at 2000°F for 200 hr (#29).

In comparison with the uncoated condition, it is seen that the Mo-47.5Re substrate hardness

was markedly higher whenever the samples were embedded in PyC/SiC (and subsequently heat

treated) whether they were MoSi2(Ge) coated or not (see Fig. 17). When the tube was embedded in

PyC/SiC but not coated, the hardness increase appeared to be related to the diffusion of both
16

elemental Si and C from PyC/SiC into the substrate. By the same token, when the sample was

coated with PyC/SiC subsequent to MoSi2(Ge) coating, the hardness increase stemmed essentially

from Si diffusion into the substrate. This is amply demonstrated in Fig. 19, which shows the x-ray

maps for the MoSi2(Ge)-coated and the PyC/SiC-embedded sample, subsequent to the 2000°F/200

hr heat treatment. It is clear from this figure that the original 17-μm thick MoSi2(Ge) coating

totally disappeared in addition to migration of large amounts of Si away from the PyC/SiC into the

substrate. The substrate appeared to be substantially enriched in Si. Also, a ~10-μm thick, freshly

formed Mo-silicide layer was observed as shown in Fig. 20. 

For the 2000°F/200 hr heat treated condition, Re-rich precipitates formed either at the grain

boundaries or within the alloy matrix (see Fig. 20). As per the Mo-Re phase diagram, this alloy

was in the two-phase region with (Mo) and χ as the equilibrium phases. (Mo) is a Mo-Re solid

solution while the χ phase is an intermetallic with the ReMo composition. The 2000°F/200 hr

heat treatment appeared to readily favor the precipitation of the χ phase, thus conforming the alloy

to equilibrium conditions of alloy stability. Grain boundary precipitation of intermetallic phases,

in general, may have a detrimental effect on eventual ductility properties of the substrate, regardless

of whether the substrate is coated or uncoated.

300

350

400

450

500

550

600

650

Hardness, KHN25g

Mo-47.5Re

uncoated MoSi2 coated

MoSi2 coated+PyC/SiC embedded+heat treated

uncoated+PyC/SiC embedded+heat treated

Fig. 17 Substrate hardness of Mo-47.5Re in various conditions.
17

Mo Si

Re BSE image

Si-depleted layer

MoSi2(Ge)

substrate

10 μm

Fig. 18 Elemental x-ray maps for MoSi2(Ge)-coated PM Mo-47.5Re (#2).

Mo-silicides

 layer

substrate

PyC/SiC

C Si

Mo BSE image

10 μm

substrate substrate

substrate

Fig. 19 Elemental x-ray maps for MoSi2(Ge)-coated PM Mo-47.5Re with PyC/SiC deposition

and heat treated at 2000°F for 200 hr (#34).
18

~10-μm thick Mo-silicide layer

Precipitates with a high

concentration of Re

Coating

Substrate

Fig. 20 SEM micrograph of MoSi2(Ge)-coated PM Mo-47.5Re with PyC/SiC deposition and

heat treated at 2000°F for 200 hr (#34).

Substrate

Mo-carbide layer

Fig. 21 SEM micrograph of uncoated AC Mo with PyC/SiC deposition (#20).

AC Mo

AC Mo tubes with and without a MoSi2(Ge) coating were characterized in the following

conditions.

• Uncoated (#10)

• Uncoated with PyC/SiC deposition (#20)
19

• MoSi2(Ge) coated (#11)

• MoSi2(Ge) coated with PyC/SiC deposition (#26)

• MoSi2(Ge) coated with PyC/SiC deposition and heat treated at 2000°F for 200 hr (#27)

The BSE image for the uncoated AC Mo sample with PyC/SiC deposition (see Fig. 21)

showed that significant levels of porosity existed in the original PyC/SiC layer. Also, a ~5-μm

thick Mo-carbide layer is evident in the Mo substrate at the original substrate/PyC/SiC interface. In

the AC Mo sample coated with MoSi2(Ge) (see Fig. 22), large transverse cracks are seen in the

~50-μm thick MoSi2(Ge) layer. Minute pores, intrinsic to the coating process, are also evident in

this layer. The presence of interdendritic shrinkage cavities in the Mo substrate attests to the metal

deficit experienced during solidification, subsequent to the arc casting of the Mo tubes. No

interaction products are evident, however, at the substrate/coating interface. 

Substrate

Coating

Substrate

Shrinkage

cavity

Fig. 22 SEM micrograph of MoSi2(Ge)-coated AC Mo in as coated condition (#11).

The SEM image together with the x-ray elemental maps for Si, C, and Mo are presented in

Fig. 23 for the MoSi2(Ge)-coated sample with PyC/SiC deposition. Note that for this condition, a

~10-μm thick Mo-silicide layer is present at the substrate/coating interface. No Mo-carbide layer

developed in the substrate, however. Large transverse cracks appear to traverse the entire thickness

comprising the PyC/SiC + Mo-silicides layers, thus compromising its mechanical integrity. From

the CKα map it is apparent that these cracks served as paths for C diffusion into the substrate

through the intervening SiC layer of the PyC/SiC deposition layer. With respect to the MoKα map,

it is clear that the PyC/SiC deposition layer is contaminated with Mo. The original MoSi2(Ge)

layer in contact with the PyC/SiC deposition may therefore not be fully chemically stable. This may

result in part of the MoSi2 disintegrating into Mo and Si. Mo may migrate to the PyC/SiC layer

while Si may diffuse into the Mo substrate. From the SiKα map, a minute amount of Si appears to

have diffused into the substrate. Subsequent to the 2000°F/200 hr heat treatment (see Fig. 24), the

transverse cracks formed within the original PyC/SiC layer appear to be wider together with a

thicker (~20 μm) Mo-silicide layer formation at the original substrate/MoSi2(Ge) interface. A

thicker Mo-silicide layer in the 2000°F/200 hr heat treated condition compared to that in the un-heat

treated condition suggests that the prolonged heat treatment promoted further deterioration of the

SiC layer, and the chemical reaction of Si with the Mo substrate. Furthermore, a mild Mo

enrichment of the SiC layer is also apparent. The CKα elemental map once again indicates that C

from PyC diffused through the transverse cracks into the Mo substrate thereby promoting the
20

formation of a Mo-C solid solution layer in the substrate. The 2000°F/200 hr heat treatment thus

had an adverse effect on the overall physical and chemical stabilities of both the coatings and the

substrate. 

Si C

Mo SEM image

SiC

Substrate

MoSi2 layer

10 μm

Fig. 23 Elemental x-ray maps for MoSi2(Ge)-coated AC Mo with PyC/SiC deposition in the

as coated condition (#26).

Nominal hardness variations for different conditions of the Mo substrate are presented in Fig.

25. With respect to the uncoated condition, the hardness of the MoSi2(Ge) coated substrate was

significantly lower, indicating possible softening of the substrate due to high-temperature effects

associated with the coating process. Subsequent to the deposition of PyC/SiC (deposition

temperature of ~ 2000°F) and the further heat treatment at 2000°F for 200 hours, the substrate

hardness increased, possibly due to minor alloying with C and/or Si. The CKα x-ray map in Fig.

