DOE/NASA/16310-10 

NASA TM-102122 

The Interface in Tungsten 

Fiber Reinforced Niobium 

Metal-Matrix Composites 

Toni L. Grobstein 

National Aeronautics and Space Administration 

Lewis Research Center 

Cleveland, Ohio 441 35 

September 1 989 

. Work performed for 

. Nuclear Energy 

U.S. DEPARTMENT OF ENERGY 

Reactor Systems Development and Technology 

Washington, D.C. 20545 

Under Interagency Agreement DE-A103-86SF16310 
ITABLE OF CONTENTS 

I. INTROOUCTION 

I I. BACKGROUND 

A. Fiber Reinforcement 

8. Diffusion 

C. Creep 

111. MATERIALS 

IV. 

V. 

A. Tungsten Fibers 

8. Matrix Materials 

PROCEDURES 

A. Fabrication 

8. Sample Preparation 

C. Creep Testing 

0. Sample Analysis 

RESULTS AND DISCUSSION 

A. Composite Creep Properties 

B. Creep Curves 

C. Multiple Linear Regression 

D. Interdiffusion Between Fiber and Matrix 

E. Effect of Matrix Alloy Additions 

i'i i 

PRECEDING PAGE BLANK NOT FILMED 

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25 

27 

27 

30 

31 

32 

37 

41 

53 
VI. 

VII. 

VI11 

F. Ki rkendall Void Formation 

G. Effect of Fiber Orientation 

CONCLUSIONS 

A. Fiber Reinforcement 

6. Fiber Orientation 

C. Matrix Alloying Additions 

0. 

E. 

REFERENCES 

APPENDICES 

A. Multiple linear Regression 

B. Vol ume Fraction Determination 

C. Rule of Mixtures Calculation 

0. 

Comparison to Monolithic Niobium Alloys 

Future Appl i cat i ons 

Extrapolation of Fiber Stress-Ruoture Data 
Advanced materials will play a major role in meeting the 

stringent weight, size, and performance requirements of future space 

power systems. 

include a service life of greater than 7 years at a temperature in 

The requirements for such a system, whicn may 

excess of 1350 K, in addition to resistance to liquid metal 

corrosion, dictate the use of refractory metals. The niobium-- 

1% zirconium (Nb-1Zr) allay has been suggested for use in such space 

power conversion applications, but does not have sufficient creep 

strength for the current design times and temperatures. Further, 

while current designs of space nuclear power systems specify niobium 

base alloys for reactor, heat pipe, and power conversion components, 

future appiications will need materials with greater hightemperature strength and increased creep resistance to meet mission 

requirements (Titran 1988). 

The goal of this research was to improve the creep 

strength/density ratio of the Nb-1Zr alloy without sacrificing 

desirable physical properties such as corrosion resistance to liquid 

alkali metals and low thermal neutron cross section. Fiber 

reinforcement has been shown to be an effective strengthening method 

(McDanels et a1 , 1960), and does not require large alloying 

1 
additions to the matrix. Therefore, the feasibility of using metalmatrix composites to meet the anticipated increased temperature and 

creep resistance requirements imposed by advanced space power 

systems is of interest. 

reinforced niobium (W/Nb) metal -matrix composites have been expiored 

by Trefilov et al. (1976) and Westfall et al. (1986), and 

prel iminary creep-rupture results were reported by Petrasek and 

Titran (1988). These studies demonstrated significant improvenents 

in the high-temperature strength/density ratio and the possibiiit;] 

for corresponding mass reductions in high-temperature space power 

systems. 

The tensile properties of tungsten fiber 

The intent of this study was to measure the creep 

properties of tungsten fiber reinforced niobium metal-matrix 

composites and to evaluate the effect of the fiber/matrix reaction 

on these properties. Two factors which may affect the fiber/matrix 

interface, and therefore, the creep properties also were evaluated; 

small additions of tungsten and zirconium to the niobium matrix, and 

variation of the orientation of the tungsten fibers relative to the 

principal stress axis. 

2 
BACKGROUND 

Fiber Reinforcement 

In a typical metal-matrix composite, the fibers reinforce and 

strengthen an otherwise weak matrix, while the matrix coats and 

protects brittle fibers which would otherwise fail due to environmental effects. Although it appears to be much simpler to coat a 

part made from the fiber material than to put fibers into a 

composite, fiber reinforcement is preferred over coating for t-,vo 

reasons. First, if even a small defect occurs in a coating, the 

entire core is exposed to the environment. In a fiber-reinforcei 

composite, however, a defect in the matrix exposes only a few 

fibers, and when they fail, the load is distributed over the 

remainder of the fibers, making the composite tough. 

fiber reinforcement is preferred versus coating is the increase in 

strength of fibrous over bulk materials. 

Another reason 

A composite differs from a monolithic material in that its 

physical and mechanical properties are defined as that of a 

structure rather than as a homogeneous or even a heterogeneous 

material. A composite does not have a characteristic modulus, 

tensile strength, or even density like a monolithic material does. 

These values depend on the volume fraction and distribution of 

3 
fibers. In a metal matrix composite with a strong interiacs gonc. 

:ne ;ensile and creep strength of the composite depends on the 

?roperties of the matrix, the fibers, and the load transfer betideen 

the two. 

This load transfer results in a very complex stress state. 

Factors such as residual stresses, mu1 ti axi a1 stresses, and 

thermally-induced stresses will impact the behavior of the 

composite. The thermal stresses are due to thermai expansion 

mismatch between the fiber and matrix. At 1500 K, the coefficient 

of linear thermal expansion for niobium is 8.75 x~O-~ K, approximately 40% higher than that for tungsten, 5.15 x~O-~. This means 

the matrix wants to elongate both axially and radially, althougn it 

is constra ned by the stronger fibers. 

structure tself can introduce additional sources for crack 

initiation or different crack paths, and can weaken conventional 

crack paths. All of these effects are somewhat minimized in this 

application since little or no thermal or mechanical cycling 

is involved, but they are not insignificant. 

In addition, the composite 

Maximum strengthening in a continuous fiber reinforced 

composite is achieved by orienting all of the fibers in the 

composite in a single direction, and then orienting the cornposit2 so 

the applied stress is along the fiber axis. However, since any real 

application will involve a multiaxial stress state, at least some 

4 
strength along more than one axis is required. 

fibers is an obvious method of boosting the transverse properties Of 

the composite. 

orientation on the internal stresses in the composite, especially in 

the interface, must be understood and quantified. 

Angle-plying the 

However, the effect of changing the fiber 

Another important factor is formation of a reaction zane 

between the fiber and matrix at elevated temperatures. Since the 

phase diagram for tungsten and niobium is a continuous solid 

solution (Figure l), interdiffusion in this system results in a 

continuous gradient betwee’n the composition of the fiber and that of 

the matrix across the interface. Obviously, the fiber/matrix 

reaction will affect and probably degrade the composite’s mechanical 

properties. 

composite material is to be used in real systems. If the 

degradation is substantial, we must look at ways to stop or slow the 

interdiffusion. 

We must be able to predict this degradation if this 

One way to impede interdiffusion in the composite is to create 

a diffusion barrier between the fiber and matrix. This technique 

was investigated by Veltri et al. (1975), who ion-plated therma1l:d 

stable coatings (such as Al2O3, Y2O3, and Hf02) onto tungsten wires. 

Unfortunately, when incorporated into a composite as coated fibers , 

small cracks in the coatings occurred due to the differential 

thermal expansion properties between the matrix, coating, and fiber. 

5 
It was discovered that these cracks in the coating allowed *ier:i fast 

diffusion of the matrix throughout the fiber/coating interface. 

leading to fracture of the fiber. 

3500 

3300 

3100 

0" 

W 

2 2900 

s 

z k2700 

2m 

23M 

ATOMIC PERCENT TUNGSTEN 

20 40 60 80 

I I I I I I 'I' 

I I I 

10 20 30 40 50 60 70 80 90 

KIGHT PERCENT TUNGSTEN 

Fig 1 Tungsten - Niobium phase diagram. 

6 
The feasibility of ion-implanted diffusion barriers has also 

been explored. Welsch et a1 (1987) investigated the concept that a 

layer of atoms on the fiber surface with considerably larger atomic 

radii than either the fiber or matrix material may impede interdiffusion. Although this technique shows promise (Kopp et ai, 1988), 

it is far from being applied. 

A different method is to make alloying additions to the 

matrix*. It is known that alloying niobium with tungsten decreases 

the niobium self-diffusion rate. Several investigation (Lyubimov 

et a1 1967; Lundy and Pawel 1969; Mundy et a1 1986; and Hehemann and 

Leber 1966) have found that increasing the tungsten content in the 

niobium decreases the diffusion kinetics exponentially across the W/Nb 

interface. It is possible, therefore, that alloying the matrix 

material with small amounts of tungsten will decrease interdiffusion 

rates in the composite. Also, the effect of zirconium in the matrix 

on the interdiffusion kinetics is of interest, given that Tuchinsky 

(1979) predicted that alloying niobium with zirconium would 

decrease its rate of diffusiou into tungsten. 

Di f fusi on 

Arcella (1974) measured interdiffusion zone growth in the W/Nb 

system, and found the empirical relationship 

*Alloying the fiber is impractical due to the complex chemistry 

which maintains the high-strength structure of the wire. 

7 
In ($1 = - 37390 - 3.869 T 

!Ahere AX is the width of the interdiffusion zone in cm, t is the 

time in seconds, and T is the temperature in Kelvin. He evaluated 

both arc-cast and CVD tungsten, both with relatively large grain 

size compared to that in a drawn tungsten fiber. 

Arcella found this equation quite accurate over a wide range 

of times and temperatures. 

light of the gross porosi$y which occurred on the niobium side of 

the interface in this system. 

across a diffusion interface, known as the Kirkendall effect, was 

first observed in the zinc/alpha brass system by Smigelskas and 

Kirkendall (1947). The Kirkendall effect occurs in a binary soiid 

solution where each of the two atomic species moves with a differenr: 

velocity. Experimental measurements have shown that the element 

with the lower melting point diffuses faster (Reed-Hill, 1973); in 

this case, the niobium. The net mass transfer results in a complex 

stress state at the interface. The area outside of the fiber which 

suffers a loss of mass is placed under a two-dimensional tensile 

stress, while the area just inside the fiber diameter is placed in a 

He noted that this was surprising in 

This phenomenon of net mass transport 

, state of two-dimensional compressive stess. The stress fields may 

8 
bring about plastic flow, with the resulting structural changes 

normally associated with plastic deformation and high temperatures; 

formtion of substructures, recrystall ization, and grain growth. 

When the two diffusing species move with greatly different 

rates during diffusion, voids or pores have been observed in that 

region of the diffusion zone from which there is a flow of mass. 

Since every time an atom makes a jump a vacancy moves in the 

opposite direction, an unequal flow in the two types of atoms just 

results in an equivalent flow of vacancies in the reverse direction 

to the net flow of atoms. 

