Patent Application
20240368737
Embodiments of the invention relate to low coefficient of thermal expansion (CTE) alloys, and in some embodiments, low CTE high strength (LCHS) alloys.
Applications for low CTE alloys include high temperature combustor-turbine section seals, containment rings, duct segments, spacer, casing, rocket nozzles, and pumps. For example, by using low CTE alloys, the gap between turbine components does not open up as temperature changes during operation. Thus, low CTE alloys can help an engine achieve better efficiency by preventing hot gas leaking through the blade or vane tip clearance (between rotating and static components during heat-up and cool-down cycles) with passive sealing. This incorporation of low CTE alloys in engines can be used in combination with active clearance control technology (ACC), with low CTE alloys reducing the amount of re-directed compressed air needed to operate ACC, thereby increasing efficiency. With low CTE alloys, less strain is caused by expansion and contraction due to temperature change during engine operation.
Currently available low CTE alloys for elevated temperature applications are limited by their strength and thermal stability with a typical service temperature ceiling at about 1200° F.-1250° F. Existing low CTE alloys, however, do not have sufficient strength or long-term microstructure stability for exposure to a service temperature of around 1400° F.
There is an industrywide need for a low CTE alloy that can be used at higher temperatures, e.g., ≥1400° F.
Embodiments of the invention include low CTE alloys, and in some embodiments, low CTE high strength (LCHS) alloys.
Advantageously, embodiments of the alloys disclosed herein have a similar or lower CTE compared to existing alloys from room temperature to 1400° F. Further, the alloys disclosed herein have a higher strength, e.g., yield strength (“YS”) and/or ultimate tensile strength (“UTS”) at high temperatures such as 1400° F. Finally, alloys disclosed herein have good thermal stabilities at high temperatures such as 1400° F., as well as better oxidation or corrosion resistance than existing alloys.
In an aspect, embodiments of the invention relate to a low coefficient of thermal expansion high strength alloy. The alloy may include or consist essentially of: chromium 7 wt. % to 10 wt. %, molybdenum 20 wt. % to 25 wt. %, tungsten 4 wt. % to 7 wt. %, aluminum 0.5 wt. % to 2 wt. %, titanium 0.5 wt. % to 2 wt. %, boron 0.005 wt. % to 0.05 wt. %, niobium ≤3.9 wt. %, tantalum ≤3.9 wt. %, vanadium 0.1 wt. % to 4 wt. %, niobium, tantalum, and vanadium, in combination 0.1 wt. % to 4 wt. %, silicon <0.5 wt. %, zirconium <0.5 wt. %, hafnium <0.5 wt. %, yttrium <0.5 wt. %, copper <0.1 wt. %, manganese <0.1 wt. %, phosphorus <0.1 wt. %, sulfur <0.1 wt. %, iron <5 wt. %, cobalt ≤15 wt. %, balance nickel, cobalt and nickel, in combination 50 wt. % to 70 wt. %, and aluminum and titanium, in combination ≥1.4 wt. %.
One or more of the following features may be included. The alloy may include: cobalt 5 wt. % to 15 wt. %. The alloy may include vanadium and niobium, in combination 0.5 wt. % to 4 wt. %. The alloy may include vanadium 0.5 wt. % to 4 wt. %. The alloy may include vanadium and titanium, in combination 0.8 wt. % to 3.5 wt. %. The alloy may include tungsten 5.5 wt. % to 7 wt. %. The alloy may include molybdenum 21 wt. % to 24 wt. %. The alloy may include molybdenum and (tungsten)/2, in combination 24 wt. % to 27 wt. %. A ratio of molybdenum wt. % to tungsten wt. % may range from 3.6 to 4.2. The alloy may include titanium and aluminum, in combination 1.4 wt. % to 4 wt. %. A ratio of titanium wt. % to aluminum wt. % may be ≥0.4. The alloy may include carbon 0.005 wt. % to 0.05 wt. %.
A coefficient of thermal expansion of the alloy may be less than 8×10−6 inch/(inch ° F.) between room temperature to 1400° F. A yield strength of the alloy may be higher than 85 ksi at 1400° F. An ultimate tensile strength of the alloy may be higher than 120 ksi at 1400° F.