24 does indicate a mild contamination of the substrate through C diffusion. The hardness data thus

appear to be sensitive enough to show minor changes in substrate chemistry.
21

Si C

Mo BSE image

Mo

MoSi2

SiC

C

10 μm

Fig. 24 Elemental x-ray maps of MoSi2(Ge)-coated AC Mo with PyC/SiC deposition heat

treated at 2000°F/200 hr (#27).

200

250

300

350

400

Hardness, KHN25g

M o Mo+MoSi2

Mo+MoSi2 +PyC/SiC

Mo+MoSi2 +PyCC/SiC+heat treated

Fig. 25 Substrate hardness of AC Mo in various conditions.
22

TiB2 Coating

Samples of arc cast Mo and PM Mo-47.5Re were coated with a TiB2 coating using the CVD

process. The coatings were applied by the Applied Technology Coatings (ATC), Fort Worth, TX.

AC Mo

Samples of AC Mo, with and without TiB2 coating, were characterized in the following

conditions:

• Uncoated (#10)

• Uncoated with PyC/SiC deposition (#20)

• TiB2 coated (#4)

• TiB2 coated with PyC/SiC deposition (#24)

• TiB2 coated with PyC/SiC deposition and heat treated at 2000°F for 200 hr (#30)

Mo-B-rich phase

Ti-rich phase

Mo substrate

B Ti

Mo 10 μm BSE image

Fig. 26 X-ray elemental maps for TiB2–coated AC Mo in the as-coated condition (#4).
23

Ti B C

Si Mo SEM image 10 μm

Ti-rich phase

Mo substrate

Si-rich particle

B4C

Si-rich area

Fig. 27 X-ray elemental maps for TiB2-coated AC Mo with PyC/SiC heat treated at 2000°F

for 200 hr (#30).

The effects of the TiB2 coating on Mo may be summarized by Fig. 26 and Fig. 27. From

Fig. 26, it is clear that the TiB2 coating was chemically unstable in contact with the Mo substrate at

the coating temperature and below. The site of the original TiB2 coating was replaced by a Ti-rich

layer, which was largely depleted in B. Furthermore, a thin layer of a Mo-B-rich phase is apparent

at the substrate/coating interface. With reference to the binary Mo-B phase diagram, the possible

Mo-B-rich phases for increasing B concentrations are Mo2B, αMoB, Mo2B5 or MoB4,

respectively. Subsequent to PyC/SiC deposition followed by the 2000°F/200 hr heat treatment, the

chemistries of the coating/substrate were further altered drastically as shown in Fig. 27. The oneto-one correspondence in the B and the C elemental maps suggests a strong affinity for B in the

TiB2 coating with C in the PyC/SiC deposition. In addition, a thin layer of Si-rich particles with a

mottled appearance is seen at the coating/substrate interface. This layer is perhaps comprised of

Mo-silicide particles. Thus, it is clear that TiB2 as a coating material for Mo suffered from severe

chemical incompatibility problems. 

PM Mo-47.5Re

PM Mo-47.5Re with and without the TiB2 coating was characterized in the following

conditions:

• Uncoated (#3)

• Uncoated with PyC/SiC deposition (#28)

• Uncoated with PyC/SiC deposition and heat treated at 2000°F for 200 hr (#29)

• TiB2 coated (#1)

• TiB2 coated with PyC/SiC deposition (#25)

• TiB2 coated with PyC/SiC deposition and heat treated at 2000°F for 200 hr (#35)

Just as in AC Mo-41Re and AC Mo with no diffusion barrier coating, a Mo-carbide layer

was also apparent in the uncoated PM Mo-47.5Re samples that were embedded in PyC/SiC. This

condition is not further discussed here since it was discussed previously.
24

The TiB2-coated condition (#1) was distinguished by a smooth TiB2/substrate interface and a

coating thickness of ~30 μm (with a hardness of ~4800±340 KHN). A few transverse cracks were

also apparent in the coating. The Mo-47.5Re substrate showed 10-15-μm wide grains. Some of

the grains contained fine-sized Re-rich precipitates (probably the χ(ReMo) phase). EDS spectra

indicated that B from the TiB2 coating diffused into the substrate to a depth of ~15 μm even during

the process of coating with TiB2. The x-ray elemental maps for this condition are presented in Fig.

28. With reference to this figure, it is pertinent to note that unlike the TiB2-coated Mo, the TiB2-

coated Mo-47.5Re alloy was substantially more stable and also distinguished by the absence of

Mo-B-rich phases at the coating/substrate interface. In contrast to unalloyed Mo, the presence of

47.5%Re in Mo thus appeared to confer superior chemical stability on the alloy vis-à-vis the TiB2

coating.

Ti B

Mo BSE image

10 μm

TiB2

Mo-Re substrate

Fig. 28 X-ray elemental maps for TiB2–coated PM Mo-47.5Re (#1).

Changes in the chemistries of the Mo-47.5Re substrate, the original TiB2 coating, and the

PyC/SiC deposition, consequent to the 2000°F/200 hr heat treatment are shown in Fig. 29, with the

substrate on the left side of each image. In the region of the TiB2 and PyC/SiC layers, the x-ray

maps of B and C correspond to each other on a one-to-one basis. This implies that a significant

amount of B from the original TiB2 coating combined with C from the PyC/SiC deposition to form

boron carbide (perhaps B4C). The presence of boron carbide is apparent throughout the TiB2 and

PyC/SiC layer. From the B-C binary phase diagram, formation of B4C needs at least 9.8 wt.% C

and this appeared to be readily available from PyC/SiC during the 2000°F/200 hr heat treatment.

The substrate enrichment with B (as indicated by the B x-ray map) accounted for the rest of the

boron in the original TiB2 layer. The corresponding x-ray map for Ti and the BSE image clearly
25

indicate the B4C regions formed islands within the network of Ti-rich regions. The Si x-ray map

together with the BSE image indicate that Si from SiC (in PyC/SiC) not only formed a ~10-μmthick Mo-silicide layer in the substrate at the substrate/coating interface, but also diffused

appreciably into the substrate. The dispersion of both B and Si from the composite layer of TiB2

and PyC/SiC into the substrate thus implies severe chemical instabilities for both TiB2 and the

PyC/SiC in the presence of each other and hence the inadequacy of TiB2 as a protective barrier.

Although TiB2 appears to have succeeded in preventing Mo-carbides formation, the significant

problem of inherent incompatibility between TiB2 and PyC/SiC still exists.

Fig. 29 Elemental x-ray maps in TiB2-coated PM Mo-47.5Re with PyC/SiC deposition and

heat treated at 2000°F for 200 hr (#35).

In the uncoated condition, the Mo-47.5Re substrate had an average hardness number of

445±35 KHN. In the TiB2-coated condition, the substrate had a moderately elevated hardness of

467±38 KHN, with no evidence of an interface reaction product layer consequent to the coating

process. However, a mild hardness gradient with respect to the coating/substrate interface was

observed in the substrate. The presence of a mild hardness gradient together with an observed

increase in the substrate hardness for the coated condition implies that B from the TiB2 may have

diffused into the substrate. Subsequent to PyC/SiC deposition, followed by the 2000°F/200 hr heat

treatment (#35), the substrate hardness increased appreciably. With reference to Fig. 29, it is clear

that for this alloy condition, significant levels of B (from TiB2 coating) and Si (from PyC/SiC

deposition) were present in the substrate, thus rendering it harder. The substrate hardness values

for the various conditions are represented in Fig. 30. 

