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.
Fig. 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). Fracture toughness values in excess of 20 MPa√m have been achieved (Metall Mater Trans A 36A (2005)). However, these toughness values come at the expense of oxidation resistance.
(a) (b) (c)
Fig. 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 (Kruzic alloy). (b) "duplex" microstructure with ~50% vol α-Mo phase produced via mechanical alloying (ULTMAT alloy). (c) duplex microstructure produced via reaction synthesis (Middlemas alloy).
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.
processing routes, such as the mechanical alloying route
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
and 50vol% intermetallic phases in a "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
At room temperature, both the ULTMAT and Middlemas alloys have much lower initiation toughnesses (7.8 MPa√m and 7.2 MPa√m) than the comparable Kruzic alloy (12.5 MPa√m). Though the volume fraction of ductile α-Mo is similar (and in fact higher for the newer alloys), neither the ULTMAT nor the Middlemas alloy displayed any stable crack growth. A number of factors reduced the damage tolerance of these alloys. The ultra-fine grains of the ULTMAT and Middlemas alloy did not prove to be a significant impediment to crack advance. Though a large volume fraction of α-Mo would imply a very high probability of the more ductile grains interacting with a moving crack and trapping it, the extremely small grain size provides a pathway by which a crack can avoid the more ductile grains without a large increase in energy. Plastic constraint of the ductile phase by the high volume fraction of hard intermetallic particles also reduced the damage tolerance of these alloys. The much harder intermetallic particles prevented the ductile α-Mo grains from plastically deforming. Instead, the high stresses which developed led to fracture of the a-Mo grains or the surrounding grain boundary material. This effect was exacerbated by segregation of Si into the grain boundaries, reducing their strength. Figs. 3a-c show the location of Si impurities of the fracture surface of the ULTMAT, Middledmas and Kruzic alloys, respectively. Large amounts of Si appear on the grain boundaries, while almost none is present in grain interiors (as depicted by grains which have fractured transgranularly. The amount of Si within the grain boundaries of the ULTMAT and Middlemas alloys was much higher than the Kruzic alloy, as the processing method used to create the Kruzic alloy removes Si in the material that becomes the grain boundaries upon sintering. The critical point here is that at low temperatures, Mo-Si-B alloys are truly brittle materials as the α-Mo phase can only provide for very limited ductility. Brittle materials can only be toughened extrinsically, and as such the coarser microstructures are able to generate toughness (more precisely crack-growth resistance) through such shielding processes as crack deflection and ductile-ligament bridging.
Fig. 3: Auger electron spectroscopy maps of impurity content on grain boundaries overlaid on the corresponding room-temperature fracture surfaces for the (a) Kruzic (b)ULTMAT and (c) Middlemas alloys. Areas of high silicon content (blue) are shown. Si segregates to grain boundaries, reducing interfacial strength and increasing the occurrence of intergranular fracture. Note the high concentrations of Si in the regions that fractured intergranularly, while almost no Si is found in regions that fractured transgranularly. Note the difference in scale for (a) as the Kruzic alloy had grains more than one order of magnitude larger than either (b) or (c).
At elevated temperature (1300°C), the alloys are above their ductile-brittle transistion temperature. As a result, the α-Mo phase is much more ductile and can readily plastically deform. Here, the volume fraction of ductile α-Mo becomes more important than its morphology. As a result, the initiation toughnesses of the ULTMAT, Middlemas and Kruzic alloys converged. All three alloys were so ductile that the underlying assumptions of linear-elastic fracture mechanics were violated and J-integrals were estimated based on optical measurement of crack-tip opening displacements. Fig. 4 plots toughness versus temperature for the ULTMAT, Middlemas and Kruzic alloys, as well as the much higher Si content alloys studied by Choe, et al. (Metallurgical and Materials Transactions, 2003, 34A) which consisted of discontinuous islands of α-Mo in an intermetallic matrix.
Fig. 4: Fracture toughness
as a function of temperature for Mo-Si-B alloys.
Crack-initiation toughnesses (closed symbols) are plotted along
with any increases in toughness with crack extension (open
symbols). The highest room temperature toughness value for
the Kruzic alloy was obtained after more than 3 mm of stable
crack growth. At low temperatures, neither the ULTMAT
(mechanically alloyed) nor the Middlemas (reaction synthesized)
alloy exhibited any stable crack growth prior to unstable
fracture. The ductile-brittle transition temperature for
these materials is ~1000°C, so only moderate gains in initiation
toughness are expected below this temperature, as demonstrated
by the ULTMAT alloy. At 1300°C, the enhanced ductility of
the α-Mo phase markedly improves the initiation toughness of
alloys containing ~50 vol.% α-Mo. At this temperature, the
volume fraction of α-Mo becomes a more important factor in
developing toughness (intrinsically from plasticity) than the
distribution and morphology of α-Mo grains (which leads to
extrinsic toughening from mechanisms such as crack bridging).
The design and development of new materials for ultrahigh temperature applications is invariably a competition between achieving excellent oxidation resistance and creep strength at service temperatures and maintaining adequate ductility and toughness at both low and high temperatures. Unfortunately, the microstructural requirements to achieve acceptable behavior in all three categories are generally mutually exclusive. This is a particularly difficult problem in Mo-Si-B alloys where the microstructures for optimal oxidation resistance, creep strength and damage tolerance (strength and toughness) are so contradictory. Specifically, for oxidation resistance, the three-phase alloys with very small discontinuous grains are best as the small grains limit the probability that an α-Mo grain will be exposed to oxygen; likewise, the small grains provide a short diffusion pathway allowing for faster passivation than in coarser-grained alloys. In direct contrast, optimal room-temperature damage tolerance is afforded by large, continuous α-Mo grains that promote extrinsic toughening by the generation of ductile ligament bridges that act to “shield” a crack tip from the full force of an applied stress, thereby inhibiting crack advance. Corresponding high-temperature toughness is also promoted by a high volume fraction of α-Mo as the ductility of this phase generates extensive plasticity which toughens the alloy intrinsically. In further contrast, optimal creep response is provided by alloys with large intermetallic grains surrounding small islands of α-Mo, which limits the number of high-diffusivity pathways such as grain boundaries; a low volume fraction of α-Mo is also desirable, as the relative ease of deformation of α-Mo at high temperatures allows individual intermetallic particles to rearrange easily. Fig. 5 shows schematic illustrations of the microstructural morphologies necessary to maximize material response for each property. Steps must be taken to limit segregation of Si in these alloys and improve the room temperature ductility of α-Mo. With increased α-Mo ductility, lower volume fractions of this phase can acheive comparable toughnesses, thereby improving the oxidation and creep resistance of Mo-Si-B alloys.
Fig. 5: Schematic illustrations of the ideal microstructures to improve oxidation resistance, creep resistance and damage tolerance of Mo-Si-B alloys. The morphological considerations for improvement in each area are mutually exclusive, so optimization of the properties of each phase is necessary.
J. K. Cochran (collaborator from Georgia Tech)
T. Weingärtner (collaborator
from Karlsruhe Institute of Technology)
R. O. Ritchie
Work on this project conducted at