The alloy may include more than 4 vol. % gamma-prime and gamma double prime phases in combination. The alloy may include less than 5 vol % eta and delta phases in combination. The LCHS alloy may include: 4 vol. %-12 vol. % mu and P phases in combination. The LCHS alloy may include at least 25 vol. % Ni2M phase, wherein M is selected from the group consisting of Cr, Mo, W, and V.
In another aspect, embodiments of the invention relate to a method of manufacturing a low coefficient of thermal expansion high strength alloy. The method includes the steps of: melting a plurality of elements including the composition of: chromium 7 wt. % to 10 wt. %, molybdenum 20 wt. % to 25 wt. %, tungsten 4 wt. % to 7 wt. %, aluminum 0.5 wt. % to 2 wt. %, titanium 0.5 wt. % to 2 wt. %, boron 0.005 wt. % to 0.05 wt. %, niobium ≤3.9 wt. %, tantalum ≤3.9 wt. %, vanadium 0.1 wt. % to 4 wt. %, niobium, tantalum, and vanadium, in combination 0.1 wt. % to 4 wt. %, silicon <0.5 wt. %, zirconium <0.5 wt. %, hafnium <0.5 wt. %, yttrium <0.5 wt. %, copper <0.1 wt. %, manganese <0.1 wt. %, phosphorus <0.1 wt. %, sulfur <0.1 wt. %, iron <5 wt. %, cobalt ≤15 wt. %, balance nickel, cobalt and nickel, in combination 50 wt. % to 70 wt. %, and aluminum and titanium, in combination ≥1.4 wt. %; homogenizing the plurality of elements; hot working the plurality of elements at a first temperature; and solution and age the plurality of elements at a second temperature to precipitate out gamma prime (γ′) and/or gamma double prime (γ″) phases and Ni2M phase. Mu phase and/or P phase may be present at the first temperature. The second temperature is higher than or equal to an intended service temperature.
The intended service temperature may be at least 1400° F.
Existing low CTE alloys for elevated temperature applications are limited by their strength and thermal stability with a service temperature that reaches an upper limit at about 1200° F.-1250° F.
As jet engine and gas turbine technologies develop, the turbine section runs at higher temperatures with increasing compression ratio, but development is limited due to limitations of existing low CTE alloys. For example, existing low CTE alloys, such as Carpenter Thermo-Span®, HAYNES 242, HAYNES 244 and INCONEL® alloy IN783 (Haynes 242 and Hayes 244 are trademarks of Haynes International, Inc. Inconel is a registered trademark of Special Metals), have a maximum application temperature of around 1250° F., and a typical service temperature at 1200° F. If exposed to a service temperature around 1400° F., these alloys may not have sufficient strength or long-term microstructure stability. Some existing low CTE alloys use the P/mu phase to control grain size while using the oP6 (Ni2M) phase for achieving bulk strength. For example, the main strengthening phase in HAYNES 242 and HAYNES 244 is Ni2M, which becomes less stable at temperatures above 1300° F. Thus, new high temperature materials having desirable performance are needed.
An exemplary composition range, in accordance with embodiments of the invention, is provided in Table 1.
The function and composition ranges of various elements in embodiments of the invention are described below, along with the criteria for selecting the levels of these elements.
Chromium (Cr): Preferably, Cr levels are in the range of 7 wt. % to 10 wt. %.
Cr is critical for oxidation and corrosion resistance. Traditional low CTE alloys such as IN909 cannot have high Cr because Cr depresses the Curie temperature (i.e., the temperature above which certain materials lose their permanent magnetic properties). In these alloys, the Curie temperature is directly related to the Invar effect (magnetic alloys having an anomalously small thermal expansion), which is used to reduce the coefficient of thermal expansion. As the LCHS alloy does not rely on the magnetostriction from the Invar effect to achieve its low thermal expansion, more chromium can be added to achieve much better oxidation and corrosion resistance at high temperatures. Cr is also an important element for forming the Ni2M strength-enhancing phase. Thus a 7 wt. % minimum is required to achieve good mechanical strength and corrosion resistance. Too much Cr reduces the stability of the microstructure and leads to formation of undesirable low ductility phases such as the sigma phase. Thus, 10 wt. % is the maximum preferred amount for Cr.
