High temperature strength and room temperature fracture toughnessof Nb�/Mo�/W refractory alloys with and without carbide dispersoids

Won-Yong Kim a,*, Hisao Tanaka a, Mok-Soon Kim b, Shuji Hanada c

a Japan Ultra-High Temperature Materials Research Institute, 573-3 Okiube, Ube-city, Yamaguchi 755-0001, Japanb School of Material Science and Engineering, Inha University, Inchon 402-751, Republic of Korea

c Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan

Received 19 March 2002; received in revised form 12 June 2002

Abstract

Room temperature fracture toughness and high temperature strength at 1773 K of Nb�/Mo�/W based alloys with and without

carbides were investigated by three-point bending and compression tests. W addition to Nb�/5Mo solid solution gives rise to an

increase in yield stress at 1773 K but to a decrease in fracture toughness at room temperature. On the basis of the experimental

results obtained, it is suggested that solid solution hardening and cleavage fracture propensity are responsible for the increase of

yield stress at 1773 K and for the decrease of fracture toughness at room temperature, respectively. The microstructure of Nb�/

5Mo�/15W�/x (Hf�/C) alloys consists of Nb-rich bcc solid solution and (Nb,Hf,Mo)C carbide with B1-structure. Yield stress at 1773

K and fracture toughness at room temperature increase concurrently with increasing content of (Hf�/C) in the Nb�/5Mo�/15W�/

x (Hf�/C) alloys. The increase of yield stress at 1773 K due to the presence of (Nb,Hf,Mo)C phase is explained by the mechanism

based on dispersion hardening, and the increase of fracture toughness at room temperature is attributable to the transition of

fracture mode from cleavage to quasi-cleavage accompanied by decohesion between carbide and matrix phase, crack branching and

deflection. Details will be discussed in relation to microstructural characteristics.

# 2002 Elsevier Science B.V. All rights reserved.

Keywords: Niobium alloy; Fracture toughness; High temperature strength; Carbide; Solid solution strengthening; Dispersion hardening

1. Introduction

In recent years, advanced structural materials have

been strongly required for applications at temperatures

above the maximum operating temperature of conven-

tional high temperature engineering materials such as

Ni-base superalloys [1�/18]. Additionally, from the

viewpoint of environmental problems, an increase in

operating temperature of electric power generators or jet

turbines is indispensable, because it can lead to high

energy efficiency, thereby reducing the emission of

harmful gases such as CO2 and NOx [3,6]. Under such

situations, Nb-based refractory alloys with a higher

melting point and a relatively low density have been

expected as candidate materials to use at ultra-high

temperature over the maximum operating temperature

of Ni-base superalloys. A drastic decrease of high

temperature strength, however, is one of the subjects

to be improved for practical application of these bcc-

based materials. Most studies have focused on the

increase in high temperature strength on the basis of

the concepts of solid solution strengthening and disper-

sion strengthening [19�/23]. In the present study, we

have considered the good balance of high temperature

strength and room temperature fracture toughness by

alloying and modification of microstructure. We chose

Mo and W as solid solution strengthening elements

because of their large elastic and modulus interactions

with Nb. To further improve mechanical properties, we

tried to introduce carbide particles in the Nb-rich solid

solution. Because of the high thermal stability and

strength at high temperature, Hf and C were added to

Nb-rich solid solution in expectation of the formation of

(Nb,Hf,Mo)C.

* Corresponding author. Tel.: �/82-32-5707-133; fax: �/82-32-5707-

102

E-mail address: kim@jutem.co.jp (W.-Y. Kim).

Materials Science and Engineering A346 (2003) 65�/74

www.elsevier.com/locate/msea

0921-5093/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved.

PII: S 0 9 2 1 - 5 0 9 3 ( 0 2 ) 0 0 5 1 5 - 4

2. Experimental procedure

Raw materials used in this study were 99.9 wt.% Nb,

99.9 wt.% Mo, 99.9 wt.% W, 99.999 wt.% Si and 98wt.% Hf. Carbon was used in a form of Nb2C with 99

wt.% purity. Alloy buttons were prepared by arc melting

in an argon atmosphere on a water-cooled copper

hearth using a non-consumable tungsten electrode.

The buttons were re-melted several times to remove

segregation and to enhance the chemical homogeneity.

