Patent Application
20140286821
1. Field of the Invention
This invention relates to novel alloys and methods of producing the alloys. More specifically, the alloys are strong and ductile microstructured alloys having lamellar structures.
2. Description of the Related Art
Basic research on alloy materials seeks to find improved materials, such as those that are lighter, stronger or less expensive than conventional metals and alloys. In other contexts, improved materials may have increased resistance to weather, chemicals or friction in an intended environment of use. Equipment that incorporates these new materials in component parts may have a longer service life, require less maintenance or achieve an improved performance level. From a cost of manufacture standpoint, it is desirable for these new materials to be made from readily available and highly affordable natural resources.
One technique that may be used to produce an alloy with enhanced strength and ductility is a eutectic transformation. A eutectic transformation occurs when components of an alloy crystallize simultaneously from a liquid solution. Products of a eutectic transformation can often be identified by their lamellar structure where spacing between lamellae is typically on the order of less than a micron to a few microns. Such structures are generally strong and ductile. For example, the most well known lamellar material is carbon steel.
Alloys of the present disclosure advance the art by providing materials with exceptional strength and ductility.
In one embodiment, an intermetallic composition formed by a eutectic transformation in at least two distinct structural phases has an average composition comprising from 25% to 35% iron, 15% to 25% nickel, 30% to 40% manganese and 10% to 20% aluminum, where the composition is described in terms of atomic percentages.
In an embodiment, the invention provides an alloy comprising at least two distinct structural phases, wherein the average composition of the alloy comprises from 25% to 35% iron, from 15% to 25% nickel, from 30% to 40% manganese, 10% to 20% aluminum and the composition is described in terms of atomic percentages. In an embodiment, the average composition of the alloy comprises from 27% to 33% iron, 17% to 23% nickel, 32% to 38% manganese and 12% to 18% aluminum. In an embodiment, one of the distinct structural phases is a face-centered cubic (f.c.c.) phase. In an embodiment, the concentration of iron and of manganese in the f.c.c. phase is greater than the amount of nickel or of aluminum. In an embodiment, another of the distinct structural phases is a body-centered cubic (b.c.c.) phase, such as B2, an ordered phase. In an embodiment, the concentration of nickel and of aluminum in the B2 phase is greater than the amount of iron or of manganese. In an embodiment, at least a portion of the B2 phase is present in the form of rod-like or plate-like structures. In embodiments, the characteristic width or thickness of the B2 phase may be from 150 nm to 750 nm or from 150 nm to 500 nm. In embodiments, the average spacing between the structures of the B2 phase can be from 400 nm to 2000 nm or from 450 nm to 1000 nm. In an embodiment, at least a portion of the B2 and f.c.c. phases are present in the form of lamellar structures. Typically at least some of the B2 lamellae alternate with the f.c.c. lamellae as is characteristic of eutectic structures. The characteristic thickness of the f.c.c. phase can be from 200 nm to 2000 nm or from 300 nm to 1000 nm.
In one embodiment, an intermetallic composition formed by a eutectic transformation in at least two distinct structural phases has an average composition according to the formula:
FeaNibMncAldMe,
where (in atomic percent) a ranges from 25 to 35; b ranges from 15 to 25; c ranges from 30 to 40, d ranges from 10 to 20, e ranges from 0 to 5 and M is selected from the group consisting of Cr, Mo, C and combinations thereof.
In an aspect, the invention provides an alloy comprising chromium in addition to Fe, Ni, Mn and Al. In an embodiment, the invention provides an alloy comprising at least two distinct structural phases, wherein the average composition of the alloy comprises from 25% to 35% iron, from 15% to 25% nickel, from 30% to 40% manganese, 10% to 20% aluminum and from greater than 0 to less than or equal to 5% Cr; from 2.5% to 5% Cr, from 3% to 5% Cr, from 4% to 5% Cr, from 4% to less than 8% Cr, from greater than 5% to less than 8% Cr, from greater than 5% to 7.5% Cr, from greater than 5% to less than or equal to 7% Cr, from 5.5% to 7.5% Cr, or from 5.5% to 6.5% Cr. In an embodiment, the average composition of the alloy comprises from 27% to 33% iron, 17% to 23% nickel, 32% to 38% manganese and 12% to 18% aluminum in addition to chromium. In an embodiment, one of the distinct structural phases is a face-centered cubic (f.c.c.) phase. In an embodiment, the concentration of iron and manganese in the f.c.c. phase is greater than the amount of nickel or aluminum. In an embodiment, another of the distinct structural phases is a body-centered cubic (b.c.c.) phase, such as B2, an ordered phase. In an embodiment, the concentration of nickel and aluminum in the B2 phase is greater than the amount or iron or manganese. In different embodiments, greater than 50 at %, from 75% to 100%, from 80% to 95%, or from 80% to 90% of the chromium is located in the f.c.c. phase.