The formation of voids, as a result of the vacancy current 

that accompanies the unequal mass flow in a diffusion couple, is 

influenced by several factors. It is generally believed that the 

voids are heterogeneously nucleated on impurity particles. The 

tensile stress that exists in the region of the specimen where the 

voids form is also recognized as a contributing factor in the 

development of voids. If this tensile stress is counteracted by a 

hydrostatic compressive stress placed on the sample during testing, 

the voids can be prevented from forming (Barnes and Mazey, 1958). 

Arcella found that the void structure, although presenting a 

reduced cross-sectional area to interdiffusion, did not appear to 

affect the extrapolation capabilities of his equation. He 

speculated that perhaps the reduced area for constituent flux was 

9 
counteracted by rapid surface diffusion effects within the voids 

Also, by the time the pores became gross, the interdiffusion 

kinetics were considerably reduced. 

Hehemann and Leber (1966) measured diffusion coefficients for 

thi s system using coup1 es prepared from polycrystal 1 i ne commerci a1 ly 

pure metals. 

in the temperature range between 1975 and 2575 K, decreased exponentially as the tungsten concentration was increased from 0 to 50 a%, 

then decreased abruptly at higher concentrations. 

coefficients measured were in the range defined by the selfdiffusion coefficients of the individual elements. However, the 

measured activation energy for interdiffusion was substantially less 

than that for self-diffusion in the pure components. It was 

concluded that short-circuit paths, structural irregularities which 

act as diffusion pipes, developed as a result of the diffusion 

process itself and contributed significantly to interdiffusion. 

relative contributions of volume and short-circuit diffusion 

appeared to change with temperature and rendered measurements of 

activation energies and frequency factors meaningless. 

observed Kirkendall porosity at all diffusion temperatures, but 

found it most severe at the lowest annealing temperature they 

evaluated, 1975 K. 

They found that the diffusion coefficients, measured 

The diffusion 

The 

They also 

10 
CREEP 

The Rule of Mixturgs 

- 

It is theoretically possible to predict the properties of a 

composite based on the properties of the individual components. The 

simplest method is the two bar model (or the rule of mixtures), 

given by 

where P is any physical or mechanical property, such as the elastic 

modulus or the yield stren,gth, and L is the volume fraction of each 

component of the composite. This equation is based on the 

assumption that the strains in the fiber and the matrix are equal to 

each other and to that of the composite. 

represented by x, include the matrix (m), the fiber (,f), and a 

reaction zone (z) formed between them, the equation is written as, 

If the components, 

This equation may be expanded to include different fibers or matrix 

materials in the same composite panel, and other components incorporated into the composite structure, such as cladding materials. 

To use this equation, we calculate the volume fraction of each 

component based on the initial conditions and interdiffusion 

kinetics. The properties of the fiber and matrix material can be 

11 
predicted through tests Of each component in isolation, and the 

properties of the interface region in the W/Nb system can be 

assumed to be somewhere between the two, since no intermetallic 

phases form. 

prediction of composite properties. 

Therefore, this simple equation can give a rough 

The rule of mixtures is most applicable when deformation is 

elastic and little or no interaction takes place between the 

components. Creep, however, involves primari ly inelastic behavior, 

and interaction between components is extensive at the emperatures 

and times of interest. McOanels et a1 (1967) has shown that the 

rule of mixtures can be applied to creep of W/Cu compos tes 

with certain modifications. 

Load transfer is critical in a composite in that the load is 

applied to the matrix and is effectively carried by the fibers. 

a ductile matrix composite with a good interface bond, the load is 

applied isostatically to the fibers by the surrounding matrix. 

However, the different properties of the fibers and matrix create a 

complex stress state during creep deformation. Stresses are induced 

by the large difference in the yield strengths of the tungsten fiber 

and the niobium matrix. This means the matrix wants to deform 

plastically under the applied stress, but is constrained by the 

fibers which remain elastic at the same stress level. 

In 

12 
Gregory and Rowe (1961) have shown that creep of niobium is a 

diffusion controlled process above 1200 K; and that the addition of 

1% zirconium retards polygonization and recrystal 1 ization during the 

1475 K standard aging treatment and leaves a forest of edge dislocations in the metal. This dislocation forest is stable to about 1375 

K, some 170 K above that for pure niobium. Two mechanisms appear to 

explain this strengthening effect of zirconium; first, ZrO precipitation in dislocations during aging, and second, the diffusion of 

zirconium atoms to edge dislocations, both of which pin jogs during 

creep, thus preventing dislocation climb. They also suggest that 

the presence of interstitial atoms such as oxygen, nitrogen, and 

carbon on the dislocations increases the binding energy of the 

zirconium atoms to the jogs. Ye et a1 (1985) confirms that much of 

the zirconium exists as a fine zirconia precipitate both within the 

grains and at the grain boundaries. 

2 

Qualitative evidence of grain boundary sliding is quite 

prominent in creep of niobium, although there is no evidence of void 

formation and intercrystalline cracking (Begley et al, 1968). At 

high temperatures, creep deformation is dominated by diffusional 

creep along grain boundaries and other fast diffusion paths. 

strain rates, the ease with which fold formation at triple points 

and grain boundary migration occurs seems to preclude stress conc2nAt low 

13 
trations arising from grain boundary sliding. Hence, creep fracture 

in niobium is transcrystall ine by void formation and coalescence, 

and the fracture elongation is quite high, 50 to 100% and more. 

Pure tungsten exhibits transcrystal 1 ine fracture with quite 

high ductility (Begley et al, 1968). Alloying with dispersed 

phases, such as thoria, however, promotes intercrystalline fracture 

which in turn leads to low creep ductility. Disparity in the shear 

strengths of the internal grain area and the grain boundary would 

tend to promote intercrystalline cracking if the strains due to 

grain boundary sliding cannot be relaxed by deformation within the 

grain. Precipitates might restrain grain boundary migration, and 

microcracks might form at the grain boundary due to low interfacial 

adhesion. However, grain boundary precipitates can exert a 

favorable influence by decreasing the free lengths of grain 

boundary, which would reduce stress concentrations. 

of doped tungsten wires has been observed to occur by nucleation of 

voids at potassium-filled bubbles (Briant, 1988). 

Creep failure 

The Continuum Approa~ 

Much of the previous work in this area has avoided this 

complex stress analysis and treated the composite as a structure -- 

the continuum approach. 

be predicted using a multiple linear regression model such as 

Larson-Mil ler (1952) or Orr-Sherby-Dorn (1954), regard1 ess of its 

The expected creep life of any material can 

14 
'- 

I* 

internal structure. These types of analyses use creep life data 

measured at several temperatures and appl ied stress levels to 

extrapolate predictions for other temperatures and stresses. 

difficult to use this approach with any composite due to varying 

volume fractions of fiber; although, if known, this can be taken into 

account. Also, in the case of W/Nb composites, an additional tirneand temperature-dependent interface reaction is taking place by 

interdiffusion between the fiber and matrix which may affect the 

strength of the composite. If this interdiffusion does affets the 

strength of the composite and is not take into account, interpretation of the regression analysis is impossible. 

It is 

15 
MATERIALS 

Tunqsten F1 berz 

When tungsten is drawn into a fine wire, the resulting microstructure may have up to five times the strength of the monolithic 

material (Yih and Wang 1979). The stabil i ty of the fibrous grain 

structure at elevated temperatures is dependent on the grainboundary pinning mechanism in the wire. The 100-hour creep rupture 

strengths of three tungsten-based alloy wires are compared to that 

for unreinforced (monolithic) Nb-1Zr in Figure 2, and micrographs of 

their longitudinal sections are in Figure 3. 

The 218CS tungsten wire is a potass um-doped, commerically 

available lamp filament. Potassium bubb es pin the grain boundaries, 

resulting in an increased recrystallizat on temperature over 

monolithic (undoped) tungsten (Welsch et a1 1979). The ST300 wire 

is dispersion strengthened with 1.5 w% thoria (-0.1 pm diameter). 

This wire is,slightly stronger than 218CS, due to the increased 

thermal stability of the thoria dispersoid compared to the potassium 

bubbles, and is also commerically available. The W-HfC wires are 

strengthened by a very fine dispersion (0.05 pm diameter) of even 

more thermally stable hafnium carbide particles, and exhibit 

superior creep Properties. Unfortunately, W-HfC wires were not 

16 
available at the time of this study; therefore, only ST300 tungsten 

fiber reinforced composites were evaluated. 

of 200 pm. 

The wire had a diameter 

40 8m 1480 K 

Nb-1 Zr 218 CS ST300 WHfC 

Comparison of the 100-hr creep rupture strength-to-density 

ratio for three tungsten-based alloy wires versus 

monolithic Nb-1Zr alloy. (From Tietz and Wilson, 1965; 

Petrasek, 1972; Petrasek and Beremand, 1987; and Petrasek 

and Signorelli, 1969.) 

17 

The matrix alloy compositions are listed in Table I. 

Additions of zirconium (1 and 2 a) and tungsten (1, 5, and 9 w%) 

were made to niobium to assess their effect on the fiber/matrix 

reaction. The small addition of zirconium results in a zirconia 

precipitate which strengthens the alloy and removes oxygen from the 

grain boundaries of the niobium*. Tungsten was added to decrease 

the concentration gradient between the fiber and matrix, and is a 

solid solution strengthener in niobium. 

The matrix materials were obtained in the form of 1.6 mrn 

diameter wire. 

tallization anneal and single tensile tests were conducted in vacuum 

for each composition at 1366 and 1589 K. 

tensile tests was .05 m/m. 

ratios for these alloys are given in Figure 4. 

Specimens of these alloy wires were given a recrysThe strain rate of the 

The ultimate tensile strength to density 

*This addition is important to lithium corrosion resistance. Lithia 

is one of the most stable metal oxides (more stable than 

zirconia), and if any oxygen is present in the grain 

boundaries, the liquid metal will leach it out. The lithium 

can eventually permeate the entire thickness of the container, 

resulting in "weeping" on the outer wall surface (Hoffman and 

Harrison, 1968). 