The following general observations may be made with regard to the influence of chemistry

differences in Mo-Re alloys on precipitate microstructures, as predicated by heat treatments. In the

Mo-47.5Re alloy sample subjected to the 2000°F/200 hr heat treatment (in contrast to the Mo-41Re

alloy heat treated in the same manner), regardless of whether the alloy was coated with either TiB2

or MoSi2(Ge), Re-rich precipitates formed both within the alloy and at grain boundaries. Based on

the Mo-Re equilibrium phase diagram, the 47.5%Re alloy (unlike the Mo-41Re which was
26

essentially a single phase solid solution of Mo and Re) was in the two phase region, (Mo) + χ.

The 2000°F/200 hr heat treatment likely hastened the precipitation of equlibrium amount of the

χ(ReMo) phase. Grain boundary precipitation of these intermetallic precipitates may have

detrimental effects on strength and ductility of the substrate. The substantial increase in the

substrate hardness of the TiB2 coated with PyC/SiC deposition and heat treatment (see Fig. 30)

may be a result of appreciable levels of B, and also the χ precipitates present in the substrate. This

observation once again points to inherent instability of the TiB2 coating and also the potential

degrading influence of the χ precipitates on the ductility properties of the substrate.

300

350

400

450

500

550

600

650

Hardness, KHN25g

Mo-47.5Re

uncoated TiB2 coated

TiB2 coated+PyC/SiC embedded+heat treated

Fig. 30 Variation in substrate hardness of PM Mo-47.5Re.

Burst Tests

Burst tests were performed on several of the refractory-metal tubes with and without diffusion

barrier coatings, and prior to and after PyC and SiC deposition by the simulated HACI C/SiC

process. The tubes were taken through the entire HACI deposition process, but were not embedded

in a carbon fiber preform. Due to nonavailability of samples, burst tests on tubes embedded in C/C

were not conducted. 

The tubes were burst tested at room temperature using pressurized water. Swage lock fittings

were used to seal the tubes at each end. In several cases, the tube diameters were not standard, and

the fittings had to be drilled out to fit the tubes. During testing, the pressure was increased in ~250

psi intervals every 30 sec. There were no long hold times at any pressure. Fig. 31 shows

photographs of the test setup and a Mo-Re tube both a) before and b) after burst testing. An
27

attempt was made to test tube lengths greater than several diameters, but that was not always

possible. 

a) b) 

Fig. 31 Photograph of the test apparatus and tube, a) prior to and, b) after burst testing.

The results of the burst tests are shown in Table 2. The specimen number, diffusion barrier

coating (if any), environment (PyC/SiC processing), time at temperature excluding coating and

processing exposure, burst pressure, and failure stress are shown in the table for each tube material.

The outer diameter (OD) and inner diameter (ID) are also given for each tube. In all cases, the

failure stress, σ, was calculated from

σ = P r / t

where P is the burst pressure, r is the average of the inner and outer radii of the tube, and t is the

tube wall thickness. In the measurement of r, contribution to thickness included any protective

coating such as TiB2 and MoSi2(Ge), as well as the PyC/SiC deposition. Several of the tubes broke

during handling and were not tested; while during other tests, fittings leaked or failed during testing

thus limiting the pressures attained. For cases where the tube did not burst, the burst pressure and

failure stress are listed in Table 2 preceded by “>”. For each test condition, only one tube was

tested. Thus, the burst strength evaluation may be considered exploratory and no claims to

statistical significance are made. 

The PM Mo-47.5Re tubes provided the most complete burst test data. These tubes were

tested in the uncoated and coated (TiB2 and MoSi2(Ge)) conditions both before and after the

PyC/SiC deposition. All of the tube lengths (distance between sleeve on fittings) were greater than

1 in. except for specimen #23 (0.603-in. long), specimen #25 (0.188-in. long), and specimen #29

(0.605 in.). From the burst test data, there appears to be a slight degradation in the tube strength

due to the application of the TiB2 and MoSi2(Ge) coatings. In addition, the PyC/SiC deposition

process appeared to degrade the strength of each of the tubes. However, the uncoated tube strength

was only reduced 8.7% by the PyC/SiC deposition, versus strength reduction of 17.2% for the

MoSi2(Ge) barrier coating and 22.5% for the TiB2 barrier coating. For the uncoated tube with

PyC/SiC processing, exposure at 2000°F for 200 hours reduced the burst strength only slightly

from the as-processed state. 

The AC Mo-41Re tubes had a 0.030-in. wall thickness, and required large pressures to burst.

As a result, in some cases the burst pressure exceeded the test capability. The length of the AC

Mo-41Re tubes were 0.99 in. for specimen #12, 0.986 in. for specimen #13, 0.394 in. for specimen

#14, and 0.658 in. for specimen #15. Due to testing limitations, specimens #12, #13, and #14 did

not burst. However, an increase in diameter of 0.001 in. was measured after the tests. Specimen

#15, which was an AC Mo-41Re tube with a MoSi2(Ge) coating, failed at approximately 130 ksi.

The burst strength of specimen #16, which also included exposure at 2000°F for 200 hours, was

significantly degraded versus the as-processed condition (#15). Though it was difficult to isolate

the contributions of tube material processing (AC versus PM) and chemistry (41%Re versus
28

47.5%Re) on the AC and PM tube strengths, it did appear that the AC tubes were significantly

stronger than the PM tubes.

Table 2. Results of Burst Tests

Spec. Coating, environment, Burst Failure

 No. and thermal exposure pressure, psi stress, ksi

PM Mo-47.5Re, OD = 0.197 in., ID = 0.175 in.

3 uncoated 11,509 97.3

28 PyC/SiC 10,507 88.8

29 PyC/SiC and 200 hr at 2000°F 10,200 86.2

1 TiB2 11,073 93.6

25 TiB2, PyC/SiC 8,577 72.5

2 MoSi2(Ge) 11,315 95.7

23 MoSi2(Ge), PyC/SiC 9,368 79.2

AC Mo-41Re, OD = 0.188 in., ID = 0.128 in.

12 uncoated > 60,700 > 158

14 PyC/SiC > 48,000 > 126.4

13 MoSi2(Ge) > 47,300 > 124.5

15 MoSi2(Ge), PyC/SiC 49,500 130.3

16 MoSi2(Ge), PyC/SiC 42,000 110.6

and 200 hr at 2000°F

CVD Re, OD = 0.195 in., ID = 0.185 in.

6 uncoated > 5946 > 113

18 PyC/SiC broke prior to testing

CVD Nb, OD = 0.204 in., ID = 0.186 in.

7 uncoated 2101 22.8

19 PyC/SiC broke prior to testing

AC Mo, OD = 0.188 in., ID = 0.137 in.