Molybdenum (Mo): Preferably, Mo levels are in the range of 20 wt. % to 25 wt. %.
The LCHS alloy derives its low thermal expansion from the low elemental thermal expansion of its constitute elements like Mo and W. Thus, a significant amount of Mo is needed for low CTE (minimum 20 wt. %). Mo also helps form the strengthening Ni2M phase. However, too much Mo reduces the microstructure stability and leads to excessive formation of undesirable phases, hence the 25 wt. % upper limit.
Tungsten (W): Preferably, W levels are in the range of 4 wt. % to 7 wt. %.
Similarly to Mo, W has low elemental thermal expansion. It also helps form Ni2M phase. W also provides solid solution strengthening and increases the creep resistance at high temperature. But too much tungsten increases the density of the alloy, thus 7 wt. % is the preferred maximum.
Aluminum (Al): Preferably, Al levels are in the range of 0.5 wt. % to 2 wt. %.
Al helps form the gamma prime (phase, which contributes to elevated temperature strength. Al is also beneficial for high temperature oxidation resistance. Too much Al reduces the hot workability of the alloy and leads to cracking during processing.
Titanium (Ti): Preferably, Ti levels are in the range of 0.5 wt. % to 2 wt. %.
Ti also forms the gamma prime phase. With more Ti in the gamma prime phase, the anti-phase boundary (APB) energy increases and makes it more difficult to shear gamma prime particles. This increases the yield strength of the alloy. Too much titanium leads to formation of the eta phase, which is not as good a strengthening phase as the coherent gamma prime phase. Also, the ratio of Ti and Al is critical to obtain stable and strong gamma prime. Just like Al, too much Ti reduces the hot workability of the alloy.
Boron (B): Preferably, B levels are in the range of 0.005 wt. % to 0.05 wt. %.
A small amount of boron (0.005 wt % to 0.05 wt. %) improves the grain boundary strength at elevated temperatures. It also forms molybdenum borides along grain boundary which helps refine grain boundary carbides by reducing the size of carbides and improves elevated temperature ductility. Too much boron, such as above 0.05 wt. %, leads to the formation of an undesirable continuous low melting temperature boron rich film along the grain boundary.
Niobium (Nb): Preferably, Nb levels are in the range of ≤3.9 wt. %.
Nb helps form gamma prime and gamma double prime phases, and sometimes a co-precipitation of gamma prime and gamma double prime phases with a compact structure. Too much Nb promotes the formation of the delta phase, which is not as good a strengthening phase as the fine dispersed gamma prime or gamma double prime phases. Excessive levels of Nb are also associated with macro segregation defects that form during solidification that negatively impact performance.
Tantalum (Ta): Preferably, Ta levels are in the range of ≤3.9 wt. %.
Ta behaves similarly as Nb, which goes into the gamma prime or gamma double prime phase. Ta also improves creep resistance and corrosion resistance. However, Ta has high cost and increases the density of the alloy, and too much Ta also promotes delta phase formation. Thus the upper limit is 3.9 wt. %.
Vanadium (V): V may facilitate the formation of the gamma prime phase, increase the APB energy in the gamma prime phase, and/or improve the tensile strength of the alloy. Further, V may help increase the diffusivity of Nb, which may promote the formation of the gamma double prime phase. Further, V increases the stability of the Ni2M phase. Preferably, V levels are in the range of 0.1 wt. % to 4 wt. %. Too much V can reduce the hot workability of the alloy.
Nb, Ta, and V, in combination 0.1 wt. % to 4 wt. %.
Nb, Ta, and V all promote the formation of the gamma prime and gamma double prime phases. These elements can change APB energy as well as lattice misfit. Too much leads to the formation of the delta phase.
Silicon (Si): Preferably, Si levels are in the range of ≤0.5 wt. %.