19 
Table I. Matrix Alloy Compositions 

Nb 

N b-1 Zr 

Nb-1Zr-5W 

Nb-2Zr-1 W 

1366 K 

1589 K 

0 5 10 15 20 

STRESS-TO-DENSITY RATIO (MPalglcc) 

Fig 4 Comparison of matrix alloy ultimate tensile 

strength-to-density ratios. (compositions in a) 

20 
PROCEDURES 

Fa br i cati on 

All of the composite panels tested in the program were 

fabricated using an arc-spray process developed at the NASA Lewis 

Research Center by Westfall (1985). This process, schematized in 

figure 5, consists of wrapping the reinforcing fibers on a drum, 

placing the drum in an air-tight chamber, evacuating the chamber. 

and back-filling it with argon. Next, wires of the matrix 

composition are brought together in the arc-spray gun, an arc is 

struck between them, and pressurized argon blows the molten matrix 

material onto the reinforcing fiber mat. The drum rotates until all 

the fibers are covered; the resulting structure is a "monotape" 

consisting of one layer of fibers in matrix (Figure 6). Layers of 

monotapes can then be hot pressed or hot isostatically pressed 

(HIP), producing a fully densified consolidated structure with 

negligible fiber/matrix interfacial reaction (Figure 7). Composite 

panels were fabricated with fibers oriented three ways, as shown in 

figure 8. The majority of panels had unidirectionally aligned, 

axially oriented (0') fibers. Panels were also fabricated with 

transversely oriented (90') fibers, and with angle-pl ied (t15") 

fibers . 

.. 

21 
3C SPRAY HEAD, 

VACUUM 

MATRIX 

SPOOL-,, CHAMBER 

OVERALL VIEW SCHEMATIC OF OPERATION 

Fig 5 Schematic of the arc-spray fabrication process. 

Fig 6 W/Nb as-sprayed monotape. 

22 
* 

> 

0 

v) 

1 

c1, 

c 

c 

m 

u 

.- 

.C 

f 

23 
Several factors determine the volume fraction of fiber in the 

finished composite panel. Most important are initial fiber spacing 

on the drum and the amount, if any, of cladding used on the panel. 

Also, the thickness of the sprayed layer, the initial and final 

density of the sprayed layer, whether or not the monotape was 

sprayed on both sides will affect the volume fraction of fiber 

somewhat. 

were specified to produce panels with 40 volume percent fiber, but 

were not precisely observed. 

around the 40 volume percent value. 

The fiber spacing and spraying parameters for this study 

This resulted in a *25% variation 

U 

LONG 

MICROGRAPH 4 TRANS 

SHORT (0) TRANSVERSE 

POLISHED lsl SECTION 

AXIAL +15O TRANSVERSE ANGLE-PLY 

Fig 8 Schematic of the fiber orientations in W/Nb creep specimens. 

24 
Additional panels without fibers were fabricated to evaluate 

the creep strength of the matrix material in the arc-sprayed 

condition*. 

-,he 1 imited amount of alloy matrix wire material. Unfortunately, 

;he matrix had to be sprayed onto a molybdenum sheet in order to 

keep it from sticking to the drum (the reinforcing wires normally do 

this when a monotape is sprayed). The molybdenum sheet had to be 

etched off of the very porous monotape with nitric acid, thereby 

increasing the interstitial element contamination of the pressed 

panel. 

influenced by this fact. 

Only niobium and Nb-1Zr were tested in this way due to 

Consideration of 'the matrix-only creep data must be 

1 e PreDarati on 

Creep specimens having a reduced gauge section 25.4 mm long 

and 6.35 mm wide were electric discharge machined (EDM) from the 

composite panels. Tungsten or tantalum tabs (0.5 mm thick) were 

attached to the ends of the samples to prevent shear pullout of the 

pinhole area. Figure 9 depicts a fractured creep specimen. These 

tabs were electron beam welded onto the samples with fibers oriented 

0" and t15" to the stress axis; this resulted in several of these 

samples breaking in the pinhole area during creep testing. 

this was partially ascribed to the fact that alloying niobium with 

Although 

*Arc-spraying produces a unique matrix microstructure; a1 though it 

is not "rapidly solidified," it is quickly cooled against the 

rotating drum, resulting in a fine-grained microstructure. 

25 
tungsten embrittles the niobium, the large amount of porosity 

observed in the weld was more likely at fault. 

ascribed to gases in the arc-sprayed material; either argon bubbles 

or dissolved oxygen and nitrogen from impurities in the argon. 

eliminate this problem, those samples with fibers oriented 90' to 

the stress axis had tantalum tabs attached to the samples during the 

HIP cycle. 

This porosity was 

To 

Fig 9 Typical failed W/Nb composites creep specimen. Tungsten 

tabs (0.5 mm thick) were electron beam welded onto the 

ends of the samples to prevent shear pullout of the 

pinhole area. 
Creep Testing 

Creep-rupture tests in a constant load conditions were 

conducted in a vacuum of 7~1O-~ Pa. 

optically via a cathetometer clamped to the furnace chamber frame 

sighting on Knoop hardness impressions placed 25.4 mm apart in the 

reduced section. The creep apparatus is fully described elsewhere 

by Titran and Hall (1965). The precision of creep strain measurements is estimated to be -0.02% for the gauge length used. The 

strain on loading was measured and was incorporated in the reported 

total creep strain. 

Creep strains were measured 

A limited number of samples were annealed in a vacuum of 10-8 

Pa without an applied stress to evaluate the effect of temperature 

without applied stress on diffusion. 

1500 and 1800 K for times up to 1000 hr. 

These pieces were annealed at 

le Analysis 

Specimens were cut from the creep tested samples as shown in 

Figure 8. 

automatic polishing equipment by the procedure in Table 11. 

Samples were mounted in epoxy and polished using Struers 

27 
These polished samples were examined on an ARL Electron Prabe 

Microanalyzer. Traces were made for niobium and tungsten in 

selected samples, usually those with longer exposures. 

Microhardness measurements on a Buehler Micromet I1 Microhardness Tester were also made in the unetched condition. Although 

were 

ne a gradient of hardness across the interd fthe lightest weight 

too large to determ 

fusion zone. 

available was used (10 g), the indentations 

For optical evaluation, the polished specimens were etched for 

tungsten, then for niobium. The etchants are given in Table 111. 

Table 111. Ftchina Solutions 

15 ml nitric acid 

5 ml hydrofl uori c 

I swab 15-30s ! 

28 
Individual fibers in the etched samples were photographed 

using differential interference contrast (DIG) using a Nikon 

Epiphot-TME Inverted Microscope. Low magnification micrographs 

(using collimated light) were taken on a Reichert MEF3 Inverted 

Metallograph. 

Stereoscan S200 Scanning Electron Microscope. 

The fracture surfaces were examined in a Cambridge 

29 
ComDosite C reeD ProDert i eq 

Fai 1 ure Anal ysi s 

Typical fracture surfaces of creep-tested specimens are shown 

in Figure 10. In these micrographs, the matrix appears to have 

undergone ductile failure, whereas the fibers failed in a brittle 

manner. In fact, the app1,ied stress on the composite exceeds the 

elastic limit for the matrix, but the fibers prevent it from 

failing. It is possible, too, that the matrix may deform 

plastically without failing, depending on how far the fibers 

elongate. In that case, the matrix functions primarily as a 

protective coating for the fibers during creep, but still 

contributes some strength to the composite. Since conventional 

failure mechanisms in the matrix are inhibited due to the constraint 

of the fibers, the matrix withstands a higher applied stress before 

it deforms plastically than it is able to in a stand-alone test. 

addition, substantial grain growth has occurred in the arc-sprayed 

matrix during exposure at elevated temperatures. 

In 

It is assumed that some grain broadening has taken place in 

the fibers. This grain broadening results in an increased grain 

width, but does not affect the aspect ratio of the grains, about 

30 
Fig 10 Fracture surfaces of W/Nb composites creep tested at 

1500 K under an applied stress of' (a) 228 MPa for 425 hr, 

(b) 80 MPa for 5353 hr, (c) 153 MPa for 937 hr, and (d) 

131 MPa for 2485 hr. SEM micrographs reveal the 

different fracture modes of the fiber and matrix. 

31 
2Oo:l. The fibers fail with little apparent necking, although 

fibers tested singly exhibit substantial necking prior to failure. 

A reason for that may be that when one fiber reaches the point of 

necking, it is constrained by the rest of the fibers which have not 

reached that point, extending the life of the fibers. 

Trefilov et al. (1976) found in tensile tests of M/Nb composites 

that fibers in the matrix achieve 1.5 to 2 time the elongation 

versus that of the fibers tested separately under the same 

conditions. 

In fact, 

Because the applied stress is divided among all of the fibers, 

when one or a few fail in creep, the load is redistributed, 

increasing the load on those remaining. Eventually, the elastic 

limit of the surviving fibers is reached and they fail transgranularly in tension; this results in a standard 45" failure angle for 

the axially reinforced composites, with all of the fibers fracturing 

on about the same plane. 

Creep Curves 

Since fiber failure dictates composite failure, the crear) 

behavior of the composite resembles that of the tungsten fibers. 

This behavior is demonstrated in Figure 11, depicting several 

typical creep curves for W/Nb composite specimens. The curves 

exhibit primary, steady-state, and tertiary creep. 

32 
1.2 r 

1.0 5 

OJJ 

5 .8 

a 

.2 

I I I I 1 

TIEE. hr 

ST300/Nb- 12 r 

1500 K/84 Wa 

tlx = 1099hr 

1, = 5353 hr' 

*ONLY FIRST 1200 hrOF DATA PLOTTED. 

L 

C 

u1 

- 

E 

TIE, hr 

ST300/Nb-lZr 

1500 ~1137 wa t = 2549 hr' 

*TEST STOPPED. 

t 1x = 349 hr 

0 

8 

f 

E 

C 

VI 

0 

0 

0 

TIE. hr 

smm-1zr tlr - 173hr 

1400 w268 pa tR = 990 hr 

0 

3- 

400 600 800 1000 

TIRE, hr 

ST300/Nb t 1x = 205 hr 

1400 K/2% MPa tR = 1012hr 

6 't 

4c 

0 

0 

0 

/ 

0 

5- 0 

4- 

3- 

0 

0 0 100 200 300 400 500 600 700 

:: 

8 

TIRE, hr 

ST300/Nb - A 

1500 ~199 ma t = 685 hr 

t 1x = 158 hr 

0 

0 

0 

,L :0 oo 

2 0200 

TIRE, hr 

t 1x = 78 hr 

tR = 715 hr 

ST300/N b 

is00 K/IS Ra 

Fig 11 Typical creep curves for W/Nb composites. The curves 

exhibit primary, steady, and tertiary creep. fitted to the steady state creep regime. 

Lines are 

33 
Three pieces of data were taken from each creep curve: the 

time-to-rupture, the minimum creep rate (or the slope of the line 

exhibiting steady state creep), and the time-to-1%-strain, generally 

a point on that line. The results for the composite samples are 

given in Table IV. 

-35 to 52 volume percent, applied creep stresses were normalized to 

50 volume percent fiber for comparison purposes by the following 

equation, 

Since the fiber content varied from 

This equation assume the entire load is carried by the fibers, 

and is therefore not a true normalization. Although it is valid for 

comparing axially reinforced composites to other axially reinforced 

composites, and for comparing angle-plied composites to other 

angle-plied composites, it may not be valid for comparing composite 

structures with different fiber orientations to each other. The 

alternative, however, is to compare the data by the applied stress 

regardless of volume fraction. The 1 atter approach is definitely 

not valid, so the creep data for axial and angle-plied composites 

plotted here has been normalized. 

samples appears to occur solely in the matrix (Figure 12), so it is 

more correct to use the actual applied stress in this case. 