In an embodiment, at least a portion of the B2 phase in the chromium-containing alloys is present in the form of rod-like or plate-like structures. In embodiments, the characteristic width or thickness of the B2 phase may be from 150 nm to 750 nm or from 150 nm to 500 nm. In embodiments, the average spacing between the structures of the B2 phase can be from 400 nm to 2000 nm or from 450 nm to 1000 nm. In an embodiment, at least a portion of the B2 and f.c.c. phases are present in the form of lamellar structures. Typically at least some of the B2 lamellae alternate with the f.c.c. lamellae. In embodiments, the characteristic thickness of the f.c.c. phase can be from 200 nm to 2000 nm or from 300 nm to 1000 nm. It has been found that the structure of alloys at higher chromium concentrations can differ from that at lower concentrations, with this change in structure being associated with decreased ductility. For example, Fe30Ni20Mn35Al15 with 8 at % Cr produced a finer more complicated structure including cuboidal particles (see Example 4 and
In an aspect of the invention, the addition of chromium to the alloy results in improved resistance to hydrogen embrittlement in air at relatively low strain rates. In embodiments, the elongation to fracture of the chromium-containing alloy may be greater than 10% or from 10% to 25% as measured in air at room temperature at a strain rate of 5×10−4 s−1. These elongation to fracture values may be obtained for dog-bone shaped samples with a gauge length of 10 mm. The ductility of the alloy may also be indicated by the fracture mode of a specimen of the alloy. In an embodiment, the fracture surface indicates ductile tearing with elongated dimples when the alloy specimen is tested in air at room temperature at a strain rate of 5×10−6 s−1. In an embodiment, the atomic percentage of chromium is from 4% to less than 8% Cr, from greater than 5% to less than 8% Cr, from greater than 5% to less than or equal to 7% Cr, from 5.5% to 7.5% Cr, or from 5.5% to 6.5% Cr.
In an embodiment, the addition of chromium to the alloy may result in some softening of the alloy, but acceptable values of yield strength, hardness and ultimate tensile strength (UTS) of the alloy may be retained. In embodiments, the yield stress of the alloy may be from 600 to 800 MPa or 700 to 800 MPa. The UTS may be from 900 to 1000 MPa. The average hardness of the alloy, as measured by the Vickers Pyramid Number (VPN), may be from 200 to 325 kg/mm2.
In one embodiment, a method of producing an intermetallic composition includes heating a mixture of metals, to create a homogenous solution, according to the formula:
FeaNibMncAldMe,
where (in atomic percent) a ranges from 25 to 35; b ranges from 15 to 25; c ranges from 30 to 40; d ranges from 10 to 20; e ranges from greater than 0 to less than 8; and M is selected from the group consisting of Cr, Mo, C and combinations thereof; cooling the homogenous solution to obtain a homogeneous solid; reheating the solid to a eutectic transformation temperature; and holding the eutectic transformation temperature for a period of time.
Alloys that are both strong and ductile at room temperature are disclosed, along with methods for making the alloys by way of a eutectic transformation. In some embodiments, observed tensile and yield strengths of the alloys are greater than those of typical stainless steels, and the alloys show ductility comparable to high-strength ferritic stainless steels.
The terms “alloy”, “intermetallic compound” and “intermetallic composition” are used interchangeably herein. They refer to compounds containing at least two different metals.
A “eutectic alloy” is an alloy that is formed when at least two different metals, as well as any non-metals, are present in suitable concentrations and held at a eutectic transformation temperature for a suitable period of time. At the transformation temperature, at least two phases can simultaneously crystallize from a liquid solution to form lamellae of the two phases.
“Room temperature” refers to 20-25° C. as used herein.