Heat treatment was carried out at 2073 K for 24 h in an

argon atmosphere followed by rapid furnace cooling.Some of the samples were also heat treated at 2273 K to

observe the microstructural change, particularly for

carbide. Samples for metallographic observation, che-

mical composition analysis, phase identification and

mechanical testing were prepared by electro-discharge

machining (EDM). Microstructural observation was

carried out using back scattered electron images (BEI)

in a scanning electron microscope (SEM) to identifyconstituent phases by contrast difference. X-ray diffrac-

tion (XRD) was performed on the heat-treated samples

to examine the crystal structure of constituent phases.

Compression specimens with 3�/3 cross section and 6

mm in height and fracture toughness specimens with

3�/6 mm cross-section and 24 mm span in three-point

bending were prepared by EDM. Specimens were then

mechanically polished with SiC paper and fine Al2O3

particles with water. Compression tests were carried out

at 1773 K in an argon atmosphere using an Instron-type

testing machine. Fracture toughness tests were carried

out in three point bending according to ASTM E399-

1987 testing method without insertion of a fatigue pre-

crack. An equation form KQ�/(PqS /BW3/2)f(a /W ),

where a /W is 0.45�/0.55, P is load, B is thickness, W

is height, S is applied load span and a is notch depth,was used to calculate the fracture toughness value.

Fracture toughness test was performed in air at room

temperature and a cross head speed of 0.5 mm min�1.

Fracture surface observation was done using SEM at the

operating voltage of 15 kV. Crack propagation under

monotonic loading was observed. Cracks were captured

using a sound signal that allows us to detect crack

nucleation and growth, and then samples are unloadedas quickly as possible. Careful attention was paid to

avoid a premature failure, since crack growth rate was

estimated to be very high.

3. Results and discussion

3.1. Microstructure

Typical microstructure of Nb-5at.%Mo-5at.%W

(hereafter denoted as Nb�/5Mo-5W) alloy annealed at

2023 K for 24 h is presented in Fig. 1. Equiaxed grains

with an average grain size of 150 mm are observed and

there is no second phase or unexpected inclusion,

indicating Mo and W are soluble completely in the Nb

solid solution. BEI microstructures of Nb�/5Mo�/15W�/

x (Hf�/C) alloys annealed at 2023 K for 24 h are shown

in Fig. 2, where x is atomic percent. In Nb�/5Mo�/15W

alloy containing 2.5Hf and 2.5C, which will be denoted

as Nb�/5Mo�/15W�/2.5(Hf�/C) hereafter, finely and

homogeneously dispersed carbide particles are found

in the Nb solid solution. The shape of carbide particles is

observed to be needle-like as shown in Fig. 2(a). With

increasing content of (Hf�/C), the volume fraction of

carbide increases as shown in Fig. 2(b), (c) and (d). In

addition, the needle-like carbide observed in the alloy

with 2.5(Hf�/C) changes to a mixture of needle-like and

blocky shapes with also increasing (Hf�/C) content. To

understand whether this microstructural change is only

attributable to the compositional dependence of (Hf�/

C) or there exists other reason such as solidification

process, we also observed the microstructures for as-cast

alloys. Interestingly, we could not find needle-like

carbide through the whole as-cast alloys studied.

Further, it appears that the microstructure is featureless

in the as-cast Nb�/5Mo�/15W�/2.5(Hf�/C) alloy. This

result indicates that the blocky carbide is produced

during solidification, probably as a primary type, in

alloys with higher (Hf�/C) content, while the needle-like

carbide is created by precipitation during annealing.

TEM micrographs of precipitated NbC at x�/2.5 and 5

are shown in Fig. 3. The results on EDX analysis taken

from matrix and carbide in thin foils of Fig. 3 are shown

in Fig. 4. Evidently, Mo and W are included in Nb solid

solution, whereas Hf, Nb and Mo are included in

carbide, forming (Nb,Hf,Mo)C.

Fig. 1. Optical micrographs of Nb�/5Mo�/5W alloy annealed at 2023

K for 24 h.

W.-Y. Kim et al. / Materials Science and Engineering A346 (2003) 65�/7466

Fig. 2. BEI micrographs of Nb�/5Mo�/15W�/x (Hf�/C) alloys annealed at 2023 K for 24 h, where x is atomic percent; (a) x�/2.5, (b) x�/5, (c) x�/

10 and (d) x�/15.

Fig. 3. TEM bright field images of (a) Nb�/5Mo�/15W�/2.5(Hf�/C) and (b) Nb�/5Mo�/15W-5(Hf�/C).

W.-Y. Kim et al. / Materials Science and Engineering A346 (2003) 65�/74 67

Fig. 4. TEM-EDX results on (a) matrix and (b) carbide in Nb�/5Mo�/15W�/2.5(Hf�/C), and (c) matrix and (d) carbide in Nb�/5Mo�/15W-5(Hf�/C).