Fracture of the transverse 

34 
TRANSVERSE 

ST300/Nb-lZr tlx = 62 hr 

1500 K/25 RPa tR = 525hr - 

LONG TRANSVERSE 

Fig 12 Fracture surfaces of a W/Nb composite with transverselyoriented fibers exhibiting ductile failure in the matrix. 

At top, a bundle of 7 fibers which has partially debonded 

from the main part of the composite is depicted. 

bottom, the matrix after the fibers have pulled away is 

shown. 

At 

35 
Table IV. Comooslte Crem e 

IT300 / A5 * 

IT3001 A5 

;T300/AS 

iT300/A9' 

iT300/A9 

;T300/A9 

iT300/A9 

iT300 /Nb 

iT300/Nb 

iT3OO/Nb 

iT3OO/Nb 

iT3OO/Nb 

iT3001Nb 

iT3OO/Nb 

iT3OO/Nb 

iT300/Nb 

;T300/Nb 

;T3OO/Nb 

iTSOO/Nb 

5T300/Nb 

iT3OO/Nb 

iT3OO/Nb 

iT300/Nb 

STJOO/Nb-A' 

jT300/Nb-A 

jT300/Nb - A 

jT300/Nb-A 

jT300/Nb-A 

jT300/Nb-A 

jTJM)/Nb-A 

ST300/Nb- A 

ST3001Nb-A 

ST3OO/Nb-A 

ST300 / Nb - A 

ST300/Nb- A 

ST300iNb-IZr 

ST3OO/Nb-IZr 

ST300/Nb- 1Zr 

ST300/Nb- 1Zr 

ST300 1 Nb- 1Z r STJOO/Nb- 1Zr 

ST300/Nb- 1Zr 

ST300/Nb- 1Zr 

ST3001Nb- 1Zr 

ST300lNb- 1Zr 

ST300/Nb- 1Zr 

ST3001Nb- 1Zr 

ST3OO/Nb- 1Zr-A 

ST300/Nb- 1Zr-A 

ST300/Nb- 1Zr-A 

ST3OO/Nb- 1%-A 

ST300/Nb- 1Zr-A 

ST300lNb- 1Zr-T' 

ST300lNb- 1Zr-T 

ST300/Nb- 1Zr-T 

1400 

1400 

1400 

1400 

1500 

1500 

1500 

1500 

1500 

1400 

1500 

1500 

1400 

1400 

1 500 

1500 

1400 

1400 

1400 

1400 

1400 

1400 

1400 

1400 

I500 

1500 

1500 

1500 

1500 

1500 

1500 

1500 

1400 

1400 

1400 

1400 

1400 

1500 

1500 

1500 

1500 

1500 

1500 

1500 

1400 

1400 

1400 

1400 

1400 

1400 

1400 

1400 

1500 

1500 

1500 

1500 

1500 

1500 

1500 

1500 

1400 

1400 

1400 

1500 

1500 

1500 

1500 

1500 

1 ao 

190 

300 

350 

130 

150 

160 

300 

220 

300 

130 

150 

190 

300 

130 

I50 

160 

iao 

iao 

190' 

300 

300 

310 

320 

95 

loo 

110 

110 

130 

130 

150 

140 

150 

150 

160 

65 

70 

90 

90 

110 

150 

160 

ao 

iao 

as 

iao 

iao 

300 

300 

230 

250 

360 

60 

100 

100 

110 

130 

130 

160 

306 

150 

190 

70 

iao 

ao 

38 

25 

30 I1 

'A1 denotes Alloy matrix Nb-ZZr-i'd 

*A5 denotes Allay matrix Nb-lZr-SU 

'A9 denotes Alloy matrix Nb-lZr-gU 

+A denotes angle-plied (215') fibers 

+T denotes transverse oriented fibers 

'formdid 

Stms 

(MP4 

239 

255 

366 

335 

169 

311 

355 

395 

250 

17 1 

193 

265 

173 

202 

356 

395 

3 64 

270 

327 

303 

110 

138 

13 1 

153 

303 

300 

19 1 

335 

247 

310 

366 

102 

119 

119 

131 

153 

19 1 

355 

390 

316 

250 

2 17 

363 

4 17 

139 

137 

155 

145 

133 

170 

217 

251 

313 

341 

144 

198 

348 

338 

2aa 

iao 

xa 

:ea 

3 ia 

a4 

36 

--- 

tlrne-Lo1% rtmn 

(h) -~ 

343 

39 

16 

10 

17 

23 

16 

3 

79 

40 

" 

I 

138 

loa 

6a 

91 

65 

305 

2 43 

333 

102 

56 

40 

705 

710 

93 

95 

88 

56 

93 

180 

31 

54 

7a 

.a', ** 

825 

178 

isa 

78 

136 

72 

34 

5 

1 

173 

79 

277 

340 

340 

70 

6 

1099 

119 

349 

313 

106 

190 

73 

17 

39 

14 

304 

65 

42 

62 

105 

238 

1 a94 

426 

266 

382 

73 

30 

3 

BO 1 

585 

331 

533 

35 1 

10 13 

315 

683 

-., 1.l- 4 

-- 

a50 

387 

3548 

3485 

161 

349 

937 

475 

436 

344 

37 3 

345 

973 

90 

369 

715 

384 

885 

a22 

3aa 

133 

44 

990 

457 

i490 

1614 

1104 

3325 

379 

33 

5353 

3 

999 

742 

92 1 

425 

53 

197 

944 

325 

680 

ai9 

3s 

3a 

3s 

za 

$8 

37 

37 

39 

37 

40 

39 

:a 

:a 

:a 

3a 

37 

31 

31 

36 

37 

31 

37 

36 

36 

37 

38 

:6 

30 

36 

39 

30 

30 

36 

30 

35 

32 

39 

36 

34 

39 

33 

39 

30 

31 

42 

46 

40 

51 

47 

31 

36 

36 

37 

36 

41 

49 

47 

47 

30 

39 

38 

-3 

.a- 

?a 

38 

40 

41 

41 

fiber 

;.~~xLo-) 

3.10~10-~ 

3.47~10-~ 

.?.39x10-7 

3.ai~io-a 

7.63~10-P 

1.53xLO-: 

S.3OxlO-; 

3.74x:o- ' 

l.71x?o-d 

?.37x:O-d 

:.35x:o-a 

3.77xio-d 

a.74x:o-g 

3.5 1x10-a 

7.aox10-9 

1.9ox 10 -a 

5.62x:O-a 

5.33~?0-~ 

1.30~ 

3.74~!0-~ 

9.70~?0-~ 

3.95~10~~ 

3.13~10-~ 

6.34~10-~ 

f.93~10-~ 

1.04~10-~ 

~.'OXIO - 9 

2.13x10-a 

3.5 1x10-a 

2.3ox10-a 

3.15x?0-8 

3.jk:O-a 

7.55x?o-a 

1.35x10-j! 

1.07x10-' 

4.33~10-~ 

1.62~10-~ 

1.63~10-~ 

3.53~10-8 

1.05x10-< 

5.49~10-1 

3.59x10-0 

1.33~?0-~ 

3.68~10-a 

7.04x10 -4 

8.43~10-~ 

926x10-9 

6.56xiO-9 

254x10-8 

i.aaxio-8 

4.11~10-7 

1.a3x10-9 

3.1 4x:o-d 

j.2axio-9 

:.:aXio-d 

1.70x:o-d 

2.77x:o-:! 

4.a3x10-a 

d.70x10-a 

5.33~10-9 

:.92x1o-a 

5.35x:o-a 

3.62x:O-S 

1.77~:0-~ 

9.65x:013 

1.17Xio-' 

-- 
The time-to-rupture versus the normalized applied stress 

plotted in Figure 13 compares the composites at 1400 K and 1500 K, 

and the minimum creep rates taken from the creep curves are compared 

in Figure 14. 

the following sections. 

Various aspects Of these data will be discussed in 

Mu1 ti pl e Li near Regression 

Multiple linear regression of the creep data for the niobium 

matrix-only samples, ST300 fibers, and ST300/Nb composites was used 

to generate equations for tR, according to the Orr-Sherby-Dorn 

(1954) analysis, of the form, 

Es 5 

-1 

where t, is a "pseudo" creep-rupture rate with units of hr taken 

as the slope of the straight line drawn from the origin to the Point Of 

failure on a plot of time versus applied stress*. 

stress in MPa is denoted by u, raised to n, the stress exponent. 

The argument of the exponential is (-Q/RT), where Q is the 

The applied 

*Similarly, tis is the slope of the line drawn from the origin to 

the point where the sample reached 1% strain. 

therefore O.Ol/hr. 

The units are 

37 
500 

a $ 300 - 

LL 2s 

I 0 200 -In iz 

aE 

U 

Fig 13 Time-to-rupture vs the normalized applied stress (50 v%) 

for W/Nb composites creep-tested at 1400 and 1500 K. 

- 

--& 

- 

1400 K 

1500 K 

a 

-\ 

- 

L 

0 

z 

W /'kOO K 

/ -&!+ ST3001Nb 

-4 It- ST3001Nb-1 Zr 

- -0-- 

A 

ST3001N b f 15' 

- ++ ST300/Nb-lZr A 

- -0- ST300/Nb*15' ANGLE-PLIED 

- 

I 1 

50 

10-9 

LL 

10-8 10-7 10 

MINIMUM CREEP RATE (s-l) 

-6 

Fig 14 Minimum creep rate vs the normalized applied stress 

(50 VX) for W/Nb composites creep-tested at 1400 and 

1500 K. 

38 
activation energy in joules/mole, R is the gas constant (8.3154 

joules/degree-mole), and T is the temperature in Kelvin. The 

computer analysis used to determine the constants in this equation 

is included in Appendix A. 

