METHOD OF PRE-AGING NITIHF SHAPE MEMORY ALLOYS AND PARTS THEREFROM WITH UNIFORM MICROSTRUCTURES AND SUPERIOR PROPERTIES

Abstract

A method to produce a high strength NiTiHf alloy, a NiTiHfZr alloy or a NiTiZr alloy are disclosed. The alloys comprise less than about 10 atomic percent of Hf, Hf+Zr, or Zr, respectively. The alloys, devices containing the alloys and methods of producing the devices are also disclosed.

Description

CROSS REFERENCE TO RELATED APPLICATIONS

This application claims priority and the benefit under 35 U. S.C. § 119(e) to U.S. Provisional Patent Application Ser. No. 62/609,848 filed Dec. 22, 2017, which is incorporated herein in its entirety by reference.

GOVERNMENT LICENSE RIGHTS

This invention was made with government support under grant number DE-SP0022534 awarded by the Department of Energy. The Government has certain rights in the invention.

FIELD OF THE INVENTION

The present invention relates to a two-step aging treatment comprised of a low-temperature pre-aging heat treatment step, followed by a typical high-temperature aging treatment, the latter similar to that established for higher Hf content alloys. The alloys produced by this treatment and products containing these alloys and methods of producing the products containing the alloys are also aspects of the invention.

BACKGROUND

NiTi-based shape memory alloys (SMAs) have been successfully used in many different fields of engineering because of their functional properties that include shape memory effect (SME) and superelasticity (SE). These properties are due to the occurrence of reversible, thermoelastic martensitic transformations. Because of inherent limitations to binary NiTi alloys, NiTiHf (i.e. nickel (Ni), titanium (Ti), and Hafnium (Hf)) alloys have primarily been proposed for use in high-temperature aerospace and automotive actuation applications where NiTi cannot perform, namely ambient environments greater than about 100° C., as Hf additions above about 10 atomic percent begin to increase the transformation temperatures. Hence, most research to date has focused on the shape memory and superelastic behavior of NiTiHf alloys with a high Hf content (generally on the order of 15-25 atomic percent.) For these high Hf containing materials, strength and shape memory behaviors are strongly enhanced by the formation of a precipitate reinforced microstructure, with the main strengthening precipitate commonly known as the H-phase. These precipitates are formed through a basic aging heat treatment, which is an effective way to increase the matrix strength in these alloys by forming fine precipitates that act as pinning sites against the movement of dislocations while still allowing the martensitic transformation to occur nearly unimpeded. The current practice for these high Hf-containing alloys is to age the solution treated material at 500 to 600° C. for between 1 and 4 hours, with the most common ageing treatment being 550° C. for 3 hours. This typical heat treatment is used to achieve high recoverable strain, high strength, excellent superelastic behavior, and microstructural and dimensional stability at temperatures greater than 100° C. in Ni-rich (>50 atomic percent Ni) NiTiHf alloys containing ≥15 atomic percent Hf. However, use of a single aging treatment, similar to that used for the high Hf alloys, has not produced the same uniform microstructures or resulted in sufficient mechanical and functional properties in Ni-rich NiTiHf alloys containing less than 10 atomic percent Hf.

SUMMARY

The present invention relates to a two-step aging treatment comprised of a low-temperature pre-aging heat treatment step, followed by a typical high-temperature aging treatment.

An aspect of the invention is a method to pre-age or heat treat low Hf (i.e. less than or equal to about 10 at. %) NiTiHf alloys, with compositions in the range of about 50 to about 53 atomic percent Ni, about 1 to about 10 atomic percent Hf, less than about 5 total atomic percent of incidental materials, which can include O, C, Fe, Mn, Cr, V, Nb, Mo, Ta, or W, or combinations thereof, and the balance of the alloy being Ti. The method comprises treating the NiTiHf alloys at a temperature between about 150° C. and about 400° C., for greater than about 10 minutes, in some embodiments between about 5 min and about 10 days, in some embodiments between about 10 minutes and 24 hours. Above temperatures of about 400° C., the method will produce primary precipitation, which is preferably avoided during this step. This pre-aging step results in homogeneous clusters of Ni and/or Hf atoms that act as nucleation sites for the precipitation of the H-phase, during subsequent aging at higher temperatures, resulting in a high strength NiTiHf alloy with a homogenous distribution of precipitates. Without the pre-aging treatment, the nucleation of H-phase tends to be heterogeneous, nucleating on such defects as grain boundaries, dislocations, and other microstructural features resulting in an inhomogeneous distribution of large precipitates (e.g. where the spacing between precipitates is more than about 1.5 times the size of the precipitates), which can be dependent upon the number of defects in the original material and can vary with processing. This inhomogeneous distribution of coarse precipitates results in poor mechanical and functional properties.

An aspect of the invention is a method to pre-age a NiTiHf alloy. The method comprises annealing the NiTiHf alloy at a temperature between about 700° C. and about 1150° C., followed by a pre-aging step at a temperature between about 100° C. and about 400° C., for between about 10 min and about 24 hours to produce a pre-aged NiTiHf alloy, and finally heat treating the pre-aged NiTiHf alloy at a temperature of between about 400° C. and about 600° C. to produce a homogenously distributed and high density of nano-sized H-phase precipitates throughout the matrix of a NiTiHf alloy. In some embodiments, the annealing can be solution annealing. The resulting alloy has improved mechanical stability and strength and at least equal to or greater than about 4% recoverable compression strain, in some embodiments between about 2% and 8% recoverable compression strain, without permanent deformation. The absence of a pre-aging step generally results in heterogeneous formations of larger H-phase precipitates, primarily clustered along grain boundaries, and correlated with poorer mechanical behavior.