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3.2. X-ray diffraction and lattice parameter

The XRD peaks for four W-added Nb�/5Mo alloys

without carbide were indexed to be bcc phase and no

additional peaks were detected. The diffraction peak

positions were found to shift more significantly to

higher 2u side with increasing W content, indicating a

decrease in lattice parameter. Measured lattice para-

meters are plotted as a function of W content in Fig. 5.

The XRD profiles for three Nb�/5Mo�/15W�/x (Hf�/C)

series alloys are shown in Fig. 6, where x is 2.5, 10 and

15. It is very evident that many peaks other than bcc

peaks are seen in the alloys containing (Hf�/C). These

extra peaks appear in a similar manner, and they are

confirmed to be (Nb,Hf,Mo)C with B1 structure. As

seen in the peaks near 348 and 578 of 2u , the diffraction

peak intensity from (Nb,Hf,Mo)C increases with in-

creasing (Hf�/C) content in agreement with the micro-

structural observation in Fig. 2. No carbide or oxide

other than (Nb,Hf,Mo)C is detected in the composition

range investigated. The lattice parameters of Nb-rich

bcc solid solution and B1-structured (Nb,Hf,Mo)C

phase are plotted in Fig. 7 as a function of (Hf�/C)

content in the Nb�/5Mo�/15W�/x (Hf�/C) alloys. No

distinct change is seen for the bcc phase, suggesting that

(Mo�/W) content in Nb solid solution does not change

significantly. In contrast, the lattice constant of

(Nb,Hf,Mo)C increases with increasing (Hf�/C) con-

tent, implying the increase of Hf content in

(Nb,Hf,Mo)C, since Taylor and Doyle found that the

lattice parameter of (Nb,Hf,Mo)C carbide increases

with increasing Hf content [27]. It is apparent, therefore,that Hf is partitioned more dominantly in (Nb,Hf,Mo)C

than in Nb-rich bcc phase, which is consistent with the

results in Fig. 6. The obtained results are consistent with

the Nb�/Hf�/C phase diagram at 2273 K presented by

Taylor and Doyle, where NbC and HfC forms a

continuous solid solution, and (Nb,Hf,Mo)C is equili-

brated with Nb solid solution [27].

3.3. Mechanical properties and fractography

Fig. 8 shows room temperature fracture toughness

and 0.2% offset yield stress at 1773 K as a function of W

content in Nb�/5Mo-xW alloys. The yield stress at 1773

K increases rapidly with increasing W content, while

room temperature fracture toughness decreases notice-

ably. The increase of yield stress at 1773 K will be

closely associated with solid solution strengthening anddiffusivity of W, if we assume that a difference in grain

size is negligible. Indeed, the grain size of the alloys

without (Hf�/C) addition was measured to be about 150

mm, which was not sensitive to W content. As shown in

Fig. 5, the lattice parameter of Nb-rich solid solution

decreases with increasing W content in Nb�/5Mo-xW

alloys. Therefore, solid solution strengthening due to

elastic interaction arising from the atomic size differencemay be one of the important strengthening mechanisms

at 1773 K in the single-phase bcc alloys. In addition, W

is well known to have low diffusivity in Nb solid

solution [31], which retards the deformation controlled

by a thermally activated process and thereby high

temperature strength will be increased. The decrease of

room temperature fracture toughness can be closely

related to the fact that Nb is a remarkably solid solutionstrengthened by Mo and W alloying, thereby decreasing

ductility [28�/30]. Fig. 9 shows the fracture surfaces

taken from (a) Nb�/5Mo-5W and (b) Nb�/5Mo�/15W

alloys heat-treated at 2023 K for 24 h. In the Nb�/5Mo-

5W alloy, a mixed fracture mode consisting of inter-

granular fracture and transgranular one is observed on

the whole fracture surface, while in the Nb�/5Mo�/15W

alloy, fracture occurs mostly in a transgranular cleavagemanner. Therefore, the difference in fracture toughness

between Nb�/5Mo-5W and Nb�/5Mo�/15W alloy would

be explained in terms of the fracture mode. The

remarkable solid solution strengthening by W addition

will reduce plastic zone at a crack tip, thereby enhancing

cleavage fracture and leading to the decrease in fracture

toughness. Fig. 10 shows the room temperature fracture

toughness and 0.2% offset yield stress at 1773 K as afunction of (Hf�/C) content in the Nb�/5Mo�/15W�/

x (Hf�/C) alloys. The yield stress increases rapidly by

(Hf�/C) addition of 2.5 at.% and then increasesFig. 5. Lattice parameters of W-added Nb�/5Mo alloys as a function

of W content.