20 PyC/SiC 15,565 49.6

The uncoated CVD Re tube withstood a stress of 113 ksi without failure while the uncoated

CVD Nb tube failed at a stress of 22.8 ksi. Both the Re and Nb tubes with the PyC/SiC deposition

broke prior to testing, and hence no meaningful trends could be established. The burst stress of

49.6 ksi is significantly below the ultimate of 109 ksi for Mo, indicating that the Mo was

detrimentally affected by the PyC/SiC deposition. In the case of the Re, Nb, and Mo burst test

specimens, only a single datum point was available, and therefore a comparison of the different

conditions is not possible.

Despite the limited number of conditions evaluated and the lack of replicates, the burst test

data indicated the following:

• Although the effects of the barrier coatings and PyC/SiC deposition on the AC Mo-41Re

could not be determined fully due to testing limitations, the MoSi2(Ge) coating appeared to

substantially degrade the burst strength relative to the uncoated condition.

• Both TiB2 and MoSi2(Ge) coatings followed by a PyC/SiC deposition, moderately degraded

the burst strength of PM Mo-47.5Re relative to uncoated PM Mo-47.5Re.

• AC Mo-41Re appeared to be significantly stronger than PM Mo-47.5Re. 
29

Concluding Remarks

Metallurgical characterization (microhardness survey and SEM examination) and burst tests

were performed on several different refractory metals and alloys, with the major effort directed

toward Mo and Mo-Re alloys. The following observations can be made:

• For the PM Mo-47.5Re alloy, burst strength degraded moderately when the alloy tubes were

coated with either MoSi2(Ge) or TiB2 (relative to the uncoated condition). The degradation in

burst strength was more drastic subsequent to the coating followed by the PyC/SiC deposition,

with contribution to degradation from TiB2 being more severe than from MoSi2(Ge).

MoSi2(Ge) thus appears to be more effective than TiB2 as a protective barrier coating for PM

Mo-47.5Re. 

• The AC Mo-41Re alloy appears to have significantly higher burst strength (though not enough

samples were available for a statically significant test) than the PM Mo-47.5Re alloy, when the

comparisons are made with respect to uncoated, MoSi2(Ge) coated, PyC/SiC deposited, or

MoSi2 coated + PyC/SiC deposited conditions. Mo-41Re is essentially a single phase alloy

whereas Mo-47.5Re is a two phase alloy. The presence of an appreciable volume fraction of the

χ(ReMo) phase at the grain boundaries appeared to be a ductility degrading factor, and may

explain the poorer burst strength of PM Mo-47.5Re compared to Mo-41Re.

• The lamellar morphology of the carbide layer evident in the AC Mo-41Re processed with C/C,

implies that more than 5.7 wt.% C reacted with Mo (in the Mo-Re substrate) to yield a eutectoid

mixture of αMo2C and MoC. This Mo-carbide layer was hard and brittle, and may severely

degrade burst strength.

• The Fe, Si, Cr-rich R512E coating, sandwiched between the Mo-41Re substrate and the C/C

layer, very effectively prevented formation of Mo-carbides in the substrate. A Mo-silicide layer

formed in the substrate, however, due to the substrate-coating interactions. This silicide layer,

though slower growing than the carbide layer (in the uncoated condition), did increase in

thickness with heat treatment time, thereby degrading the useful engineering cross-section of the

Mo-Re substrate. The development of lateral and transverse cracks in the original R512E

coating may diminish the efficacy of the coating.

• Even though coatings seemed to reduce carbide formation, they did so with the creation of a thin

silicide layer which seemed to be detrimental to the burst strength.

• TiB2 was inherently unstable when it was in contact with the Mo substrate, with B possibly

tending to form the Mo2B phase at the coating/substrate interface. In contrast, TiB2 appeared to

be substantially more stable in contact with the Mo-47.5Re substrate. 

Acknowledgements

The authors would like to thank the Metals and Thermal Structures Branch at NASA Langley

Research Center for funding this work under Contract No. NAS1-96014. The valuable assistance

of Mr. Jim Baughman of Analytical Services & Materials, Inc. in the SEM studies of the

investigation and Honeywell Advanced Composites, Inc. in the PyC and SiC deposition is

gratefully acknowledged.
30

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34

Appendix: Literature Survey

The issues regarding material compatibility between ceramics and refractory metals have been

a major focus of research over the past several years. In this regard, the need for diffusion barrier

coatings for preventing the formation of deleterious carbides and silicides has been well

documented. Advanced high-temperature cooling applications such as heat-pipe-cooled leading

edges may often require the elevated-temperature capability of carbon/silicon carbide (C/SiC) or

carbon/carbon (C/C) composites in combination with the hermetic capability of metallic tubes. In

view of this, a critical review of past research efforts in the area of refractory metals incompatibility

appears to be germane to identifying 1) potentially useful metal systems and 2) diffusion barrier

coatings. Mo-Re alloys particularly are deemed potentially viable materials as heat-pipe-cooled

leading edges in hypersonic vehicles. Compatibility of Mo-Re with silicon carbide and carbon is

therefore a specific area of interest. By discussing previous research efforts on material

incompatibilities, it is hoped that a clearer understanding would emerge as to which class of

materials and coatings holds promise.

Refractory Metal Compatibility with Carbon and Silicon Carbide

Most refractory metals form carbides when in the presence of carbon at high temperatures.

Nb, Ta, and Mo readily absorb C.2-3 Carburization is expected for Mo above 2012°F, for W above

2552°F, and for Ta above 1832°F. The solid solubility of C in Mo varies between 0.005 and

0.0922 wt. % in the temperature range of 3002-3992°F,4 with an activation energy, Q, for C in Mo

being 163 Btu/mole in the temperature range of 2192-2912°F.5 Re is the only refractory metal that

does not form a carbide.6-7 However, solubility of carbon in solid Re increases with temperature,

resulting in excellent bond strength between carbon and Re. Re absorbs about 0.9 wt. % C when

heated in methane at 1472-3992°F.

Lundberg8 studied the interaction of SiC with W in the temperature range of 2916-3276°F.

An electron microprobe analysis revealed the formation of W5Si3 and W2C layers between the SiC

deposition and the W substrate, i.e., a layered structure was formed comprised of SiC-W5Si3-

W2C-W. The two reaction layers appeared to grow at the same rate with no W diffusing into the

SiC. Similarly, in graphite that was CVD coated with W and SiC a W2C reaction layer was found

between the W and the graphite, with no W in the graphite. From these studies, it was determined

that SiC and C react with W at similar rates. The reaction layer growth could be predicted by

x = 4.03 x 104 (t) exp(-113,400/T) 

where x is the reaction layer thickness in inches, t is the heating time in seconds, and T is the

absolute temperature in degrees Rankine. As evident from the equation, reaction layer growth is

highly temperature dependent.

Ranken and Lundberg9-10 studied several different refractory metals for use in refractorymetal-lined ceramic tubes for heat pipes. The ceramics they chose for study were SiC and Al2O3.