Silicon can be beneficial for elevated temperature oxidation resistance, but too much silicon can lead to excessive formation of undesirable phases like the brittle Laves phase.
Zirconium (Zr): Preferably, Zr levels are in the range of <0.5 wt. %.
A small amount of Zr can be beneficial for strengthening grain boundaries. Excessive Zr forms undesirable phases with low melting temperatures.
Hafnium (Hf): Preferably, Hf levels are in the range of <0.5 wt. %.
A small amount of Hf can improve oxidation resistance as well as grain boundary strength.
Yttrium (Y): Preferably, Y levels are in the range of <0.5 wt. %.
A small amount of Y can improve the oxide by preventing oxide layer spallation. Too much Y leads to processing issues, such as causing incipient melting during homogenization.
Copper (Cu): Preferably, Cu levels are in the range of <0.1 wt. %.
Cu is an undesirable impurity.
Manganese (Mn): Preferably, Mn levels are in the range of <0.1 wt. %.
Mn is an unintentional impurity.
Phosphorus (P): Preferably, P levels are in the range of <0.1 wt. %.
P is an unintentional impurity.
Sulfur (S): Preferably, S levels are in the range of <0.1 wt. %.
Sulfur is an undesirable impurity; hence the lower the concentration, the better.
Iron (Fe): Preferably, Fe levels are in the range of <5 wt. %.
A small amount of Fe can reduce the cost of the alloy without having a detrimental effect on performance.
Cobalt (Co): Co is a solid solution strengthening element that can increase high temperature strength and creep resistance of the alloys described herein. The stacking fault energy in FCC matrix is lowered with increasing cobalt, which reduces creep rate. Co can also reduce the solubility of gamma prime phase forming elements (such as Al and Ti) in the FCC matrix, thereby increasing the equilibrium gamma prime volume fraction. Accordingly, higher strength can be achieved with the same amount of Al and Ti when Co is added, which may avoid the challenge of lowered hot workability associated with higher Al and Ti content. Co can raise the solidus of the alloy, which improves hot workability by extending the hot working temperature range. Adding Co can also improve oxidation and/or hot corrosion resistance. In alloys with high Mo content, Co can increase the solvus of the P phase and/or mu phase, and thus can be useful for grain size control during hot working (e.g., hot forging/hot rolling). Too much Co suppresses formation of the Ni2M phase. Preferably, Co levels are in the range of ≤15 wt. %.
Nickel (Ni): Ni is the main alloy element, and constitutes the balance of the composition, with Co and Ni in combination ranging from 50 wt. % to 70 wt. % Aluminum and titanium, in combination ≥1.4 wt. %, e.g., 1.4 wt. % to 4 wt. %.
Presence of Al and Ti leads to the formation of the strengthening gamma prime phase in the Ni-based matrix. The gamma prime phase may precipitate on grain boundaries (GB), which can be beneficial for breaking up the P or mu phase, thereby avoiding the formation of a continuous band around the GB (which may cause the alloy to be brittle). The resulting microstructure can enhance both the strength and ductility of the alloy. Further, addition of Al and/or Ti may increase the stability of the Ni2M phase.
A minimum of 1.4 wt. % of Al+Ti is needed to achieve the targeted equilibrium gamma prime volume fraction during service. A 4 wt. % maximum ensures that the hot workability of the alloy is still good.
Al/Ti: A ratio of A/Ti in weight percentage is preferably higher than 0.4, e.g., 0.5-5, to ensure the stability of the gamma prime phase, which is thereby the main strengthening phase, rather than other phases such as the eta phase.
Mo+W/2: Preferably, Mo+W/2 levels are in the range of 24-27 wt. %, and the ratio of Mo/W (wt. %) is between 3.6-4.2. Mo and W each have low CTE individually. A sufficient amount of W and Mo is needed for lowering CTE, but too much W and/or Mo may not only increase the density but also lead to formation of undesirable phases. The preferred ratio helps ensure that the balance between Mo and W increases the stability of the Ni2M phase at elevated temperatures (e.g., around 1400° F.).