W.-Y. Kim et al. / Materials Science and Engineering A346 (2003) 65�/74 69

Fig. 6. XRD profiles for Nb�/5Mo�/15W alloyed with (Hf�/C).

Fig. 7. Lattice parameters of Nb solid solutions and (Nb,Hf,Mo)C

carbide plotted as a function of (Hf�/C) content in Nb�/5Mo�/15W�/

x (Hf�/C) alloys. All samples were heat treated at 2023 K for 24 h.

Fig. 8. Relationship between room temperature fracture toughness

and 0.2% offset yield stress at 1773 K plotted as a function of W

content in Nb�/5Mo-xW alloys.

W.-Y. Kim et al. / Materials Science and Engineering A346 (2003) 65�/7470

gradually with further increasing content of (Hf�/C).

Interestingly, the room temperature fracture toughness

increases with increasing (Hf�/C) content up to 5% and

then no distinct change is observed. Fig. 11(a), (b), (c)

and (d) shows fracture surfaces of Nb�/5Mo�/15W�/

x (Hf�/C) alloys at x�/2.5, 5, 10 and 15, respectively.

In Nb�/5Mo�/15W�/2.5(Hf�/C), cleavage fracture sur-

faces with river patterns and traces showing decohesion

between matrix and carbide are observed. This decohe-

sion will be derived from debonding at needle-like

carbide (see Fig. 2(a)). With further increasing (Hf�/

C) content, cleavage fracture becomes less dominant,

while decohesion is often observed depending on volume

fraction of carbide, as shown in Fig. 11(b), (c) and (d).

In Nb�/5Mo�/15W�/x (Hf�/C) series alloys, the relation-

ship between yield stress at 1773 K and room tempera-

ture fracture toughness as a function of (Hf�/C) content

appears to be quite different from that as a function of

W content in Nb�/5Mo�/15W alloy. It is very interesting

to note that 0.2% offset yield stress and room tempera-

ture fracture toughness increase simultaneously with

increasing (Hf�/C) content up to 5 at.%. The increase of

yield stress may be interpreted in terms of both solid

solution strengthening and dispersion (or precipitation)

strengthening. However, no apparent change is seen in

the lattice parameter of Nb solid solution as shown in

Fig. 7, suggesting that W�/Mo content is not signifi-

cantly changed in the matrix. Moreover, Hf is not solved

in Nb solid solution as shown in Fig. 4. Therefore, solid

solution strengthening is ruled out. Thus, the increase of

yield stress will be explained by dispersion strengthening

of carbide based on Orowan mechanism [24�/26]. It is

well known that the strengthening by dispersoids

depends on volume fraction (f), size (r) and mean

interparticle spacing (l) of particles, where l is usually

expressed as a function of r and f , l�/4(1�/f)r /3f . Here,

it should be noted that the yield stress of (Hf�/C)-added

alloys increases by about 80 MPa with increasing (Hf�/

C) content from 0 to 2.5%, while it increases by about 12

MPa with increasing (Hf�/C) content from 2.5 to 5

at.%. The most distinguishable difference in microstruc-

ture between 2.5 and 5 at.% for x is in the morphology

of carbide; fine needle-like carbide dispersoids in the

Nb�/5Mo�/15W�/2.5(Hf�/C) and needle-like and blocky

carbide dispersoids in the Nb�/5Mo�/15W-5(Hf�/C).

Kelly et al. suggested that rod- and plate-types of

particles strengthen dispersion hardening alloys about

twice as much as spherical particles at an equal volume

Fig. 9. SEM micrographs of fracture surface of (a) Nb�/5Mo-5W and (b) Nb�/5Mo�/15W.

Fig. 10. Relationship between room temperature fracture toughness

and 0.2% offset yield stress at 1773 K plotted as a function of (Hf�/C)

content in Nb�/5Mo�/15W�/x (Hf�/C) alloys.

W.-Y. Kim et al. / Materials Science and Engineering A346 (2003) 65�/74 71

fraction [24]. Therefore, the difference in increment of

yield stress mentioned above can be related to these

microstructural features. As shown in Fig. 7, the lattice

parameter of carbide increases with increasing (Hf�/C)

content, implying the substitution of Hf for Nb in NbC

phase. This result indicates that Hf is solved preferen-

tially in NbC phase. This alloying behavior is well

consistent with the ternary Nb�/Hf�/C phase diagram

constructed based on lattice parameter change and

microstructural observation [27].