The following equation was found for the ST300/Nb composites 

(from data in Table IV), 

50000 

5.11. -- r t, = 2.05. u e 

The creep data for the niobium matrix-only creep samples is 

listed in Table V. Regression analysis produced the following, 

Table V. Wix Cre. an Rata 

4.17. -- t,=4.28~10~~*~ e' 

Temp 

(K) 

1400 

1400 

1400 

1400 

I400 

1400 

1500 

1500 

1500 

1500 

I400 

1400 

1400 

1400 

I500 

I500 

I500 

Applied 

Stress 

(MPa) 

IO 

15 

18 

20 

20 

25 

5 

10 

10 

15 

15 

20 

25 

35 

5 

IO 

20 

time-to- 1% 

strain (h) 

347 

233 

54 

17 

12 

5 

56 

3 

51 

6 

1 

99 5 

7 

- 

i 

time- torupture (h) 

754 

60 

242 

29 

523 

79 

93 

69 

I607 

728 

592 

minimum 

creep rate 

b-1) 

7.69~ 10-9 

1.06~10-7 

5.1 7x10-8 

1.OoX10-~ 

1.35xIO-7 

5.OOx IO-; 

3.9OX 10-3 

;.:ox 10-7 

2.99~ 10-7 

:.sox 10-3 

1.19~ I04 

I .:ox 10-7 

7.76~ IO-: 

l.59xlO-8 

4.16~ 10-1 

I.25.x 10-8 

39 
Stress-rupture data, for ST300 wire (Petrasek and Titran 1988, 

Petrasek and Beremand 1987) is listed in Table VI. The creep 

equation derived from this data was, 

1400 

1400 

I400 

I400 

1400 

1400 

1400 

1400 

1400 

1400 

I400 

I400 

1400 

1478 

1478 

I478 

1478 

I478 

1478 

1 SO0 

1 SO0 

1 SO0 

I500 

1 SO0 

1500 

1500 

1500 

I500 

1500 

1500 

I’ 

ZTOM) 

-12. 5.58. -- t, = 8.42s 10 u eT 

Table VI. ST700 FibTemp 

(K) 

1255 

1255 

I255 

I255 

I366 

I366 

1366 

I366 

1366 

1366 

1400 

1400 

I400 

Applied Stress 

(MP3) 

-3 

-- 1-4 

1-4 

758 

793 

483 

552 

552 

62 1 

690 

758 

414 

448 

448 

4733 

483 

483 

517 

552 

586 

586 

620 

620 

620 

690 

690 

753 

276 

245 

4 14 

483 

552 

62 I 

230 

310 

345 

245 

279 

379 

414 

414 

448 

48 3 

517 

time- torupture (h) 

213 

I26 

191 

89 

548 

85 

87 -. 

!i 

-- 

I 

402 

333 

I92 

99 

124 

131 

64 

34 

33 

i5 

2Q 

15 

5 

11 

I069 

I72 

49 

I7 

6 

524 

242 

IO5 

I 39 

32 

40 

20 

34 

12 

6 

3.9 

-- 

-J 

J 

-! - 

40 
In addition to creep-tested composite specimens, samples were 

annealed for 100, 500, and 1000 hrs to evaluate interdiffusion in 

the absence of an applied stress. 

zone for both types of specimens was measured by electron ?robe 

microanalysis. A sample trace is shown in Figure 15. 

these microprobe traces are plotted in Figure 16, where they show 

good correlation to diffusion data found in the literature. 'de 

might expect differences due to the different microstructure of the 

The width of the interdiffusion 

The resuits of 

ORIGINAL 

FIM INTERFACE 

ST300lN b 

1400 K 

833 h 

W 

FIBER 

f- INTERDIFFUSION 

Fig 15 Typical microprobe trace across the fiber/matrix interface 

in a ST300/Nb composite exposed to 1400 K in vacuum for 

833 h under a normal ized appl ied stress of 238 MPa. 

extent of interdiffusion is indicated taking the 

1 imitations of the instrumentation into account. The 

position of the original fiber/matrix interface imp1 ies 

that the mjority of the diffusion occurs from the matrix 

into the fiber. 

The 

41 
106 

lo4 

2 102 

n 

\ 

N 

E 

5 

7 10 N 

U 

TEMPERATURE. K 

2000 1800 1500 1400 

I I I I 

t .--O- ARCELLA 

-0.- THIS STUDY 

I I I I I I 

.45 .so .55 .60 .65 .70 .75 

VT (VK x 103) 

Comparison of diffusion results to Arcella (1974). 

Circles represent composite specimens with a niobium 

matrix; squares, a Nb-1Zr matrix. 

Fig 16 

tungsten used; the arc-cast and CVD tungsten Arcella evaluated is 

quite different from the fine-grained tungsten in a drawn fiber. It 

appears, however, that diffusion is controlled by bulk processes at 

these temperatures (Hehemann and Leber 1966). Also, the radius of 

curvature of the interface in the composite may influence diffusion 

coefficients. However, these effects have been found to be minimal 

with 200 pm fibers (Kopp et a1 1988). The measured data fits the 

equation, 

26692 

In ($) = - T + 16.4 

42 
where 1-y is the width of the interdiffusion zone in prn, t is the 

time in hours, and T is the temperature in Kelvin. 

relates the extent of the diffusion to the t me and temperature of 

exposure. In addition, a plot of the diffus on at 1400, 1500, 1800 

K versus the square root of time is given in Figure 17, which 

exhibits traditional diffusion behavior according to the t1/2 

dependence. No difference was observed between Nb and Nb-1Zr matrix 

composites. 

This equation 

60 

50 

40 

h 

5 30 u - 

20 

10 

100 500 1000 2000 5000 , I I I/ / I I 

/ 

A //1800 K 

/ 

- 

A/ / 

/ 

/ 

/ 

- 

/ 

// b A 

- A/ 

A/ / 

/A /- 

1500 K -e &' a / 

L/ 0 

/- 

/ 

- 

/ 

/ 

/ % 1400 K 

/ 

I 

" U 

o' 0 0 

'/ 

I 1 I 1 

Fig 17 Comparison of the extent of interfacial diffusion at 

various temperatures. 
The interdiffusion between the fiber and matri.t wnich occ3r-s 

upon elevated temperature exposure i s a1 SO apparent i n the fracture 

nicrographs (Figure lo), where the measured diameter of the fibers 

including the reacted areas does not appear to have changed from the 

as-fabricated condition (-200 W). In fact, microprobe traces of 

the fiber/matrix interface indicate that the majority of the 

diffusion is into the fiber (Figure 15). When the polished sections 

are examined (Figure la), a recrystallized zone is evident within 

the original outer diameter of the fiber. These facts show that 

niobium diffuses into the fiber and lowers its recrystallization 

temperature, resulting in a recrystall ized and tungsten-depleted 

area within the original fiber diameter. 

In Figures 13 and 14, a negative deviation from the 

established slope of each of the lines is apparent for specimens 

subjected to relatively low applied stresses. Since these spec mens 

were tested for correspondingly long times in a relatively poor 

vacuum, the negative deviation may be due to interzitital react on 

in the matrix--niobium forms a volatile oxide at these temperaturis. 

However, it may also be attributable to the growth of the 

fiber/matrix reaction zone. 

no longer has the high-strength fibrous grain structure, and no 

longer contributes to the strength of the composite in the capacity 

it did at the start of the test. 

This recrystallized area of the fiber 

44 

We can use a simple calculation to assess the effect of the 

strength of the interface on the strength of the composite. 

the rule of mixtures, 

Recall 

which, as was mentioned previously, does not account for inelastic 

behavior of the components. However, positive or negative deviation 

from this equation can he1.p us to understand the composite creep 

behavior. We can calculate the volume fractions of each of the 

components by making some assumptions about the interdiffusion and 

the initial volume fraction. These calculations are detailed in 

Appendix B. 

Eq 10 

Eq 11 

Eq 12 

u, = 0.5 

(Do - 2~)~ 

20: 

Uf = 

(Do -2x)2 

u- = 0.5- 

Now, Eqn 9 gives an expression for Ax, the width of the 

reaction zone as a function of time and temperature. Since we also 

46 
know the original fiber diameter Do, we now have expressions for the 

volume fractions of each component as a function of time and 

temperature. 

We also have expressions for the creep strength of the niobium 

matrix material (Eqn 6) and the ST300 fiber (Eqn 7). Now, the only 

unknown in Eqn 3 is the creep strength of the interface zone. In 

order to better estimate these properties, microhardness measurements were taken on the fiber, matrix, and the interface zone. 

Figure 19 contains several micrographs displaying the position of 

the indentations and their corresponding hardness readings. Only a 

few samples had a reaction zone of sufficient width to accomodate 

the entire indentation. Because of the composition gradient, an 

indentation which encompassed the recrystallized zone exactly might 

best represent the hardness of the interface layer. The measured 

hardnesses indicate that the reaction zone is somewhere in between 

the properties of the fiber and the matrix, but actual estimations 

are impossible due to the placement and limited number of the indentations. 

hardening of the matrix phase; this was ascribed to extensive 

oxidation of the matrix during testing. 

Some of the samples exposed for very long times exhibited 

Further, we can calculate the creep strength of the composite 

using Eqn 3 while varying the properties of the interface zone from 

47 
Fig 19 Microhardness measurements of composite sections after 

creep testing. 

48 
that of the fiber to that of the matrix*. Three cases are plotted 

in Figure 20, where the properties of the interface are set equal to 

those of the matrix, to those of the fiber, and to an average 

between the two. It is seen that at higher stresses, where the 

samples are exposed to temperature for correspondingly shorter 

times, the difference between the three cases vanishes. Therefore, 

one can see that at shorter times, the width of the reaction zone is 

small, so the value for u, is also small; this minimizes any 

differences in u:. At lower stresses (and longer times), the 

difference becomes appreciable. 

However, when the creep strengths of all the components are 

plotted together, as in Figure 21, the small differences arising 

from the strength of the interface zone become secondary. 

what value is used for interface strength, the measured creep 

properties exceed those calculated from the rule of mixtures. 

Therefore, the creep strength of the composite is greater than the 

combined creep strengths of its components. 

No matter 

There may be several 

reasons for this synergistic effect. 

The factors present in the composite that do not exist in 

stand-alone tests include inelastic strains due to differential 

thermal expansions and diffusion effects. It is unlikely that 

*Cal cul ati on detai 1 ed in Appendix C. 

49 
;F 40 

In[-] 

o a =om 2 

0 oz-- - +am 

2 

az = af 

0 

0 

0 

0 

0 

e 

on 

000 

0 00 

00 00 

Fig 20 Calculated strength of the composite using Eq 10 varying 

the properties of the reaction zone. 

20 ' 

CALCULATE? 

, 

0 

I I I I 

0 '2 4 6 8 10 

In a 

Fig 21 Comparison of matrix and fiber independent creep 

properties versus calculated and measured composite 

strengths. 

50 
thermal stresses contribute to this effect for two reasons; first, 

that these stresses will relax after the initial heat-up under creep 

conditions, and second, that the difference will tend to put more 

stress on the fiber and on the matrix, not less. 

One effect of interdiffusion is the loss of a portion of the 

high-strength fiber microstructure. This effect has been shown to 

be minimal, however, compared to the overall strength of the 

composite (Figure 21). 

Diffusional effects will aid in strengthening for one main 

reason: interdiffusion promotes a strong and ductile fiber/matrix 

bond. Once diffusion occurs, even to the extent that it does during 

fabrication, there is little chance of debonding taking place. Tnis 

bond may be the most important contributor to the synergistic 

strengthening of the composite. By constraining the fibers and the 

matrix to the same amount of strain, we realize maximum strength 

contributions from each component in the composite. 

not deform plastically even though it would without the fibers, and 

the fibers are kept from necking by the constraint of the rest of 

the fibers which are maintaining their integrity. 

