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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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