One
of the most difficult problems in materials engineering today is the
development of higher temperature structural materials for use in
applications
such as gas-turbine engines. The current material of choice,
single-crystal nickel-based superalloys, has reached its technological
limit;
indeed, as these alloys melt at temperatures between 1200°-1400°C, they
are
unsuitable for structural use above ~1100°C (Fig. 1). High
melting-point (>2000°C) materials, based on refractory metals such
as
molybdenum, represent a higher-temperature alternative but have been
plagued by
oxidation and brittleness problems. Additions of silicon
and boron
to molybdenum to form silicides and borosilicides have shown promise in
improving the oxidation resistance; however, the silicide compounds are
quite
brittle and will provide little fracture resistance for most structural
applications without significant active toughening mechanisms.
While
several alloys have been produced containing the more ductile α-molybdenum
phase in addition the hard but brittle intermetallic phases Mo3Si
and Mo5SiB2 (the T2 phase), many of these show
only
marginal improvements in toughness relative to the monolithic
intermetallic
phases which have fracture toughnesses of ~ 3 - 4 MPa√
m (e.g., Metal.
Mater. Trans.,
34A, 2003, p. 225). This suggests that the key to
achieving high
fracture resistance in these materials may lie in making more effective
use of
the "ductile" α-Mo phase, in a manner not unlike
the way that nickel-based (γ-γ')
superalloys obtain high fracture toughness
with a similarly
high fraction of intermetallic (γ') precipitates.

Figure 1 Plot showing the improvement since 1940 in the temperature capacity of metallic alloys, specifically nickel-based superalloys, for gas-turbine engine applications and demonstrating the need for new materials, such as molybdenum based superalloys, in order to achieve further technological gains.
Accordingly,
to achieve improved fracture resistance, the approach currently being
undertaken is to develop alloys where the intermetallic phases are
completely
surrounded by a continuous "ductile" α-Mo
phase. Molybdenum-based alloys
have been processed with Mo3Si and T2 particles in a
continuous α-Mo
matrix using a novel powder processing route. Specifically, to
obtain the
continuous α-Mo phase, ground powders of Mo3Si
and T2 phase
(composition Mo-20Si-10B at%) were vacuum annealed to remove silicon
from the
surface and leave a α-Mo coating on each
particle. These surface-modified
powders were then hot isostatically pressed to achieve alloys with a
continuous
α-Mo
matrix, reinforced by the intermetallic phases, Mo3Si and T2
(Fig.
2a). Full processing details may be found in (Scripta Mater.,
46,
2002, p. 217). Preliminary results have indicated fracture
toughness
values in excess of 20 MPa√m
may be achieved, while
future work will focus on understanding the specific relations between
the
microstructure and the fracture and fatigue behavior of these alloys.


Figure
2 (a) Microstructure of a Mo-Si-B alloy with a continuous α-Mo
matrix (~ 46% vol.) produced by the surface modified powder metallurgy
method.
(b) "Quasi-continuous duplex" microstructure with 50% vol α-Mo
phase produced using the ULTMAT process.
The
focus of this work is to make a significant advance in the development
of
Mo-Si-B alloys, specifically by tailoring the composition, morphology
and
volume fraction of the major phases of these alloys (α-Mo,
MoSi3 and Mo5SiB2
(T2)) to achieve an optimum balance of low- and high-temperature
damage-tolerance with creep and oxidation resistance. Unlike
Mo-Si-B
materials based entirely on intermetallic compounds, these alloys
contain the
metallic α-Mo
phase which provides some degree of fracture resistance and ductility.
Furthermore, the silicide and borosilicide phases provide creep and
oxidation
resistance, the latter of which is the result of a borosilicate glass
scale
which forms in situ on the metal surface.
Newer
processing routes, developed as part of ONERA's ULTMAT program, produce
alloys
with compositions near those of the work above, namely
Mo-3wt%Si-1wt%B.
The resulting alloys consist of roughly 50vol% α-Mo
and 50vol% intermetallic
phases in a "quasi-continuous duplex" microstructure. See Figure
2b. Grains in this material are much smaller (15-20 mm) than previous materials.
Nanoscale
yttria particles have been dispersed in the grain boundaries to limit
grain
growth and impede creep. See Jehanno et. al., Materials Science and Engineering A 463
(2007) pgs. 216-223 for full processing details.
A third
processing
route, described by Middlemas and Cochran, JOM July (2008) pg
19-24,
will also be studied. The microstructure of these Mo-3wt%Si-1wt%B
alloys
is nearly identical to that of the ULTMAT materials.
Preliminary results show initiation toughnesses on the order of 8.5 MPa√m for the ULTMAT material and 7.2 MPa√m for the Middlemas & Cochran (GT material), with no stable crack growth apparent for either material. Compare these results with the results of Kruzic, et. al., Metall Mater Trans A 36A (2005), where similar volume fractions of α-Mo produced initiation toughnesses up to 12 MPa√m and significant rising R-curve behavior. The difference in behavior becomes more pronounced at higher temperatures. Consult Figure 3 below.

Figure
3: Initiation toughness for ULTMAT (blue) GT (red) and Kruzic, et. al. materials. The ULTMAT
and GT
materials show no stable crack growth, while Kruzic et. al.
were able to
grow cracks stably for several millimeters, even at room
temperature.
They achieved room temperature toughnesses above 20 MPa√m.
Fractography of the fracture surfaces and crack path-microstructure studies will be performed to better understand the origin of these disparate crack growth behaviors.
J. A. Lemberg
J. H.
Schneibel
(collaborator from ORNL)
M.
R. Middlemas (collaborator from Georgia Tech)
J. K. Cochran (collaborator from Georgia Tech)
Work on this project conducted at