The matrix does 

51 
.. Ef feet of Matrix Alloy Additions 

The effect of interdiffusion is further demonstrated if we 

compare the creep properties of the composites to that of the fiber 

itself. In Figure 22, the stress-rupture properties of ST300 fibers 

have been extrapolated using the Monkman-Grant (1956) re1 ationship 

(which states that the time-to-rupture and the minimum creep rate 

are inversely proportional) to plot the theoretical time-to-1%- 

strain at 1500 K for a bundle of fibers assuming no surrounding 

matrix*. On the same graph, the time-to-1%-strain for composites 

with several matrix compositions is plotted. 

S 

strength 

nce the fibers contribute almost all of the composite 

it is expected that the values for the composite and the 

fibers would be parallel. It is apparent from Figure 19 that, at 

the beginning of the test, they are comparable. 

however, the composite properties diverge negatively from the 

theoretical values calculated for the fibers. This can be 

attributed to two factors. First, the growing reaction zone 

removes an increasing amount of the high-strength fiber microstructure; therefore, a smaller volume fraction of fiber remains to 

contribute to the strength of the composite. Second, those samples 

As time progresses, 

~~~ - 

*This calculation is detailed in Appendix D. 

52 
r 

MATRIX COMPOSITIONS 

+ 

-0- Nb 

Nb-1Zr 

A Nb-1Zr-5W 

-+- Nb-2Zr-1W 

Q Nb-lZF9W 

I I I. 

. 

1 10 100 1000 

TIME TO 1% STRAIN (h) 

Fig 22 Comparison of the creep strength of ST300 fiber reinforced 

composites vs that calculated for a bundle of fibers with 

an equivalent cross-sectional area of fiber at 1500 K. 

exposed for longer times adsorb proportionally more interstitial 

contaminants which may affect the strength and integrity of the 

matrix. 

The three lines plotted for the composite properties in 

Figure 22 represent composites with 0, 1 and 2 w% zirconium 

additions to their niobium-base matrices. There is an increasing 

53 
negative deviation from the ideal fiber strength with increasing 

zirconium matrix content, the converse of the tensile results shown 

in figure 4, where an increasing zirconium content strengthened the 

natrix alloys. 

enhances the fiber/matrix reaction, resulting in a decreas2 in the 

volume fraction of fiber in the composite, and reducing the strength 

of the composite at long times. 

zirconia precipitate phase encourages the nucleation of Kirkenaall 

voids at the interface: their formation is discussed in the next 

section. 

It is possible that the addition of zirconium 

Another possibility is that the 

A small number of creep data points for samples with 5 and 

9 w% tungsten are also plotted in Figure 22. 

difference in the creep strength between Nb-1Zr with and without 

5 w% tungsten, whereas the addition of 9 w% tungsten appears to 

increase the creep resistance of the composite by -10% at a time-to1%-strain of 100 hr, and up to 24% at a time-to-1%-strain of 3000 

hr. This may be ascribed to the higher strength of the matrix (as 

shown in Figure 4), to a decrease in the rate of diffusion between 

the tungsten fiber and the niobium alloy matrix, or a combination of 

the two. 

There is little 

These results indicate that at least 9 w",; tungsten must be 

added to see any eif?ct of the tungsten addition. (Therefore, the 

lilditlorl of less t!iiln 1 w'i, tungsten can be reoardea as negligible. 

54 
and the decrease in strength of the Nb-2Zr-1W can be mainly 

attributed to the zirconium addition.) 

that an addition of 9 w% tungsten to the matrix increases the 

density of a 50 VX fiber composite by approximately 3%; hence, the 

increase in real strength of the composite correlates to a somewhat 

lesser increase in its specific strength. 

tungsten may impart enough additional creep resistance to the 

composite to make it worthwhile, longer term tests will be necessary 

in order to confirm and extrapolate this data. 

It is interesting to note 

While the addition of 

In addition to recrystallization of the interdiffusion zone, 

the formation of Kirkendall voids occurs at the fiber/matrix 

interface. 

the fiber causes a ring of very small voids to form outside of the 

original fiber diameter. Over time, these microvoids accumulate, 

and the microstructure shown in Figure 23a results. 

noted that these voids nucleate on second phase particles, so the 

presence of a zirconia precipitate in the matrix may encourage their 

formation. These voids at the fiber/matrix interface may affect the 

composite properties given that the integrity of the interface area 

in any ductile metal-matrix composite is essential for load transfer 

from the matrix to the fiber. 

In the composite, the net mass transfer of niobium into 

It has been 

Although it is apparent from the 

55 
Lc) W 

L 

=z 

0 

E n 

v 

+I 

v 

n 

r 

\ 

0 

0 

m 

a 

I 

In 

I\ 

c0 

0 

E 

E: 03.- .r 0 x 

Lm mu 

c 

12 h 

m 

v a 

z 

\ 

0 

0 

M 

c 

v) 

is L 

In 

00 

3 

h( 

LC-r v) 

56 
creep results that these voids do not grossly affect the axial 

properties of the composites, it is probable that they will have a 

detrimental effect on the transverse strength. 

Kirkendall void formation was apparent when the fibers were 

4 

angle-plied _+15*, but only on two "sides" of the fiber. This is 

illustrated in Figurq 23b, where the long transverse direction of 

the composite panel is in the vertical direction in the micrograph. 

When the fibers are angle-plied, compressive stresses are induced on 

two "sides" of the fiber due to the forces which are attempting to 

pull the fibers parallel to the strain axis. These compressive 

stresses inhibit the formation of Kirkendall voids in some area 

around the fiber diameter (Barnes and Mazey 1958). 

the remaining areas just as they do in the unidirectionally 

Voids form in 

reinforced composites. Angle-plying of the fibers can therefore 

augment the transverse properties of the composite in two ways. 

First, some,of the applied stress 

transferred to the longitudinal d 

the induced compressive stresses 

in the transverse direction is 

rection of the fibers. Second, 

ppear to improve the fiber/rnatrix 

bond, thereby increasing load transfer across the interface. 

Micrographs of the fracture surface in composites with transversely oriented fibers (Figure 12) indicate that the failure is 

due to ductile deformation in the matrix rather than failure at the 

fiber/matrix interface or cleavage of the fiber. Furthermore, the 

57 
creep strength of the specimens with transverse fibers are slightly 

stronger than matrix-only Samples. However, these specimens were 

not exposed for very long times, and the void formation may not be 

as developed as it would be at longer exposures. However, the 

applied stress on the fiber/matrix interface in the transverse 

oriented samples would certainly encourage void production. At 

longer exposures, excessive void formation at the interface may 

dominate failure of the composite. 

The creep strength of the composite decreases as the fiber 

orientation with respect to the stress axis changes (Figure 24). 

Conversely, the transverse properties will increase. This type of 

data is especially useful when designing tubing made from composite 

materi a1 s. 

r 

I I I I I I I 

0 10 20 30 40 50 60 70 80 90 

FIBER ORIENTATION WITH RESPECT TO STRESS AXIS 

Fig 24 Effect of fiber orientation with respect to the applied 

stress on the creep properties of W/Nb composites. 

58 
CONCLUSIONS 

Fiber Reinforceme& 

It has been shown that reinforcement with ST300 fibers 

increases the creep resistance of these niobium alloys significantly. In fact, the measured properties of the composite exceeded 

those calculated by the rule of mixtures. 

of the individual components was apparently increased in the 

composite versus in stand-alone tests. 

the constraint the fibers exert on the matrix and on each other, 

preventing plastic deformation. The fiber/matrix reaction, however, 

tends to degrade the composite creep-rupture properties at long 

times. 

zation of a reaction zone in the fiber, it may also be due to 

related Kirkendall void formation which can affect load transfer. 

The strength properties 

This may be attributable to 

Although this degradation is partially due to recrystalliFiber Orientation 

The strength of composites with fibers angle-plied 215" is 

decreased about 20% compared to that for composites with unidirectional, axially aligned fibers. This is apparent when comparing the 

time-to-rupture (Figure 13) and the minimum creep rate (Figure 14). 

Short term creep tests of transverse fiber reinforced composites 

revealed failure in the matrix, and the properties (only about 10% 

of the reinforced material) reflected this. 

59 
Matrix Allovina Additions 

Matrix alloying additions of 9 w% tungsten appear to increass 

the creep resistance of W/Nb composites, although this increase is 

gartially offset by an increase in the density of the composite. 

There is no indication of an effect of smaller additions of tungsten 

to the niobium matrix on the composite properties, but small 

additions clf zirconium apparently cause a slight increase in the 

fiber/matrix interdiffusion rates, thereby lowering the creep 

resistance of the composite. 

Figure 25 compares the time-to-1%-strain of unidirectionally 

reinforced composites to angle-plied composites, monolithic Nb-lZr, 

and monolithic PWC-11 alloy (Nb-1Zr-0.06C) at 1500 K (Titran et al, 

1988). Composites with angle-plied fibers (+15.) are approximately 