The relatively high Cr, Mo content as well as addition of Al can help the alloys disclosed herein achieve better oxidation and/or corrosion resistance than existing low CTE alloys such as IN909 and Thermo-Span®. This may be demonstrated with cyclic oxidation tests as well as sulfidation tests.
Advantageously, the alloys disclosed herein can have a similar or lower CTE, from room temperature to 1400° F., compared to existing alloys such as HAYNES 244 or IN783. Furthermore, the alloys disclosed herein can have a higher strength (e.g., yield strength and/or ultimate tensile strength) at high temperatures such as 1400° F. than HAYNES 244, IN783 or Thermo-Span®. Finally, the alloys disclosed herein can have good microstructural stabilities at high temperatures such as 1400° F., as well as better oxidation or corrosion resistance than existing alloys such as IN903, IN907, IN909 and Thermo-Span®.
For example, alloys disclosed herein can have a CTE lower than 8 inch/inch/° F. (i.e., less than 8×10−6 inch/inch/° F. from room temperature (RT) to 1400° F. Under some conditions (e.g., a properly heat treated state), alloys in accordance with embodiments of the invention have a YS higher than 85 ksi, and a UTS higher than 110 ksi at 1400° F.
Alloy LCHS-F was aged and exposed to an elevated temperature for an extended period, and tested for micro-hardness. Table 2 provides the microhardness results (Rockwell hardness, scale C) for alloy LCHS-F.
An aging process can be used on solution heat treated alloys to increase their strength and hardness, while retaining their ductility. A suitable two step aging process may include a first step of heating the alloy to 1300-1500° F. for 8-32 hours followed by a slow cool (about 100° F./hour cooling rate) to the second heating step at a temperature 1100-1300° F., and dwelling at the 2nd temperature for 8-48 hours with air cool or a faster quench.
As shown in Table 2, aged alloy LCHS-F samples exhibit micro-hardness with an average of about 51 HRC, while alloy LCHS-F samples that have been exposed to an elevated temperature for an extended period exhibit micro-hardness with an average of 47.2 HRC. Thus, after long term thermal exposure, minimal reduction of hardness was observed due to the coarsening of precipitates. The minimal change in hardness shows that the strength is still relatively high. For instance, AMS5709J specification for Waspaloy only requires hardness of 32-42 HRC after aging. Furthermore, no cracking was associated with the hardness indents, showing that the alloy's ductility did not degrade.
Alloys in accordance with embodiments of the invention have a higher oxidation resistance than currently available Ni—Fe—Co low CTE alloys (e.g., INCOLOY® Alloys 907 and 909). Alloy LCHS-F exhibits less specific weight change due to oxidation at a temperature between 1400° F. and 1500° F. and for a time up to 1000 hours with isothermal or cyclic exposure.
Table 3 illustrates the phase transformation temperatures obtained from DSC (differential scanning calorimetry) analysis for the alloy LCHS-F. An exemplary DSC heat ramp is provided in
DSC may be used for determination of the solvus temperatures of strengthening phases using the techniques, e.g., of Verma et al. as illustrated and described in
Embodiments of the invention provide elevated temperature phase stability by raising the solvus temperature to 771° C. (1420° F.) for the Pt2Mo phase and 1049° C. (1921° F.) for the gamma prime strengthening phase. The higher temperature stability of these phases contributes to the unique combination of high temperature strengthening and low thermal coefficient of expansion of the compositions disclosed herein.
Field emission SEM and EDS were used to characterize the aged alloy LCHS-F, and the results are discussed in the Examples section.
In some embodiments (e.g., a properly heat treated state), the alloys disclosed herein include mu or P phase on the grain boundaries, which control grain size. The mu or P phase's solvus can be higher than 2000° F.
The alloys disclosed herein may also contain other phases such as eta or delta phase for grain size control, as well as other strengthening mechanisms, such as Laves phase, gamma double prime, and co-precipitation of γ′/γ″ or Ni2M/γ″.