The alloys disclosed herein may be used in the manufacture of machine, building and industrial parts. The alloys may be particularly suitable for applications requiring high-strength, wear resistant parts including but not limited to: engines, bearings, bushings, stators, washers, seals, rotors, fasteners, stamping plates, dies, valves, punches, automobile parts, aircraft parts, building materials, and drilling and mining parts. Further, the alloys can be used in any known application currently utilizing stainless steel or any high-strength, ductile alloy.
In a particular embodiment, an alloy contains iron, nickel, manganese and aluminum to which may be added chromium, molybdenum, carbon and combinations thereof. Such an alloy is represented by a macroscopic average formula:
FeaNibMncAldMe, Formula (1)
where M is an alloying addition of any element or combination of elements;
a ranges from 25 to 35;
b ranges from 15 to 25;
c ranges from 30 to 40;
d ranges from 10 to 20;
e ranges from 0 to 5; and
where a-e are expressed on an atomic percent basis.
In one aspect, M may be a metal or combination of metals. For example, M may be chromium, molybdenum, carbon and combinations thereof. In some embodiments, the portion of the alloy that is allocated to M may also range from 0.05 to 4% or in other aspects from 0.5% to 3%.
A narrower formulation that is within the general scope of Formula (1) is:
FexNi50-xMn50-yAly, Formula (2)
wherein x ranges from 25 to 35 (atomic percent basis) and y ranges from 10 to 20 (atomic percent basis).
In another aspect, the composition of FeaNibMncAldMe may be within the ranges:
a ranges from 27 to 33;
b ranges from 17 to 23;
c ranges from 32 to 38;
d ranges from 12 to 18; and
e ranges from 0 to 2.5;
where a-e are expressed on an atomic percent basis and M is an alloying addition of any element or combination of elements.
The alloy may be formed by a heat treatment process that results in a eutectic transformation leaving at least two intermetallic phases of different structure and stoichiometry. The macroscopic formulas above pertain to the overall composition, but the macroscopic composition has nanostructure or microstructure of localized phase variances in composition and ordering. The presence of two phases present as lamellae results in ductility along the planes of the lamellae. This ductility may be measured as percent elongation.
In an embodiment, at least one of the phases present in the alloy is in the form of an elongated or lamellar structure. In different embodiments, the phase may be rod-like, plate-like, or a combination thereof. As used herein, a rod-like structure need not be perfectly circular in cross-section, but has a width less than its length. In different embodiments, the elongated structure may have a characteristic length which is greater than its characteristic width, greater than twice its characteristic width, greater than five times its characteristic width, or greater than ten times its characteristic width. As used herein, a lamellar or plate-like structure has a thickness less than its width or length. A characteristic dimension of a structure may be determined from image analysis of a polished section of the alloy. For example, the characteristic width or thickness of a first phase may be determined by measuring the linear distance (within the first phase) between intersections of that phase with a second phase. The distance between structures of the first phase may be determined by measuring the linear distance (outside the first phase) between intersections of the first phase with a second phase.
The rod-like or plate-like structures need not be perfectly straight. However, in an embodiment, adjacent regions of a structure do not form an angle of 90 degrees or less with respect to each other. Structures forming an angle of 90 degrees or less are visible in
The shape and size of the phases in the alloy can influence the mechanical properties of the alloy. If the B2 phase is rod-like in form, the alloy may be viewed as comprising B2 rod-like structures in a matrix of the f.c.c. phase. If both the B2 and f.c.c. phases are plate-like in form, the alloy may be viewed as comprising alternating lamellae of B2 and f.c.c. In an embodiment, the chromium containing alloys of the invention may comprise both plate-like and rod-like B2 structures. In an embodiment, the volume fraction of the plate-like structures is greater than the rod-like structures. In an embodiment, a plurality of the rod-like structures are aligned with one another when viewed in a section approximately parallel to their longitudinal axes. In an embodiment, a plurality of the plate-like structures are aligned with one another when viewed in a section approximately transverse to the plate thickness. In different embodiments, at least 5 or at least 10 adjacent structures may be aligned with each other along at least a portion of the structure. In an embodiment, elongated structures are aligned with one another when their corresponding axes are aligned within ten degrees. In another embodiment, the length of the elongated structures may be greater than 1 micrometer, greater than 2 micrometers or greater than 5 micrometers. In embodiments, the average width or thickness structures of the B2 phase may be from 150 nm to 750 nm or from 150 nm to 500 nm. In embodiments, the average spacing between the structures of the f.c.c. phase can be from 400 nm to 2000 nm or from 450 nm to 1000 nm.