20% weaker than unidirectionally reinforced panels, but exhibit 

approximately the same stress exponent, indicating that the same 

creep mechanisms are at work. 

approximately an order of magnitude stronger than mono1 ithic Nb-IZr, 

and are -4 times stronger than the PWC-11 alloy at this temperature. 

in addition, the creep data generated in this study can be ussd to 

predict the allowable stresses of W/Nb composite materials for a 

lifetime of seven years (Figure 26). 

All of the composite materials are 

60 
30 r 

1 n 

A 

0 

UNIDIRECTIONAL 

A 

--- 

PWC-11 

-\ -\ 

-\ ---\. -\ 

-\ Nb-1Zr ST3OOlN b 

ST300/Nb-1 Zr -\ 

-- ST3001Nb ANGLE-PLIED ---\ 

0.3 I I I 

1 10 100 1000 

TIME TO 1% STRAIN (hrs) 

Fig 25 Comparison of the time-to-1%-strain for ST300 fiberreinforced composites and the monolithic Nb alloys Nb-IZr 

and PWC-11 (Nb-1Zr-0.06C) (after Titran et a1 1988). All 

data is plotted on a stress-to-density basis, and the 

composite data was normalized to 50 VX fiber. 

34 r / 

24 I I 

0 100 200 300 400 500 

34 - 

32 - 

30 - 

1400 K: 7 YR 

26 - 1500 K: 7 YR 

I 

0 100 200 300 400 500 

APPLIED STRESS. WA 

Fig 26 Orr-Sherby-Dorn plot showing tJe allowable stress levels 

to limit W/Nb composites to less than 1% strain in 60000 

hours at 1400 and 1500 K. 

61 
Future ADD1 1 cat1 on2 

Tungsten fiber reinforcement is a viable method of strengthening niobium alloys. It is apparent, however, that degradation of 

the creep properties of the composite occurs with long term elevated 

temperature exposure. Small additions of tungsten and zirconium to 

the matrix do not appear to affect the diffusion kinetics sufficiently to allow the use of this material for very long times (up to 

7 years), although it is possible that larger additions may be more 

effective. This would, however, defeat the objective of not 

altering the chemistry of the matrix alloy significantly. It seems 

clear, therefore, that more direct methods of stopping or slowing 

the interdiffusion must be developed before this composite can be 

applied in a long-life application requiring creep resistance at 

temperatures above 1400 K. 

62 
REFERENCES 

Arcella, F.G.: Interdiffusion Behavior of Tungsten or Rhenium and 

Group V and VI Elements and Alloys of the Periodic Table, Part 

1. (WANL-M-FR-74-005-PT1, Westinghouse Astronuclear 

Laboratory, NASA Contract NAS3-13231) NASA CR-134490, 1974. 

Barnes, R.S.; and Mazey, D.J.: The Effect of Pressure Upon Void 

Formation in Diffusion Couples. Acta Met., vol. 6, no. 1, 

Jan. 1958, pp. 1-7. 

Begley, R.T.; Harrod, D.L.; and Gold, R.E.: High Temperature 

Creep and Fracture Behavior of the Refractory Metals. 

Refractory Metals Alloys, I., Machlin, R.T. Begley, and 

E.D., Weisert, eds., Plenum Press, New York, 1968, pp. 41-83. 

Bradt, R.C.; Lenel, F.V.; and Ansell, G.S.: Kirkendall-Effect 

Densification in Sintered Tungsten-Molybdenum Powder Compact. 

Acta Met., vol. 14, No. 4, Apr. 1966, pp. 533-534. 

Braint, C.T.: On the Growth of Creep Voids from Potassium-filled 

Bubbles in Tungsten Wire. Scripta Met., vol. 22, no. 10, Oct. 

1988, pp. 1665-1669. 

Gregory, D.P.; and Rowe, G.H.: Mechanisms of Creep in Columbium 

and Columbium-1% Zirconium Alloy. 

D.L. Douglass, and F.W. Kunz, eds., Interscience, New York, 

Columbium Metallurgy, 

1961, pp. 309-341. 

Hehemann, R.F.; and Leber, S.: Chemical Diffusion in the 

Columbium-Tungsten System. Trans. AIME, vol. 236, 1966, 

pp. 1040-1044. 

Hoffman, E.E.; and Harrison, R.W.: The Compatibility of Refractory 

Metals with Liquid Metals. Report GESP-15, General Electric 

Nuclear Systems Division, Cincinnati, OH, 1968. 

Diffusion Barriers in High Temperature Metal Matrix 

Composites. Scripta Met., vol. 22, no. 9, Sept. 1988, 

pp. 1527-1530. 

Kopp, M.W.; and Tien, J.K.: A Preliminary Study of Ion-Implanated 

Kopp, M.W.; Tien, J.K.; and Petrasek, D.W.: Reaction Kinetics 

Between Fiber and Matrix Components in Metals Matrix 

Composites. Superalloys 1988, S. Reichman, D.N. Duhl; G. 

Maurer, S. Antolovich and C. Lund, eds., TMS, Warrendale, PA, 

1989, pp. 193-201. 

63 
65 

Larson, F.R.; and Miller J.: 

Rupture and Creep Stresses. 

1952, pp- 765-775. 

Lundy, T.S.; and Pawel, R.E.: 

Diffusion Coefficience Measurements. Trans. AIME, vol. 245, 

no. 2, Feb. 1969, pp. 283-286. 

Time-Temperature Relationship for 

ASME Trans., vol. 74, no. 5, June 

Effects of Short-circuiting Paths on 

Lyubimov, V.D.; Gel'd. P.V; Shveykin, G.P.; and Sutina, YuA,: 

Self-Diffusion of Niobium in its Alloys with Tungsten. 

Isvestiya AN SSSR, Metally, vol. 2, 1967, pp. 84-87. 

(Translated in Russian Metallurgy, 1967, pp. 40-43.) 

McDanels, D.L.: Analysis of Stress-Rupture and Creep Properties of 

Tungsten Fiber Reinforced Copper Composites. 

Fiber-Strengthened Metallic Composites, ASTM-STP-427, 1967, 

ASTM, pp. 124-148. 

McDanels, D.L.; Jech, R.W.; and Weeton, J.W.: Metals Reinforced 

with Fibers. Met. Prog., vol. 78, no. 6, Dec. 1960, 

pp. 118-121. 

Monkman, F.C.; and Grant, N.J.: An Empirical Relationship Between 

Rupture Life and Minimum Creep Rate in Creep-Rupture Tests. 

ASTM Proc., vol. 56, 1956, pp. 593-605. 

Mundy, J.N.; Ockers, S.T.; and Smedskjaer, L.C.: Dehancement of 

Impurity and Self-Diffusion in Niobium by Tungsten Additions. 

Phys. Rev. B, vol. 33, no. 2, Jan. 15, 1986, pp. 847-853. 

Orr, R.L.; Sherby, O.D.; and Dorn, J.E.: Correlations of Rupture 

Data for Metals at Elevated Temperatures. Trans. ASM, vol. 

46, 1954, p. 113. 

Petrasek, D.W.; and Signorelli, R.A.: Stress-Rupture and Tensile 

Properties of Refractory-Metal Wires at 2000° and 2200 OF 

(1093" and 1204 "C). NASA TN 0-5139, 1969. 

Petrasek, D.W.: High-Temperature Strength of Refractory-Metal 

Wires and Consideration for Composite Applications. NASA TN 

D-6881* 1972. 

Petrasek, D.W.; and Beremand, G.: Unpublished Data. NASA Lewis 

Research Center, C1 eve1 and * OH, 1987. 

Petrasek, D.W.; and Titran, R.H.: Creep Behavior of 

Tungsten/Niobi um and Tungsten/Niobi um-1% Zi rconi um 

Composities. DOE/NASA/16310-5, NASA TM-100804, 1988 

64 
Reed-Hill, R.E.: Physical Metallurgy Principles. 2nd ed., Van 

Nostrand, New York, 1973, pp. 378-429. 

Smigelskas, A.D.; and Kirkendall, E.O.: Zinc Diffusion in Alpha 

Brass. Trans. AIME, vol. 1. 171, 1947, pp. 130-142. 

Tietz, T.E.; and Wilson, J.W.: Behavior and Properties of 

Refractory Metals. Stanford University Press, Stanford, CA, 

1965. 

Titran, R.H.; and Hall, R.W.: High-Temperature Creep Behavior of a 

Columbium Alloy, FS-85. NASA TN D-2885, 1965. 

Titran, R.H.; Stephens, J.R.; and Petrasek, D.W.: Refractory 

Metal Alloys and Composities for Space Nuclear Power Systems. 

DOE/NASA/16310-8 Report, NASA TM-101364. 

Trefilov, V.I.; Karpinos, D.M.; Tuchinskii, L.I.; Mosieev, V.F.; 

Gorb, M.L.; Vishnyakov, L.R.; Pechkovskii, E.P.; Vasil'ev, 

A.D.; and Kokhanyi, V.A.: Plastic Deformation and Fracture of 

Niobium Reinforced with Tungsten Fibers. Poroshkovaya 

Metallurgiya, vol. 5, 1976, pp. 63-67. Translated in Soviet 

Powder Metallurgy and Metal Ceramics, vol. 15, no. 5, Oct. 

1976, pp. 381-384. 

Tuchinsky, L.I.: Thermodynamic Method of Calculating the Effect of 

Alloying Additives on Interphase Interaction in Composite 

Materials. Fizika i Khimia Obrabotka Materialov 1, 1979, 

pp. 121-125. Translated in Report NASA TM-88391, 1986. 

Veltri, R.D.; Paradis, E.L.; and Douglas, F.C.: Investigation to 

Develop a Method to Apply Diffusion Barriers to High Strength 

Fibers. NASA CR-134719, 1975. 

Welsch, G.; Young, B.J.; and Heheman, R.F.: Recovery and 

Recrystallization of Doped Tungsten. Strength of Metals and 

Alloys, Vol. V, P. Haasen, V. Gerold; and G. Kostorz, eds., 

Pergamon, New York, 1979, pp. 1693-1698. 

Welsch, G.; Kim, K.T.; and Wang, J.J.: Potassium-Implanted Tungsten 

Fibers in Nickel-Base Superalloy for Study of Reaction 

Barrier. Final Report for NASA NAG 3-67, Case Western 

University, Cleveland, OH, 1987. 

Westfall, L.J.: Tungsten Fiber Reinforced Superalloy Composite 

Monolayer Fabrication by an Arc-Spray Process. NASA TM-86917, 

1985. 

65 
Westfall, L.J.; McDanels, D.L.; Petrasek, D.W. ; and Grobstein, T.L. : 

Preliminary Feasibility Studies of Tungsten-Niobium Composites 

for Advanced Space Power Systems Applications. NASA TM-87248, 

1986. 

Ye, T.T.; Botta, W.J.; Labun, P.A.; Christian, J.W.; and Taylor, 

G.: Structure and Orientation of Zirconia in a Niobium Metal 

Matrix. Acta Met., vol. 33, 1985, pp. 477-486. 

Yih, S.W.H.; and Wang, C.T.: Tungsten. Plenum Press, New York, 

1979, pp. 270-299. 

66 
The following is a sample MINITAB calculation. The columns 

contain the following data: 

cl Temperature (K) 

c2 Normal ized appl ied stress (MPa) 

c3 time-to-l%-strain (h) 

c4 1/T 

c5 In (0) 

c6 In ( -01 / tl% 1 

c7 temperature compensated time 

c9 counter; in the plot "A" = 1400 K; "B" = 1500 K 

MTB > print cl-c9 

ROW c1 c2 c3 

1 

2 

3 

4 

5 

6 

7 

8 

9 

10 

11 

12 

13 

14 

15 

16 

1400 

1400 

1400 

1400 

1400 

1400 

1400 

1400 

1500 

1500 

1500 

1500 

1500 

1500 

1500 

1500 

216 238 

217 340 

218 277 

250 240 

263 70 

268 173 

290 79 

417 6 

84 1099 

133 190 

137 349 

139 119 

145 106 

155 212 

179 72 

217 17 

c4 

0.0007143 

0.0007143 

0.0007143 

0.0007143 

0.0007143 

0.0007 143 

0.0007 143 

0.0007 143 

0.0006667 

0.0006667 

0.0006667 

0.0006667 

0.0006667 

0.0006667 

0.0006667 

0.0006667 

c5 

5.37528 

5.37990 

5.38450 

5.52146 

5.57215 

5.59099 

5.66988 

6.03309 

4.43082 

4.89035 

4.91998 

4.93447 

4.97673 

5.04343 

5.18739 