They looked for a refractory metal with a similar CTE as that of the ceramics. They determined that

Nb and Al2O3 have nearly identical CTE’s. In addition, W and either α- SiC or SiC containing 5-

10 % free silicon (KT-SiC), have close CTE’s. The CTE of β- SiC (the form typically obtained by

CVD) can be fairly well matched by alloys of W and Mo. Similarly, CTE compatibility between

CVD Niobium and Al2O3 was demonstrated by the absence of defects related to debonding and

thermal stresses. Due to the uncertainty about the reaction between W and SiC, tests were

performed on SiC with a CVD layer of W. Upon heating to 2916°F, three phases were revealed in
35

the reaction layer,W5Si3, W6Si2C, and W6SiC2. Si could not be detected in the W nor W in the

SiC.

Merrigan and Keddy11 studied SiC tubes with a CVD W liner for use in heat pipes for waste

removal. SiC was used because it has excellent thermal shock resistance, good thermal

conductivity, low volatility, and resistance to both oxidizing and reducing atmospheres. In addition,

SiC has an extremely low hydrogen permeability and its CTE is similar to that of W. Previous

investigations have shown the W-SiC reaction layer growth followed a parabolic law of the form x2

= 2Kt, where x is the layer thickness, K is the reaction rate constant, and t is time.

Rovang, et al.,12-13 reported on the development of a C/C heat-pipe internally coated with a Nb

liner for the SP-100 project. Their approach was to use an interlayer of 0.039-0.197 mil (1-5 μm)

thick Re (applied by CVD) between the C/C and the Nb. This interlayer would: a) provide a

modest gradation of the CTE mismatch, b) assume a portion of the load from the induced stress, c)

improve the coating adhesion, and d) serve as a carbon diffusion barrier to allow the Nb to remain

ductile.

Yaney and Joshi14 studied the reaction between Nb and SiC. A Nb layer, 0.0394-0.0787 mil

(1-2 μm) thick, was sputter deposited on SiC substrates and heated to temperatures ranging from

1472 to 2323°F. An Auger electron spectroscopy depth profile revealed extensive reaction between

the Nb and the SiC with the typical reaction layer sequence being

SiC/Nb5Si4C/Nb5Si3/Nb2C/NbO/Nb. Joshi, et al.,15 studied the reaction of Nb and Ta with SiC

and Al2O3 up to 2192°F. Nb and Ta were sputter deposited on SiC and Al2O3 substrates to a

thickness of approximately 0.0394 mil (1 μm). The main thrust of their study related to the Nb/SiC

system. They observed that the reactions of Nb with SiC was much more uniform for the single

crystal SiC than for the polycrystalline SiC. Also, annealing the Nb/SiC system at 2012°F for 4 hr

resulted in reaction of the entire Nb film. In contrast, there was no evidence of reaction in the

Nb/Al2O3 system after heating at 2192°F for 4 hr, indicting that Al2O3 may be a good diffusion

barrier between Nb and SiC.

Fries, et al.,16 evaluated the diffusion of carbon into W and Mo. They found the diffusion of

carbon through the carbide phase as the rate-determining step. Though a thin WC skin was

observed, only the W2C layer growth was prominent below 4352°F. The diffusion of W into

carbon was negligible compared to that of C into W. Similarly, carbon diffusion into Mo was

found to proceed rapidly, and the metal quickly became saturated with carbon.16

Reagan, et al.,17-18 constructed tubes with a trilayer structure of SiC, graphite, and W. While

the inside of the graphite tube was coated with W through CVD deposition, the outside of the tube

was coated with CVD-deposited SiC. The mechanical integrity of the resulting structure was

assessed by performing thermal shock, thermal cycling, and pressure tests. In one series of tests,

one end of the structure was heated by radio frequency induction heating to 3853°F in air and then

water quenched to below 1963°F. This process was repeated 10 times, with the structure remaining

leak-tight. In another series of tests, the structure was cycled from 981°F to 2916°F in air by rf

induction heating 150 times. Again, the structure remained leak tight. Next, the structure was

cycled from 1000°F to 3177°F in air 102 times. A vacuum of 10-6 Torr was maintained inside the

structure. During the last few cycles, a crack was noticed in the SiC, and testing was stopped to

prevent oxidation of the W. The structure was also pressurized to 500 psi for several hours and

remained leak tight.

The feasibility of coating C/C tubes internally with refractory metals for use as heat pipes was

studied by Hartenstine et al.19 Re was used as a carbon diffusion barrier coating between the C/C

and the refractory liner (either W or Nb). Four different coating techniques were studied. A 0.002
36

in. thick layer of carbon was applied between the C/C and the Re to transition the CTE between

them and thereby reduce the interfacial thermal stresses. Though Re does not form a stable carbide

at their use temperature of 2804°F, carbon was found to diffuse through the Re and react with the

W or the Nb substrates, to form a W-carbide or a Nb-carbide. As the coatings were applied at a

temperature of 1832°F or higher, the most severe thermal stresses were experienced during the

cooling of the systems. Re/W and the C/Re/W coatings cracked severely during thermal cycling

due to the embrittlement of the W as a result of carbon diffusion through the Re. The Re/W

coating separated from the C/C wall, while the C/Re/W coating did not. The Re/Nb and C/Re/Nb

coatings completely separated from the wall to form an almost stand-alone tube with only minor

cracking. Again, the carbon diffused through the Re, this time resulting in Nb-carbides. 

Rabin20 reviewed the interaction of SiC with various metals (Ni, Fe, Ti, Cr, Al, Cu, Pd, Mo, W,

V, Ta, Nb, Hf and Zr) with relevance to SiC joining. In the SiC/Mo system, Mo2C formed as a

result of C diffusion into the Mo, together with a small amount of Mo5Si3 in between the SiC and

the carbide layer. In yet another study, a Mo2C layer, a mixed layer of Mo2C + Mo5Si3, and a

layer of Mo5Si3C were found to grow between the Mo and SiC. In the SiC/W system, existence of

a mixed layer of W-carbide + W-silicide was noticed.

In graphite coated with Mo, growth of an Mo2C intermediate layer was observed between the

graphite and Mo21. This carbide layer had a parabolic growth rate with annealing time, at

temperatures as high as 1832°F. Cross-sections of the Mo coating revealed the Mo grains to be

columnar in structure, with grain size increasing with deposition temperature. The diffusion

coefficient of C in Mo was found to increase with decreasing columnar grain size; and

experimentally determined diffusion coefficients as well as carbide growth rates for a Mo coating

on graphite, and for a Mo sheet coupled to graphite, were documented. The study concluded that

for Mo coatings on graphite, grain boundary diffusion of C in the carbide layer played a significant

role in the carbide layer growth.

Diffusion couples of Mo and Si were used by Tortorici and Dayananda22 to study the growth

of silicides and interdiffusion in a Mo-Si system. The diffusion couples were heat-treated in either a

H2 atmosphere at temperatures between 1652°F and 2012°F, or in an Ar-5% H2 atmosphere at

temperatures up to 2462°F. Layers of MoSi2 and Mo5Si3 were observed in each case. The MoSi2

grains exhibited a preferred direction of growth parallel to the diffusion direction. Also, the MoSi2

layer exhibited a columnar microstructure and had one to two orders of magnitude greater thickness

than the Mo5Si3 layer. The interdiffusion coefficients of Mo and Si were also one to two orders of

magnitude greater than in both MoSi2 and Mo5Si3. No Mo3Si was found, however, despite the fact

the Mo-Si phase diagram allows its formation. Since Mo3Si was observed by other researchers to

have formed at temperatures above 2552°F, Tortorici and Dayananda suggest that barriers for the

nucleation of Mo3Si are perhaps overcome at higher temperatures.