In some embodiments (e.g., a properly heat treated state), the alloys disclosed herein may include dispersed Ni2M and gamma prime phases inside the grains for high strength. The solvus of gamma prime can be higher than 1900° F. and the equilibrium γ′ fraction at 1400° F. can be higher than 4%. The solvus of the Ni2M phase can be higher than 1400° F. Thus, the alloys disclosed herein can be stable for extended service at 1400° F. or higher temperatures.
In terms of microstructure, in the equilibrium state corresponding to long term exposure at 1400-1500° F. (last aging temperature or service temperature), the volume fraction of the equilibrium gamma prime phase can be higher than 4%, and lower than 12%. In some embodiments, the mu and/or P phase fraction can be higher than 4% and lower than 12%. A sufficient amount of gamma prime is needed for achieving good mechanical strength, but too much gamma prime makes the alloy crack prone and difficult to process, especially for the LCHS alloys that already contain a significant amount of Ni2M phase. Too much gamma prime (>12%) in combination with Ni2M (which can be as high as 30%) makes it very difficult if not impossible to process the material with a conventional wrought process. Similarly, a sufficient amount of mu/P phases can help control the grain size, while excess mu/P phases reduces the ductility of the alloy.
In terms of other phases, the volume fraction of equilibrium gamma prime and gamma double prime phases in combination can be higher than 4%. The volume fraction of equilibrium eta and delta phases in combination can be less than 5%. The volume fraction of equilibrium Ni2M phase can be equal to or more than 25%, with M being Cr, Mo, W, or V.
The processing or manufacturing of the alloys disclosed herein may include the following steps: 1) homogenization after melting, 2) hot working at a temperature at which the mu phase or P phase is still present (e.g., below the solvus of the mu/P phases) to control the grain size, 3) solution and age to precipitate out the gamma prime phase and the Ni2M phase. The aging temperature may be higher or equal to an intended service temperature. The alloys can also be further strengthened by cold work prior to final aging. Moreover, the alloys can be produced in powder form with gas atomization and used for powder metallurgy or additive manufacturing.
Table 4 illustrates the processing parameters in which the mu or P phase controls the grain size during forging and solution to maximize the strength derived from Hall-Petch relationship. The aging process is configured to precipitate out Ni2M and gamma prime phases for strengthening the alloys as described herein. The exposure heat treatment is configured to obtain alloys with long-term thermal stability.
Examples of articles that may be made of the alloys described herein and processed in accordance with embodiments of the invention can be but not limited to high temperature combustor-turbine section seals, containment rings, duct segments, spacer, casing, rocket nozzles, and pumps with reduced thermal stress and improved tolerance with temperature variation. The alloys may also be used to make high temperature fasteners, preventing relaxation due to thermal expansion.
Tables 5A and 5B list the compositions of LCHS alloys A-F that include different levels of Al and Ti for introducing the strengthening γ′ phase. More specifically, Table 5A lists the composition ranges in weight percentage, atomic ratio of Al to Ti, and targeted γ′ phase in volume fraction (vf). Table 5B lists the composition ranges in atomic percentage for the LCHS alloys A-F as shown in Table 5A, and atomic percentage of Al and Ti in combination. As shown in Table 5B, the alloys disclosed herein may include Al+Ti>2.5 at % and/or Al+Ti>1.4 wt. %. Thermodynamic modeling was performed using ThermoCalc (trademark of ThermoCalc Software), JMatPro(trademark of Sente Software Ltd) and MatCalc(trademark of MatClac Engineering GmbH) to down-select composition for trials. The properties of alloy LCHS-F are described above.
Table 6 provides compositions in weight and atomic percentages of an actual melting target for a button heat for alloy LCHS-F, whose composition is given in Tables 5A and 5B.
Table 7 lists the strength (including yield strength (YS) and ultimate tensile strength (UTS), elongation percentage, and reduction in area percentage (RA %)) for alloy LCHS-F at elevated temperatures of 1200° F. and 1400° F.
The alloy has very high strength at 1200° F., which is the typical service temperature for existing low CTE alloys. LCHS-F still retained good strength at 1400° F., which exceeds the application temperature for traditional low CTE alloy, because of the thermal stability of the strengthening phases. Also, LCHS-F is stronger than traditional superalloys such as Waspaloy, etc, while also having an ultimate strengths of 115-130 ksi at 1400° F. and significantly higher CTE.