5.37990 

C6 

- 10.0774 

- 10.4341 

-10.2292 

-10.0858 

-8.8537 

-9.7585 

-8.9746 

-6.3969 

- 11.6073 

-9.8522 

- 10.4602 

-9.3843 

-9.2686 

-9.9618 

-8.88 18 

-7.4384 

c7 

28.8168 

28.4602 

28.665 1 

28.808s 

30.0406 

29.1358 

29.9197 

32.4974 

24.6940 

26.4491 

25.8411 

26 - 9170 

27.0327 

26.3396 

27.4195 

28.8630 

c9 

1 

1 

1 

1 

1 

1 

1 

1 

2 

2 

2 

2 

2 

2 

2 

2 

*Minitab computer Software, Release 5, December 1986. Minitab, 

Inc., 3081 Enterprise Drive, State College, PA 16801. 

67 
MTB > regress c6 2 CS c4 

Tire regression equation is 

Pr ed i c t or Coef S tdev t -rat i o 

Constant 2.920 3.893 0.75 

4.7832 0.5182 9.23 

C4 - 53452 8209 -6.63 Ci 

~6 = 2.92 + 4.78 cs - 54452 c4 

= 0.4798 R-sq = 86.9% R-sq(adj) = 84.9% 

Anal ys i s of Vari ance 

DF ss MS 

9.9353 

0.2302 

SOURCE 

Regression 2 19.8706 

Error 13 2.9924 

Total 15 22.8630 

SOURCE DF SEQ SS 

c5 1 9.7417 

c4 1 10.1289 

MTB > let c7 = c6 + ~4~54452 

MTB > lplot c7 c5 c9 

A 

AA 

2 AA 

2 

B 

BB 

BB 

B 
The accepted method to determine volume fraction of a 

composite is to first weigh the composite structure, etch away 

either the matrix or fiber material, then weigh the remainder; the 

volume fraction can then be computed. This technique is most useful 

when the fibers are discontinuous or have a very small diameter. 

Its main disadvantage lies in measurements when there are more than .. 

two components, such as an interface layer between the fiber and 

natr i x . 

In the case of tungsten fiber reinforcement, a simpler 

approach is to measure the cross sectional area of the panel, count 

the number of fibers in that cross section, and (knowing the fiber 

diameter), calculate the area fraction, which equals the volume 

fraction. The amounts of any other phases in the composite whose 

area can be measured, such as a reaction zone, can be measured and 

taken into account in this calculation. This technique assumes two 

things; first, that the fibers are unbroken and extend through the 

length of the panel, and second, that the diameter of the fiber is 

relatively invariant. 

Calculation of Volume Fractions 

Assumptions: 

69 
le Mill assume that the entire affected reaction zone is inside the 

original fiber diameter, as shown in Figure 27. This assumption is 

reasonable considering the microprobe results (Figure 15) and the 

aopearance of the recrystallized reaction zone (Figure 18). 

since ST300 wire is a powder metallurgy product, it may contain 

enough porosity to accomodate the influx of mass from the matrix 

(Bradt et a1 1966). 

Also. 

Do = original fiber diameter 

s = reaction zone width 

D = unreacted fiber diameter ( = Do - 2.u ) 

Schematic of 

reacted fiber in 

a matrix. 

In addition, we assume that initially, u, = uf = 0.5, and that L'. = 

0. If we assume a single fiber in Area A, the initial volume 

fraction of matrix is, 

70 
This expression equals 0.5 since the volume fraction of the matrix 

remains constant. This can be solved to find an expression for A, 

no: 

.-i = - 2 

The volume fraction of fiber is equal to, 

n (D, - 2s)z 

Uf - - - -(0,,-2.~)~ = - lKD2 

4 .-t 4 .-i 20; 

and the volume fraction of. the interface zone is, 

L 

71 
: 

2 

j 'neasured, and the creep strength of the reaction zone is varied' 

1 'between that of the fiber and the matrix by the variable R' 

'This program calculates the creep strength of a W/Nb composite using' 

'the rule of mixtures. The creep strength of the fiber and matrix are' 

10 for R = 0 to 1. step .5 

!5 'Cycle temperature to generate data points' 20 for T = 1000 to 2000 step 100 'T is temperature in K' 

25 'Set constants' 

30 let D = 200 . '0 is original fiber diameter' 

40 let tm = 100 ' 'tm is time in hours' 

35 'Calculate the width of the reaction zone' 

50 let A = exp (-0.5 * (37390/T + 3.869)) 

60 let x = A tm"0.5 * 6 * lO"5 

65 'Calculate the volume fractions of matrix, fiber, and zone' 

70 let vm = .S 

80 let vf = ((0 - 2*x)"2)/(2*DA2) 

90 let vz = ((D"2 - ((D - 2*x)"2)))/(2*DA2) 

95 'Calculate creep strength of matrix and fiber' 

100 let sm = (I/tm)*.24 * 3.4 * lo"(-5) * exp (20384/T) 

110 let sf = (l/tm)".152 * 48.2 * exp (4153/T) 

115 'Estimate the creep strength of the reaction zone' 

116 'If R=O, sz = sm; if R=l, sz = sf' 

120 let sz = sf * R + sm * (1.-R) 

125 'Calculate the creep strength of the composite' 

130 let sc = sm * vm + sf * vf + sz * vz 

140 let y = SOOOO/T - log(tm) 

150 let x = log(sc) 

160 print T,tm,R,y,x 

190 next T 

200 next R 

72 
t ernp time ratio 

1000 

1100 

1200 

1300 

1400 

1500 

1600 

1700 

1900 

2000 

1000 

1100 

1200 

1300 

1400 

1500 

1600 

1700 

1900 

2000 

1000 

1100. 

1200 

1300 

1400 

1500 

1600 

1700 

1900 

2000 

1 aoo 

1 aoo 

1800 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

100 

0 

0 

0 

0 

0 

0 

0 

0 

0 

0 

0 

0.5 

0.5 

0.5 

0.5 

0.5 

0.5 

0.5 

0.5 

0.5 

0.5 

0.5 

1 

1 

1 

1 

1 

1 

1 

1 

1 

1 

1 

Y 

45.39483 

40.84937 

33,85637 

28.72816 

24 .a0659 

20.39483 

45.39483 

33. 85637 

28.72816 

26.64483 

37 .0615 

31.10912 

26.64483 

23.17261 

21.71062 

40.84937 

37.0615 

31.10912 

24.80659 

23.17261 

21.71062 

20.39483 

45.39483 

37.0615 

33.85637 

31.10912 

26.64483 

24.80659 

23.17261 

21.71062 

40. 84937 

28.72816 

20.39483 

A 

a. 470425 

7.048043 

5.786394 

5.209218 

6.24327 

5.473395 

4.941931 

4.622714 

4.177294 

3.441844 

1.78867 

a. 47038 

7.04801 

5.485968 

4.7891 15 

3.927832 

a. 470335 

6.243983 

5.790243 

5.242216 

5.018569 

4.532011 

4.234017 

7.047977 

6.244695 

5.794077 

5.274159 

5.089749 

4.931736 

4.670425 

5.498384 

4.793318 

4.560298 

C 

73 
appendix D. ExtraDo iation of Fiber Stress -RuDture Da ta 

?he squation for the minimum creep rate line on a typical creep 

curve is 

y=rnx+b or 1 -(rncr)(t,06)-b 

In addition, it has been determined from the experimental creep 

curves for ST300/Nb and ST300/Nb-lZr (values in Table X) that 

b = 0.292. Therefore, 

(mcr)(t = 1 - 0.292 = 0.708 

In addition, Monkman and Grant (1956) found that, 

(mcr)(t,) = 0.036 

0.708 0.036 

t 1% ti? 

-=- .. 

In addition, the approximate relationship between the time-torupture, applied stress, and temperature for the ST300 fibers was 

found to be (Eqn 7), 

Z7WO 

6.58 -- t, = (8.42~ 10-'*)a e ' 

74 
1168 

!Je now have an approximate relationship between the normalized 

applied stress and the time-to-1%-strain for the ST300 fibers at 

1500 K. 

i 

75 
cw\s/\ 

National Aeronautics and 

17. Key Words (Suggested by Author(@) 

Interdiffusion; Creep; Matrix-alloy effects; 

Fiber orientation; Angle ply; Refractory metals 

Report Documentation Page 

18. Distribution Statement 

Unclassified - Unlimited 

Subject Category 26 

DOE Category UC-25 

SDace Administration 

19. Security Classif. (of this report) 20. Security Classif. (of this page) 

Unclassified Unclassified 

2. Government Accession No. NASA TM- 102 122 1. Report No. 

21. No of pages 22. Price' 

76 A03 

DOE/NASA/ 163 10- 10 

4. Title and Subtitle 

The Interface in Tungsten Fiber Reinforced Niobium 

Metal-Matrix Composites 

7. Author(s) 

Toni L. Grobstein 

9. Performing Organization Name and Address 

National Aeronautics and Space Administration 

Lewis Research Center 

Cleveland, Ohio 44135-3191 

12. Sponsoring Agency Name and Address 

National Aeronautics and Space Administration 

Washington, D.C. 20546-0001 

3. Recipient's Catalog No. 

5. Report Date 

September 1989 

6. Performing Organization Code 

8. Performing Organization Report NO. 

E-4754 

10. Work Unit No. 

586-01-1 1 

11. Contract or Grant No. 

13. Type of Report and Period Covered 

Technical Memorandum 

14. Sponsoring Agency Code 

15. Supplementary Notes 

Final Report. Prepared under Interagency Agreement DE-AI03-86SF 163 10. This report was submitted by 

Toni L. Grobstein as a thesis in partial fulfillment of the requirements for the degree Doctor of Philosophy to 

Case Western Reserve University, Cleveland, Ohio. 

16. Abstract 

The creep resistance of tungsten fiber reinforced niobium metal-matrix composites was evaluated. The interface 

region between the fiber and matrix was characterized by microhardness and electron probe microanalysis 

measurements which indicated that its properties were between those of fiber and matrix. However, the measured 

properties of the composite exceeded those calculated by the rule of mixtures even when the interface zone was 

assumed to retain all the strength of the fiber. The composite structure appeared to enhance the strengths of both 

the fibers and the matrix above what they exhibited in stand-alone tests. The effect of fiber orientation and matrix 

alloy composition on the fibedmatrix interface were also evaluated. Small alloying additions of zirconium and 

tungsten to the niobium matrix affected the creep resistance of the composites only slightly. A decrease in the 

creep resistance of the composite with increasing zirconium content in the matrix was ascribed to an increase in 

the diffusion rate of the fibedmatrix interdiffusion reaction, and a slight increase in the creep resistance of the 

composite was observed with an addition of 9 w% tungsten to the matrix. In addition, Kirkendall void formation 

was observed at the fibedmatrix interface; the void distribution differed depending on the fiber orientation 

relative to the stress axis. 

NASA FORM 1626 OCT 86 *For sale by the National Technical Information Service, Springfield, Virginia 221 61