Martinelli et al.23, studied diffusion couples of Mo and SiC. The Mo-SiC joints were made

by hot pressing at temperatures between 2192oF and 2732oF for times between 15 minutes and 2

hours. The reaction layer that formed consisted of a mixed layer of Mo5Si3 and Mo2C. Hotpressing for longer times at lower temperatures (i.e., 2 hr at 2372oF) resulted in a separate and

distinct layer of Mo2C being extended into the Mo region. On the other hand, on hot pressing at

2552oF for long times or at 2732oF, a layer of Mo5Si3C formed against the SiC interface. Shear

strengths of the joints as a function of hot pressing times of 15 min, 30 min, 1 hr, and 2 hr were

determined. At 2192oF, shear strength increased as bonding times increased. At 2372oF, shear

strength increased for times up to 1 hr and decreased thereafter for the 2 hr exposure. While at

2552oF, no specific trend was apparent, at 2732oF, shear strengths showed a continuous decline

with increasing hot pressing times. The authors attributed these trends in shear strength to the

variations of elastic moduli and the CTE of the Mo, SiC, and inter-layers. The Mo2C layer was said

to be compatible with SiC. Therefore strength increased with time at low temperatures where Mo2C
37

formed. In contrast, the SiC and Mo5Si3 layers, as well as the SiC and Mo5Si3C layers were

deemed mutually incompatible, and hence strength decreased for the higher temperature exposures

where these incompatible silicides formed23.

The phase equilibria of Si3N4 with Cr, Mo, W, and Re were studied by Schuster24, at

temperatures of 1832°F and 2552°F in an argon environment. In addition, he complemented his

study with a critical review of the available literature. In the Cr-Si-N system, Si3N4 reacted with

chromium at temperatures as low as 1652°F. Various chromium silicides, such as CrSi2, CrSi,

Cr5Si3, and Cr3S formed. In the absence of external nitrogen pressure, Cr2N did not form.

However, it did form when there was a sufficient nitrogen partial pressure resulting from the

decomposition of Si3N4 by chromium. In the Mo-Si-N system, in the absence of a nitrogen

pressure, MoSi2 and Mo5Si3 formed and co-existed with Si3N4 at 1832°F. Reactions between

Si3N4 and Mo were slow. At temperature of 2912°F and higher, however the co-existence between

Si3N4 and Mo5Si3 was disturbed and only MoSi2 co-existed with nitrogen gas. In the W-Si-N

system under an argon environment, reactions between tungsten and Si3N4 did not occur at the

temperature of 1832°F. At a higher temperature of 2143°F, however, reactions caused the

formation of nitrogen and W5Si3; and at a still higher temperature of 2732°F, nitrogen also coexisted with WSi2. In the Re-Si-N system, Si3N4 did not react with Re in argon at the temperature

of 1832oF. However, at 2552°F, Re17Si9 and nitrogen gas formed. In assessing the viability of

Si3N4 as a diffusion barrier coating in these metallic systems, Schuster concluded that the

formation of the various embrittling silicides, together with the undesirable chemical reactions

occurring during thermal cycling, restrict the potential of such joints for structural applications at

elevated temperatures.

Carbon Diffusion Barriers

Potential diffusion barrier coatings for refractory metals in contact with carbon, appear to be

Pt, Ir, Re, SiC, TiB2 and TiC. Platinum forms low melting silicides at rather low temperatures, and

this restricts its use in a C/SiC composite.25 Carbon forms a eutectic with Pt at 3150°F; with Ir at

4177°F (but this temperature is significantly lowered in the presence of Si to at least 3902°F); and

with Re at 4471°F. It is thought that while Re does not form a carbide, carbon has been found to

diffuse through a Re coating and cause embrittlement of the substrate metal.19 However, the

diffusion of carbon through Re is temperature dependent, and is not thought to be significant up to

3000°F. The TiB2 coating, on the other hand, appears to prevent diffusion of carbon through it, but

its efficacy has not been tested at high temperatures and for extended time periods. TiC coating is

thought to be an effective diffusion barrier for carbon, but its utility is limited by its severe

brittleness.

Platinum, Iridium, and Rhenium

Interactions of Pt and Ir with carbon were studied by Whitcomb, et al.26 and Selman et al27.

Molten Pt (and others in its group) dissolves substantial amounts of carbon, which in turn

depresses the melting point of Pt. Pt and C form a eutectic at a temperature slightly below the

melting point of Pt. The eutectic solidification is characterized by the formation of a mixture

containing a near C-free Pt solid solution and graphite. Pt containing the eutectic graphite is

brittle.28 In contrast to most reports that the solubility of C in Pt is small26 (~1 wt.% max. at

1700°C)29, Pt dissolves substantial amounts of C after heating in contact with graphite powder for a

few hours even at temperatures as low as 2192°F.27 While the dissolved C in solid Pt has little

effect on the hardness of Pt, the dissolved C in Pd considerably increases the hardness of Pd. It is

also reported that at high temperatures, carbon diffuses very rapidly in solid Pt.27
38

Pt and Si form three eutectics over the entire compositional range with melting temperatures

ranging from 1497-1810 °F.30 Pt that contains sufficient Si to form the low melting eutectic phase

is brittle.

Chou31 studied the reaction between Pt and SiC at 1652°F and reported interfacial melting

between Pt and SiC due to a reaction product with a low melting point. The further annealing at

1652°F caused the reaction zone to penetrate into the SiC. All the SiC decomposed in that region

converted into a mixture of Pt3Si (with a melting point of 1526°F) and C. In another study,

extensive reaction between SiC and Pt were observed to take place in the temperature range of

1652-2012°F.32 The reaction products were graphite and Pt-silicides. Local melting followed by

solidification was observed at the interface, with formation of Pt-silicides. Due to its low solubility

in Pt-silicides, the carbon precipitated out of the Pt-silicide matrix.

Searcy and Finnie33 studied the ternary systems of Si and C with Re and the Pt metals. They

observed that neither the Pt metals nor Re formed stable compounds with C at temperatures up to

2912°F. Rhenium silicides were observed to be less stable than silicides of the Pt metals. 

 The CTE of Ir closely matches with that of Mo, W, and Ta, and can therefore be

recommended as a good oxidation diffusion barrier.34 Ir does not form a carbide at temperatures

up to 3812°F. The diffusion rate of carbon in Ir at temperatures up to 3452°F is low; but above this

temperature, the diffusion rate is high.

A strong bond can be formed between Ir and C by allowing molten Ir to wet the surface of a

graphite substrate.34 The bonding results from a dissolution of graphite in the molten Ir (at

~3992°F) with subsequent reprecipitation of the graphite upon solidification of the metal. The

graphite crystallites form an interlocking network with the Ir and the graphite substrate.