The following examples set forth preferred materials and methods for use in making the disclosed alloys. These examples teach by way of illustration, not by limitation, and should not be interpreted as unduly narrow.
A quaternary alloy of Fe30Ni20Mn35Al15 composition was prepared by well known arc melting and casting techniques. A quantity of material including 24 g Fe, 17 g Ni, 27 g Mn and 5 g Al was placed in a water-cooled copper mold and heated until molten using the arc melting technique. Ingots were flipped and melted a minimum of three times under argon to ensure mixing. Quenching was done by allowing the alloy to rapidly cool in the copper mold to a temperature of ˜30° C. in approximately 10 minutes. A eutectic transformation was carried out by holding the quenched ingots at about 1215° C. for about 30 minutes. For this composition, the eutectic transformation temperature, as shown in the differential thermal analysis curve, was between about 1210-1290° C., or between about 1212-1250° C. or between about 1214-1230° C. In some embodiments, a 5% excess of Mn may be added to the starting materials because Mn accounts for the majority of weight loss during casting, which results from brittle sharding and evaporation.
The resulting alloy had microstructure in the form of lamellae formed as two intermetallic phases. One phase was a B2 (ordered body-centered cubic, b.c.c.) phase having a composition of Fe7Ni47Mn18Al28 in terms of atomic percent. The other phase was a face-centered cubic (f.c.c.) phase having a composition of Fe50Ni7Mn37Al6 in terms of atomic percent. The widths of the body-centered cubic and face-centered cubic phases were 200 nm and 500 nm, respectively.
The alloy was characterized using analytical techniques that are well known in the art. For example, water displacement analysis was used to determine that the alloy had a density of about 7.02 g/cm3, and chemical composition was determined by energy dispersive spectroscopy (EDS). As discussed above, the overall composition, Fe30Ni20Mn35Al15, was based on microstructured phases of Fe7Ni47Mn18Al28 and Fe50Ni7Mn37Al6.
Structural data was obtained using a Siemens D5000 X-ray Diffractometer with a Kevex silicon detector in the range of 20-110° 2θ, using an instrument that was calibrated against an alumina standard purchased from the National Institute of Standards (NIST).
Room temperature hardness of the two phase alloy, Fe30Ni20Mn35Al15, was determined by taking the average of five measurements from a Leitz Microhardness Indentor with a 200 g load. The average Vicker's hardness was 310 kg/mm2.
Differential thermal analysis (DTA) was performed on a Perkin Elmer Pyris Diamond TG/DTA. A typical DTA curve is shown in
Transmission electron microscopy (TEM), performed on either a JEOL 2000FX or a Philips CM 200, indicated that, after formation, the eutectic microstructure was stable up to at least 760° C. Optical microscopy confirmed that no distinct differences were observed between samples that had been annealed for 30 minutes at temperatures between 327-727° C. and then subsequently quenched.
Yield strength of the eutectic alloy was determined using a MTS 810 mechanical testing system. The two phase alloy was subjected to mechanical testing at temperatures shown in
The stress versus strain curve of
As shown in
In contrast,
Various alloys are cast with a composition:
FexNi50-xMn50-yAly, Formula (2)
where x ranges from 25 to 35 atomic percent plus or minus 5%, and y ranges from 10 to 20 atomic percent plus or minus 5%.
The alloys are cast using the aforementioned arc melting technique and heated to a eutectic transformation temperature range of between about 1210-1290° C., or between about 1212-1250° C. or between about 1214-1230° C. The alloys are expected to be strong and ductile with a range of mechanical properties that can be manipulated by composition variations within the disclosed range.
Table 1 gives the percent elongation to failure, yield stress and ultimate tensile strength measured for alloys with less than 15 at % aluminum. (strain rate 5×10−4 s1). Table 2 gives hardness measurements for alloys with less than 15 at % aluminum.