Concerning the fracture toughness result, it is clearly

shown that (Hf�/C) addition to Nb�/5Mo�/15W im-

proves fracture toughness in this study. The presence of

carbide particles hinders straight crack propagation and

eventually leads to crack deflection and branching by

which fracture toughness is increased. As seen in Fig. 11,

the region showing cleavage fracture appearance is

reduced and the areas exhibiting decohesion at interface

between carbide and matrix or stretching of matrix are

observed more frequently with increasing content of

(Hf�/C). Thus, the increase of fracture toughness by

(Hf�/C) addition may be ascribed to the changes of

volume fraction and shape of carbide.

3.4. Crack propagation behavior

Crack propagation under monotonic loading is pre-

sented in Fig. 12(a) and (b) for Nb�/5Mo�/15W and

Nb�/5Mo�/15W-5(Hf�/C), respectively. In Nb�/5Mo�/

15W, crack propagates straight without showing any

shear deformation, crack deflection, branching or ar-

resting and renucleation as shown in Fig. 12(a). In

contrast, crack deflection and crack branching are seen

in Nb�/5Mo�/15W-5(Hf�/C). This is likely due to the

presence of carbide particles, which hinder straight

crack propagation. Therefore, these differences in crack

propagation behavior may result in the difference of

fracture toughness between both alloys, suggesting that

carbide distribution in the Nb-rich solid solution alloy is

Fig. 11. SEM micrographs of fracture surface of (a) Nb�/5Mo�/15W�/2.5(Hf�/C), (b) Nb�/5Mo�/15W-5(Hf�/C), (c) Nb�/5Mo�/15W-10(Hf�/C) and

(d) Nb�/5Mo�/15W-15(Hf�/C). All samples are arc-melted and annealed at 2073 K for 24 h.

W.-Y. Kim et al. / Materials Science and Engineering A346 (2003) 65�/7472

beneficial to improve the fracture toughness in the

present alloy system.

3.5. High temperature strength and room temperature

fracture toughness

In Fig. 13, we summarize the relationship between

0.2% offset yield stress and room temperature fracture

toughness in this study. In the Nb-rich solid solution

alloys, fracture toughness is decreased rapidly with

increasing yield stress at 1773 K, indicating that solid

solution hardening by Mo and W addition leads to the

remarkable decrease of fracture toughness. In contrast,

in the dispersion-hardened alloys, yield stress at 1773 K

is increased without sacrificing fracture toughness at

room temperature. It is concluded from the present

experiments that the incorporation of dispersoids having

high thermal stability and strength in a toughened

matrix is a promising technique to enhance both room

temperature fracture toughness and high temperature

strength in Nb base alloys.

4. Conclusions

High temperature strength at 1773 K and room

temperature fracture toughness of Nb�/Mo�/W based

refractory alloys with and without carbide were inves-

tigated as a function of (Hf�/C) and W content. The

results obtained are summarized as follows.

1. Microstructures of Nb-W-Mo-x (Hf�/C) alloys

consist of bcc Nb-rich Nb�/Mo�/W solid solution as a

matrix phase and (Nb,Hf,Mo)C carbide dispersoids

with B1-structure. With increasing content of (Hf�/C),

volume fraction of carbide increases and the shape of

carbide changes from a needle to a mixed type consisting

of needle and blocks.2. With increasing content of (Hf�/C), the lattice

parameter of Nb-rich solid solution alloys is left un-

changed, but that of (Nb,Hf,Mo)C increases in the Nb�/

5Mo�/15W�/x (Hf�/C) alloys.

3. W addition to Nb�/5Mo increases remarkably high

temperature strength at 1773 K, but decreases fracture

toughness at room temperature.4. Carbide dispersion increases simultaneously both

fracture toughness at room temperature and strength at

1773 K.

Fig. 12. BEI micrographs showing crack propagation in Nb�/5Mo�/

15W (a) and Nb�/5Mo�/15W-5(Hf�/C) (b).

Fig. 13. Relationship between yield stress at 1773 K and room

temperature fracture toughness for Nb�/5Mo-xW and Nb�/5Mo�/

15W�/x (Hf�/C) alloys.

W.-Y. Kim et al. / Materials Science and Engineering A346 (2003) 65�/74 73

Acknowledgements

This work is supported by a grant from the New

Energy and Industrial Technology Development Orga-nization (NEDO) of Japan.

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