Interactions of Ir-base alloys with carbon were studied by Harmon.25 There is an Ir-C eutectic

at 4177°F (melting point of Ir is 4429°F), and this temperature is significantly lowered in the

presence of silicon to at least 3902°F.

The effectiveness of a Re coating as a diffusion barrier between a graphite substrate and a Mo

coating was investigated. The Re coatings were 1.5-20 μm while the Mo coatings were 10 μm

thick. Graphite samples coated with Mo, and with both Re and Mo, were annealed at temperatures

from 1652°F - 2192°F. It was found that the growth of the Mo2C interlayer depended greatly on

the thickness of the Re layer, with the reaction rate decreasing exponentially with increasing Re

layer thickness. In both the single layer coating of Mo and the double layer coating of Re/Mo, the

growth of the intermediate carbide layer obeyed a parabolic law with annealing temperature. At

lower temperatures, the Re layer was more effective in slowing down the growth rate of the carbide

layer. For the case of a 10 μm Re interlayer, the carbide layers were 4, 23, 56, and 69% as thick as

for the a single Mo layer, at annealing temperatures of 1652, 1832, 2012, and 2192°F respectively.

Grain size or thickness of its columnar structure also had an effect on the reaction rate of carbide

growth. The carbide layer increased in thickness with decreasing grain size of the Re layer. It is

believed that columnar grain boundaries in the Re layer act as fast diffusion paths for carbide layer

growth in the Mo layer35.

Arnoult and McLellan36 studied the solid solubility of carbon in several refractory metals

such as Rh, Ru, Ir, and Re. Metal foil samples in contact with graphite were annealed at

temperatures in the range of 1508°F-2293°F, until equilibrium had been reached. Carbon content

was determined by a combustion C-analysis method using NBS (National Bureau of Standards)

standards for calibration. The solubility of carbon in the four metals was low and increased with

temperature. Solubility is defined as the atom ratio of carbon to metal. Carbon was the least

soluble in Ir, followed by that in Rh, Ru, and Re, with carbon solubilities in Ir being several times
39

lower than for the other metals, with the difference being predominant at high temperatures. In

rhenium, the solubility of carbon ranged from 8.52E-4 at 1508oF to 170.0E-4 at 2293oF. The

solubility of carbon in rhodium ranged from 8.70 x 10-4 at 1508oF to 75.0 x 10-4 at 2293oF with a

peak of 118.0 x 10-4 occurring at the intermediate temperature of 2156oF. In the carbon-iridium

system, the solubility of carbon ranged from 3.28 x 10-4 at 1508oF to 25.2 x 10-4 at 2293oF. The

solubility of carbon in ruthenium ranged from 9.95 x 10-4 at 1508oF to 125.0 x 10-4 at 2293oF. 

Hamilton et al37. characterized diffusion-related changes in samples taken from a Mo thruster

manufactured by CVD. The thruster consisted of a Mo mandrel and had a 50 μm Ir coating

followed by a 2 mm Re coating on it. The Mo mandrel was etched away after the thruster was

fabricated. In an un-annealed sample taken from the thruster, a 3 μm Ir-Re interdiffusion layer, and

two interdiffusion layers of IrMo and Ir3Mo (2-3 μm in total thickness) were observed. In addition,

adjacent to the Ir/Re interface, a 40 μm region of the Re coating, containing small voids was also

evident. X-ray diffraction analysis of the un-annealed sample showed the average grain size in the Ir

coating to be 10 μm while that in the Re layer to be ~400 μm. From three annealed samples

(2552°F for 14 hours, 3092°F for 8 hours, and 3452°F for 8 hours), it was determined that no

significant grain growth occurred during extended anneals at temperatures up to 2452°F. However,

additional void growth was observed for the sample annealed at 2452oF. The width of the

interdiffusion zone in which Re diffused into Ir increased with temperature and had widths of 10

μm, 20 μm, and 40 μm at temperatures of 2552°F, 3092°F, and 3452°F respectively.

Titanium Diboride

TiB2 coatings on Mo-Re, to serve as diffusion barriers for use between Mo-Re and C, were

developed by the Advanced Technology Coatings, Ft. Worth, TX. Results of the study on

interdiffusion of carbon were reported for Mo-Re coated with carbon, Mo-Re coated with TiB2/C,

and Mo-Re coated with TiB2/C/SiC.38 These alloy samples had not been heated except during the

CVD process. Results of this study are summarized as follows. In the Mo-Re sample coated only

with C, Mo diffused into the C coating, but no C diffused into the substrate. In addition, no cracks

were detectible at the coating-substrate interface. In the Mo-Re sample coated with TiB2/C, no C

was detected in the substrate; also, no cracks were visible at the coating-substrate interface. Next to

the substrate, a Ti-rich layer was seen. Neither C or Mo was observed in the TiB2 layer. Diffusion

of Mo from the substrate into the C was not apparent. Detection of B in the inter-layers and also

the substrate was extremely difficult. In the Mo-Re sample coated with TiB2/C/SiC, the outer

surface showed only Ti. It was believed that the C and SiC layers had broken off during

metallographic sample preparation. No C or Ti was detected in the substrate, however.

Molybdenum Disilicide

Govindarajan et al39, investigated coating systems that had the potential to provide oxidation

protection and prevent the C and Si diffusion into a Mo substrate. They pointed out the problems

associated with using MoSi2 as a diffusion barrier. The CTE of MoSi2 does not match that of a

Mo substrate (CTE at 1832oF for MoSi2 is reported to be 5.2 x 10-6 F-1 while for Mo it is 3.7 x 10-6

F-1). In addition, MoSi2 exhibits low creep strength at high temperatures, with the diffusion of Si

into Mo promoting the embrittling silicides.39 A multilayer coating of Mo and Si was used to

alleviate the CTE mismatch. The reasoning was that silicon, in excess of that necessary for forming

stoichiometric MoSi2, would minimize the formation of silicides. However excessive silicides did

form.40 Efforts at matching the CTE of Mo by using a layer of MoSi2 + 1.96 moles SiC was not

beneficial, either. Annealing at 1832oF caused the formation of Mo3Si and Mo5Si3 with a resulting

deleterious change in the CTE of the silicides layer.41 Based on this work, a new coating system
40

was developed. The new coating system consisted of a Mo substrate that was coated with a

diffusion barrier (based on the Mo-Si-C-N quaternary system) to prevent the diffusion of C and Si

into the substrate. On top of the diffusion barrier layer, a layer of MoSi2 + 1.96 moles SiC was

applied. This was followed by a thick layer of B- and Ge-doped MoSi2. The doping helps to lower

the viscosity of the glassy, self-healing oxide (SiO2) that forms on the surface as the MoSi2 layer

oxidizes. This silica scale serves as a diffusion barrier for oxygen. The diffusion barrier applied

onto the Mo substrate was chosen based on promising results by a research group at Los Alamos

National Laboratory and the Technical Research Center of Finland in which MoSi2Nx prevented the

formation of silicides up to 1652oF (the highest temperature investigated).42-46 The diffusion barrier

used by Govindarajan et al.39, was reactive sputtered onto the Mo substrate. It was comprised of