Characterization of a Phase Diagram Near a Eutectic Transformation
A portion of a phase diagram near a eutectic transformation may be constructed by varying percentages of Fe, Ni, Mn, Al and M as described in the context of Formula (1), except the subscripts a, b, c, d, and e, may be any value. The constituents are processed as described in Examples 1 and 2 to ascertain the presence or absence of eutectic transformation products. The preferred metals include combinations of Fe, Ni, Mn, and Al, in which case the ranges for x and y shown in Formula (2) may be any value. When adjusting the respective subscripts a, b, c, d, e, x and/or y, it is suggested to increase or decrease the individual ranges or combinations of ranges in steps of five percent from the values shown in Formulas (1) and (2), at least until the resulting alloy does not show evidence of a eutectic transformation. For alloys that contain four or five constituents, it is routine in the art that several hundred castings are needed to fully characterize the phase diagram around a eutectic transformation.
Fe30Ni20Mn35Al15 with different additions of Cr (0-8 at. %) were prepared by arc melting a mixture of elemental Cr, Fe, Ni, Mn and Al with purities >99.8% in a water-cooled copper crucible under an argon atmosphere. The ingot was melted and flipped three times using a probe inside the chamber to ensure homogeneity, followed by cooling rapidly to room temperature.
X-ray diffraction (XRD) was performed on the alloys using a Rigaku D/Max 2000 diffractometer with Cu Kα radiation operated at 40 kV and 300 mA. Measurements were performed by step scanning 2θ from 20° to 120° with a 0.02° step size. A count time of 1 s per step was used. Specimens, cut from the ingot, were polished using increasingly fine grade of silicon carbide papers followed by 0.3 μm alumina powder to a mirror finish. The surface was etched using 4% nitric acid for ˜5 s followed by rinsing in water. Specimens were examined in a FEI XL-30 field emission gun scanning electron microscope (SEM) operated at 15 kV. Discs of 3 mm in diameter and 100 μm thick, from either before or after ˜5% strain under compression, were electropolished using 30% nitric acid in methanol at 253 K in a Struers Tenupol 5, and washed alternatively in ethanol and methanol for three cycles followed by a final rinse in methanol. The typical voltage and current were 11 V and 100 mA, respectively. The resulting thin foils were examined using an FEI Tecnai F20 field emission gun transmission electron microscope (TEM) equipped with energy dispersive X-ray spectrometry (EDS), operated at 200 kV.
Flat tensile specimens of dog bone geometry with a gauge length of 10 mm and thickness of ˜1.27 mm were produced for tensile tests. The specimens were polished using silicon carbide papers and finished with 0.3 μm alumina powder to reduce surface defects. Tensile tests were performed at different strain rates (5×10−6 s−1−5×10−1 s−1) at room temperature and at a strain rate of 5×10−4 s−1 at temperatures from 300 K to 1000 K.
Secondary electron (SE) images of Fe30Ni20Mn35Al15 with different Cr additions are shown in
XRD patterns of Fe30Ni20Mn35Al15 with 6 at. % Cr and 8 at. % Cr, see
The effects of Cr on the room temperature strength and
ductility of Fe30Ni20Mn35Al15 are illustrated in
BF TEM analysis of Fe30Ni20Mn35Al15 with 6 at. % Cr after ˜5% strain showed dislocations were generated within f.c.c. lamellae and piled up at interphase interfaces. However, no dislocations were observed in the B2 phase. It is evident that the deformation mainly occurred in f.c.c phase, while the B2 phase acted as obstacles to moving dislocations. Room temperature tensile tests were also performed as a function of strain rate for Fe30Ni20Mn35Al15 with 6 at. % Cr, see
Fracture surfaces from the 6 at. % Cr-modified Fe30Ni20Mn35Al15
tensile tested at a strain rate of 5×10−6 s−1 at room temperature showed ductile tearing with elongated dimples (See
The fracture surfaces of the 6 at. % Cr-modified Fe30Ni20Mn35Al15 tensile tested at elevated temperature at strain rate of 5×10−4 s−1 were also examined. Examination of typical fracture surfaces at 500 K and 800 K showed ductile, dimple-type rupture mode was found at both temperatures, but the size and depth of the dimples was greater at 800 K compared to those at 500 K, which is to be expected since the tests at 800 K showed maxima in both elongation and reduction in area.