Analysis of Grain Structures in Alloys in Accordance with an Embodiment of the Invention
Low CTE alloy LCHS-F in accordance with an embodiment of the invention was analyzed with SEM and EDS. The LSCHS-F samples included an aged sample and a long-term (i.e., 500 hours) and elevated temperature (i.e., 1500° F.) exposed sample. The SEM and EDS results are shown in
For the long-term and elevated temperature exposed sample (1500° F./500 h), the SEM micrographs shown in
The aging of LCHS-F included 1400° F./16 hours furnace cooling (FC) to 1200° F. and 1200° F./32 hours air cooling (AC).
The forging (pancaking) in this example was performed at 2050° F., at which temperature the P phase was stable and contributed to the grain size control. This was verified from the SEM micrograph of the aged LCHS-F sample (e.g.,
Based on
Another observation is that along the GB, a dark phase (from backscattered electron image (BSE)) was mixed with the bright phase. Based on EDS, the large GB dark phase is identified as gamma prime. It is expected that the growing of the GB phase (P or mu) depleted the nearby Mo and W and enriched the matrix with Al and Ti, which triggered the precipitation of large incoherent gamma prime on the grain boundary. This structure may be beneficial for the formation of GB gamma prime which breaks up the P/mu phase, preventing the P/mu phase from forming a continuous band around the GB (which may cause the alloy to be brittle). The gamma prime phase has a cubic Bravais lattice structure, and has many deformation systems common with the matrix, which gives more ductility to the GB. Accordingly, a hierarchical microstructure can be present, which is beneficial for achieving both high strength and good ductility in the alloy.
Overall, as the bright GB phase (P or mu) has a much higher solvus with the Al and Ti addition, it can be present at the typical forging temperature (2050° F.), at which temperature the material is soft and easy to work with, while the GB phase controls the grain size and prevents grain growth.
After aging, both fine gamma prime and Ni2M phases serve as strengthening phases. The addition of Al, Ti can further increase the stability of Ni2M phase. The Ni2M phase has a lower symmetry with 6 variants and a lath-like morphology; thus, it has a stronger strengthening effect in comparison to the gamma prime phase. The low misfit of the Ni2M phase helps it maintain coherency, which slows down coarsening and further strengthens with coherency stress.
Some Ni2M phases were visible in certain grains. The length of the Ni2M phase is below 1 micrometer.
The aged sample was then exposed at 1500° F. for 500 hours, resulting in a plurality of precipitates. Further, the needle-like Ni2M phase was coarsened (several micrometers to up to 10 micrometers in length).
Low CTE alloy LCHS-72 in accordance with an embodiment of the invention was scaled up to a ˜100 lbs vacuum induction melting (VIM) heat and forged into bars with a reduction ratio between 25:1 and 16:1. It was processed in accordance with the method described above in the section “Method of manufacturing”. The actual composition of LCHS-72 in weight % is listed in Table 8. The room temperature and elevated temperature tensile properties of LCHS-72 are listed in Table 9. In comparison to alloy LCHS-F, a small amount of strength was traded for higher ductility, especially at higher temperatures. The coefficient of thermal expansion of LCHS-72 was measured with two samples as shown in
All references, issued patents and patent applications cited within the body of the specification are hereby incorporated by reference in their entirety, for all purposes.
While the present invention has been described herein in detail in relation to one or more preferred embodiments, it is to be understood that this disclosure is only illustrative and exemplary of the present invention and is made merely for the purpose of providing a full and enabling disclosure of the invention. The foregoing disclosure is not intended to be construed to limit the present invention or otherwise exclude any such other embodiments, adaptations, variations, modifications or equivalent arrangements; the present invention being limited only by the claims appended hereto and the equivalents thereof.
This application claims priority to U.S. Provisional Application No. 63/464,083, filed on May 4, 2023, which is incorporated by reference herein in its entirety and for all purposes.
| Number | Date | Country | |
|---|---|---|---|
| 63464083 | May 2023 | US |