Mo-Si-C-N, and had an undetermined stoichiometry of MoSixNyCz. Annealing was conducted in a

vacuum at 1832oF (the highest temperature investigated) for 30 minutes. Silicides and carbides did

not form in the Mo substrate. The effectiveness of the diffusion barrier layer in preventing the

diffusion of C and Si was attributed to its amorphous nature. A powder form of the Mo-Si-C-N

diffusion barrier as well as a sample of Mo foil coated with the Mo-Si-C-N diffusion barrier layer

and the MoSi2 + 1.96 moles SiC composite layer were stable up to temperatures of 2300oF in an

oxidizing environment. The authors39 suggest that a promising area of investigation would be the

amorphous layers of nitrided refractory metal silicides in the form of M5Si3Nx, where M represents

Ta, W, Mo, etc. It was recommended that Ta5Si3 would be a good silicide to nitride since it is one

of the best Si diffusion barriers among these silicides, which are more effective Si diffusion barriers

than the MSi2 silicides.47

Overview

Incompatibility between Mo-Re and other refractory materials, in the form of joints and

coatings, can degrade the mechanical properties of Mo-Re. Two major issues of incompatibility

include chemical incompatibility and the CTE mismatch. While chemical incompatibility can lead

to embrittlement of Mo-Re through formation of carbides and/or silicides, the CTE mismatch can

lead to a poorly protected or an exposed Mo-Re substrate with inadequate protection against a

hostile oxidizing service environment.

The literature review conducted hitherto has facilitated identification of compatibility issues,

classes of materials that poorly match with Mo-Re, and those materials that hold promise as

diffusion barriers. As a general rule, all sources of C and Si that lead to formation of detrimental

carbides and silicides in Mo-Re should be avoided. However, a silicon-bearing system such as the

M5Si3Nx may still be worth considering as a diffusion barrier.47 Ceramic coatings of nitrides or

borides are also an area worth exploring. Nitrogen has been observed to be compatible with several

refractory metals.23 Potentially useful nitrides are ZrN, TiN, HfN, and BN. These nitrides have

fairly high melting points (2980°C, 2930°C, 2852°C, and 2300oC respectively), and have relatively

low densities compared to nitrides such as the TaN.48 Some work has been done with TiN

coatings on Mo and SiC. It appears that the relatively high CTE of TiN may preclude its use, unless

an appropriate intermediate layer that provides a transition in thermal expansions is used. 49-50

Possible useful borides are HfB2, TaB2, W2B2, ZrB2, and TiB2. A critical literature review of these

borides, with particular reference to their chemical and thermal expansion compatibilities, may help

narrow the scope of future developmental studies.

Issues of concern in targeting elements from the Pt group as possible diffusion barrier

coatings are density, CTE, temperature capability, and material compatibility. Even though a certain

metal does not react with C and/or Si, it may not necessarily render itself as a suitable candidate

material for diffusion barrier coating. For example, Re is not very reactive with C or Si, but C and

Si can easily diffuse through the Re layer into the Mo substrate, and form substrate Mo carbides

and/or silicides.
41

From the Pt group of metals, Os, Ir, and Ru (with respective density, CTE, and melting points

of 22.2 g/cm3, 2.8 ppm/oF and 5513oF for Os; 22.4 g/cm3, 3.6 ppm/oF and 4370oF for Ir; and 12.2

g/cm3, 3.6 ppm/oF and 4190oF for Ru51), are potential candidates as diffusion barrier coatings

between Mo-Re, carbon, and SiC. Ir is a viable diffusion barrier coating due to its proven track

record as a rocket thruster coating material, and also due to its favorable compatibility with Re52-54.

Pt and Pd, on the other hand, may not serve as useful candidate diffusion barrier coatings materials.

This is because they have relatively low melting points and relatively high CTE (with respective

density, CTE, and melting point of 21.4 g/cm3, 4.9 ppm/oF and 3222oF for Pt; and 12.0 g/cm3, 6.5

ppm/oF and 2826oF for Pd51). Also, carbides that form through interaction of Pt and C have even

lower melting points than Pt.
REPORT DOCUMENTATION PAGE Form Approved

OMB No. 0704-0188

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Washington, DC 20503.

1. AGENCY USE ONLY (Leave blank) 2. REPORT DATE

December 2001

3. REPORT TYPE AND DATES COVERED

Technical Memorandum

4. TITLE AND SUBTITLE

Effectiveness of Diffusion Barrier Coatings for Mo-Re Embedded in C/SiC

and C/C

5. FUNDING NUMBERS

WU 242-33-03-20

 

6. AUTHOR(S)

David E. Glass, Ravi N. Shenoy, Zengmei Wang, and Michael C. Halbig 

7. PERFORMING ORGANIZATION NAME(S) AND ADDRESS(ES)

 NASA Langley Research Center

 Hampton, VA 23681

8. PERFORMING ORGANIZATION

REPORT NUMBER

L-18116

9. SPONSORING/MONITORING AGENCY NAME(S) AND ADDRESS(ES)

National Aeronautics and Space Administration

Langley Research Center

Hampton, VA 23681-2199

10. SPONSORING/MONITORING

AGENCY REPORT NUMBER

NASA/TM-2001-211264

11. SUPPLEMENTARY NOTES

Glass: Langley Research Center, Hampton, VA; Shenoy, Lockheed Martin Engineering & Sciences, Hampton,

VA; Wang, Chromealloy Gas Turbine Corporation, Orangeburg, NY; Halbig, Glenn Research Center,

Cleveland, OH.

12a. DISTRIBUTION/AVAILABILITY STATEMENT

Unclassified-Unlimited

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13. ABSTRACT (Maximum 200 words)

Advanced high-temperature cooling applications may often require the elevated-temperature capability of

carbon/silicon carbide or carbon/carbon composites in combination with the hermetic capability of metallic

tubes. In this paper, the effects of C/SiC and C/C on tubes fabricated from several different refractory metals

were evaluated. Though Mo, Nb, and Re were evaluated in the present study, the primary effort was directed

toward two alloys of Mo-Re, namely, arc cast Mo-41Re and powder metallurgy Mo-47.5Re. Samples of these

refractory metals were subjected to either the PyC/SiC deposition or embedding in C/C. MoSi2(Ge), R512E,

and TiB2 coatings were included on several of the samples as potential diffusion barriers. The effects of the

processing and thermal exposure on the samples were evaluated by conducting burst tests, microhardness

surveys, and scanning electron microscopic examination (using either secondary electron or back scattered

electron imaging and energy dispersive spectroscopy). The results showed that a layer of brittle Mo-carbide

formed on the substrates of both the uncoated Mo-41Re and the uncoated Mo-47.5Re, subsequent to the C/C or

the PyC/SiC processing. Both the R512E and the MoSi2(Ge) coatings were effective in preventing not only the

diffusion of C into the Mo-Re substrate, but also the formation of the Mo-carbides. However, none of the

coatings were effective at preventing both C and Si diffusion without some degradation of the substrate.

14. SUBJECT TERMS

Coatings; Composites; Carbon diffusion barrier; Silicon

15. NUMBER OF PAGES

46

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