Chromium additions up to 6 at. % have three intriguing effects on the room temperature mechanical behavior of Fe30Ni20Mn35Al15, while producing little change in the microstructure.
First, Cr decreases the yield strength. Since plastic deformation at room temperature is accommodated solely by plastic deformation of the f.c.c. phase and as shown by EDS, most of the Cr partitions into the f.c.c. phase (FIGS. 14D-E)., it is evident that the effect of Cr is to soften this phase.
The second intriguing effect is the increase in ductility with increasing Cr (up to 6 at. %). Previous work showed that dislocation pile-ups in the f.c.c. phase at the f.c.c./B2 interface produced cracking in the B2 phase (Liao and Baker, 2011, Mater. Sci. Eng.: A 528, 3998-4008), which ultimately led to failure. Without wishing to be bound by any particular belief, a reduction in the length of dislocation pile-ups at a given stress due to Cr additions may delay this cracking until larger strains are reached. The third effect is that the Cr addition (particularly for 6 at. % Cr) suppresses the environmental embrittlement problem observed in unalloyed Fe30Ni20Mn35Al15 when tested at slow strain rates at room temperature (Liao et al., 2011). Without wishing to be found by any particular believe, there may be several explanations for the effect of chromium on the environmental embrittlement in Fe30Ni20Mn35Al15. First, as noted above, the Cr addition may increase the stacking fault energy and lead to wavy slip, as observed by McKamey et al. in Fe3Al (J. Mater. Res. 1989, 4, 1156-1163; Scr. Metall, 1988, 22, 1670-1681). This would reduce dislocation pile-ups, and, hence, stress concentrations. It may also reduce the ability of gliding dislocations to transport atomic hydrogen, produced by the reaction between water vapor and Al, into the material. However, it is worth noting that Schramm and Reed (1975, Metall. Trans. A, 6, 1345-1351) suggested that Cr increased the stacking fault energy of austenitic stainless steels.
Second, the Cr addition may result in a change in oxide chemistry and properties, or a change in the kinetics of oxide formation (Mckamey and Liu, 1990, Scr. Metall. Mater. 24, 2119-2122). It is likely that mixed oxides are present on the surface when Cr is present. Both of these mechanisms will minimize the environmental embrittlement by reducing the water-vapor reaction. Finally, the possible mechanisms responsible for the beneficial effect of Cr addition on the embrittlement of Ni3(Si, Ti) alloys suggested by Ma et al. (1995, Scr. Metall. Mater. 32, 1025-1029), i.e. blocking site occupation and/or the diffusion of hydrogen along grain boundaries; affecting both the water decomposition process and the subsequent hydrogen absorption processes; and enhancing the cohesive strength of the grain boundaries, may also be important in the alloy tested here.
The temperature dependence of the yield strength of the 6 at. % Cr modified alloy is essentially the same as the Cr-free alloy (Liao and Baker, 2011, J. Mater. Sci. 246, 2009-2017). Differential thermal analysis showed that the melting point of 6 at. % Cr modified Fe30Ni20Mn35Al15 was ˜1500 K. The yield strength of Cr modified Fe30Ni20Mn35Al15 is only weakly dependent on temperature up to 600 K, which is ˜0.45 Tm. The rapid decrease in yield stress with increasing temperature above 600 K has been observed in many B2 compounds above 0.45 Tm and is associated with the onset of diffusive processes (Baker, 1995, Mat. Sci. Eng. A-Struc. Mater. Prop. Microstruct. Process. 192, 1-13). The B2 phase in Fe30Ni20Mn35Al15 experiences a brittle-to-ductile transition with increasing temperature, as in many B2 compounds (Baker, 1995). At lower temperatures, the B2 phase is hard and acts as an obstacle to dislocation motion, eventually fracturing due to stress concentrations from dislocation pile-ups in the f.c.c. phase (Liao and Baker, 2011, Mater. Sci. Eng. A 528, 3998-4008). In contrast, at higher temperatures the B2 lamellae undergoes plastic deformation, and the overall yield stress the elongation to failure and reduction in area at 800 K followed by a decrease to their lowest values at 900 K and 1000 K. Dimple-type ductile fracture was observed at all temperatures. In contrast, B2 compounds typically show increasing ductility with increasing temperature (Baker 1995) For example, for boron-doped (500 ppm) B2 Fe-45Al the elongation to failure increased from —3% at room temperature with increasing temperature up to 40% at 800 K, but this was followed by a sharp drop to 20% above that temperature when necking occurred prior to failure (Klein and Baker, 1994, Scr. Metall. Mater., 30, 1413-1417). However, in contrast to the present work, the reduction in area increased continuously with increasing temperature reaching a value of 90% at 1000 K (Klein and Baker, 1994). The reason for the decrease in elongation at higher temperatures is because the FeAl no longer showed any work-hardening. Thus, once a neck formed during deformation, there was no capacity for work hardening to offset the reduction in area at the neck.
The lack of a significant increase in ductility up to 700 K in Fe30Ni2Mn35Al15 may reflect the fact that the B2 phase does not undergo plastic deformation up to that temperature as indicated by the slight change in yield strength with increasing temperature up to that point. The peak at 800 K is presumably related to the fact that the B2 phase now undergoes plastic deformation as reflected in a substantial decrease in yield strength at that temperature. Similar to FeAl, the reduction in elongation observed at 900 K and 1000 K is due to the lack of work-hardening at these temperatures, which produces unstable necking. However, it is somewhat surprising that, unlike FeAl (Klein and Baker, 1994), the reduction in area also decreases at 900 K and 1000 K. This probably reflects the extreme cavitation in the neck.
Additional details may be found in Meng et al., 2013, Mater. Sci. Engr. A, 586, 45-52, hereby incorporated by reference.
It is understood for purposes of this disclosure, that various changes and modifications may be made to the disclosed embodiments that are well within the scope of the present compositions and methods. Numerous changes may be made which will readily suggest themselves to those skilled in the art and which are encompassed in the spirit of the compositions and methods disclosed herein and as defined in the appended claims.
All references throughout this application, for example patent documents including issued or granted patents or equivalents; patent application publications; and non-patent literature documents or other source material; are hereby incorporated by reference herein in their entireties, as though individually incorporated by reference, to the extent each reference is at least partially not inconsistent with the disclosure in this application (for example, a reference that is partially inconsistent is incorporated by reference except for the partially inconsistent portion of the reference).
All patents and publications mentioned in the specification are indicative of the levels of skill of those skilled in the art to which the invention pertains. References cited herein are incorporated by reference herein in their entirety to indicate the state of the art, in some cases as of their filing date, and it is intended that this information can be employed herein, if needed, to exclude (for example, to disclaim) specific embodiments that are in the prior art. For example, when a compound is claimed, it should be understood that compounds known in the prior art, including certain compounds disclosed in the references disclosed herein (particularly in referenced patent documents), are not intended to be included in the claim
As used herein, “comprising” is synonymous with “including,” “containing,” or “characterized by,” and is inclusive or open-ended and does not exclude additional, unrecited elements or method steps. As used herein, “consisting of” excludes any element, step, or ingredient not specified in the claim element. As used herein, “consisting essentially of” does not exclude materials or steps that do not materially affect the basic and novel characteristics of the claim. Any recitation herein of the term “comprising”, particularly in a description of components of a composition or in a description of elements of a device, is understood to encompass those compositions and methods consisting essentially of and consisting of the recited components or elements. The invention illustratively described herein suitably may be practiced in the absence of any element or elements, limitation or limitations which is not specifically disclosed herein.
This application is a continuation in part application of U.S. application Ser. No. 12/867,712, international filing date Feb. 13, 2009 which is hereby incorporated by reference and which is the national stage of International Application No. PCT/US09/34123, filed Feb. 13, 2009, which claims the benefit of priority to U.S. Provisional Patent Application Ser. No. 61/028,809, filed Feb. 14, 2008, which is incorporated by reference herein.
The United States Government has rights in this invention under Contract No. NSF-DMR-0505774 and NSF DMR-0905229 between the National Science Foundation (NSF) and Dartmouth College and also under Contract DE-FG02-07ER46392, between the U.S. Department of Energy and Dartmouth College.
| Number | Date | Country | |
|---|---|---|---|
| 61028809 | Feb 2008 | US |
| Number | Date | Country | |
|---|---|---|---|
| Parent | 12867712 | Nov 2010 | US |
| Child | 14222230 | US |