An aspect of the invention is a method to produce a high-strength NiTi(Hf/Zr) alloy with homogeneous distribution of H-phase precipitates. The recoverable compression strain is between about 2 and about 8%. The method includes processing a NiTi(Hf/Zr) alloy by at least one method of casting, additive manufacturing, drawing, forging, extrusion, powder metallurgy, or combinations thereof. The composition of the NiTi(Hf/Zr) alloy includes between 50 and 53 atomic percent of Ni, between about 1 and 10 atomic percent of Hf, Zr, or combinations thereof, less than about 5 atomic percent total% of incidental materials, and the balance of the composition being Ti. The NiTi(Hf/Zr) workpiece is annealed to produce an annealed NiTi workpiece, which is then pre-aged at a temperature between about 100° C. and about 400° C. for a pre-aging time period between about 1 and 24 hours to produce the high-strength NiTi alloy.

An aspect of the invention is a NiTi alloy. The recoverable yield of the alloy is between about 2 and about 8%, and the compressive yield strength of the alloy is greater than about 1.5 GPa.

An aspect of the invention is a biomedical implant. The implant includes a NiTi alloy, wherein a composition of the NiTi alloy is Ni_50.3-53.Ti_49.7-xHf_xZr_y wherein x is between 6 and 9, 3 and 6, or 1 and 3, wherein x+y is between 1 and 10, and wherein y is between 0 and 10.

BRIEF SUMMARY OF THE DRAWINGS

The patent or application file contains at least one drawing executed in color. Copies of this patent or patent application publication with color drawing(s) will be provided by the Office upon request and payment of the necessary fee.

FIG. 1A illustrates the mechanical behavior of Ni_50.3Ti_42.7Hf₆ (at. %) alloys in compression with pre-aging treatment at 300° C. for 12 hours followed by normal aging at 550° C. for 3.5 hours for all samples. The solid arrows show the amount of strain recovered upon heating the deformed samples to 150° C.;

FIG. 1B illustrates the mechanical behavior of Ni_50.3Ti_41.7Hf₈ (at. %) alloys in compression with pre-aging treatment at 300° C. for 12 hours followed by normal aging at 550° C. for 3.5 hours for all samples. The solid arrows show the amount of strain recovered upon heating the deformed samples to 150° C.;

FIG. 1C illustrates the mechanical behavior of Ni_50.3Ti_41.2Hf_8.5 (at. %) alloys in compression with pre-aging treatment at 300° C. for 12 hours followed by normal aging at 550° C. for 3.5 hours for all samples. The solid arrows show the amount of strain recovered upon heating the deformed samples to 150° C.;

FIG. 1D illustrates the mechanical behavior of Ni_50.3Ti_40.7Hf₉ (at. %) alloys in compression with pre-aging treatment at 300° C. for 12 hours followed by normal aging at 550° C. for 3.5 hours for all samples. The solid arrows show the amount of strain recovered upon heating the deformed samples to 150° C.;

FIG. 1E illustrates the mechanical behavior of Ni_50.3Ti_42.7Hf₆ (at. %) alloys in compression without pre-aging treatment at 300° C. for 12 hours, but including normal aging at 550° C. for 3.5 hours for all samples. The solid arrows show the amount of strain recovered upon heating the deformed samples to 150° C.;

FIG. 1F illustrates the mechanical behavior of Ni_50.3Ti_41.7Hf₈ (at. %) alloys in compression without pre-aging treatment at 300° C. for 12 hours, but including normal aging at 550° C. for 3.5 hours for all samples. The solid arrows show the amount of strain recovered upon heating the deformed samples to 150° C.;

FIG. 1G illustrates the mechanical behavior of Ni_50.3Ti_41.2Hf_8.5 (at. %) alloys in compression without pre-aging treatment at 300° C. for 12 hours, but including normal aging at 550° C. for 3.5 hours for all samples. The solid arrows show the amount of strain recovered upon heating the deformed samples to 150° C.;

FIG. 1H illustrates the mechanical behavior of Ni_50.3Ti_40.7Hf₉ (at. %) alloys in compression without pre-aging treatment at 300° C. for 12 hours, but including normal aging at 550° C. for 3.5 hours for all samples. The solid arrows show the amount of strain recovered upon heating the deformed samples to 150° C.;

FIG. 2A illustrates a conventional bright field image of Ni_50.3Ti_41.2Hf_8.5 (at. %) after pre-aging at 300° C. for 12 hours. Corresponding selected area diffraction pattern along [111] zone axis is illustrated in the lower left inset;

FIG. 2B illustrates a high resolution TEM micrograph along [111] zone axis of Ni_50.3Ti_41.2Hf_8.5 (at. %) alloy pre-aged at 300° C. for 12 hours and subsequently aged at 550° C. for 3.5 hours. Some of the H-phase precipitates are indicated by arrows marked with the letter “P”. The corresponding Fast Fourier transform (FFT) pattern is illustrated in the lower right inset;

FIG. 3A illustrates conventional Bright Field (BF) images of Ni_50.3Ti_42.7Hf₆ (at. %) at 50 nm after aging at 550° C. for 3.5 hours without the pre-aging treatment, with corresponding Selected Area Electron Diffraction (SAED) patterns along [111] zone axis for each region presented in the inset;

FIG. 3B illustrates conventional BF images of Ni_50.3Ti_42.7Hf₆ (at. %) at 200 nm after aging at 550° C. for 3.5 hours without the pre-aging treatment, with corresponding SAED patterns along [111] zone axis for each region presented in the inset;

FIG. 3C illustrates conventional BF images of Ni_50.3Ti_41.2Hf_8.5 (at. %) at 100 nm after aging at 550° C. for 3.5 hours without the pre-aging treatment, with corresponding SAED patterns along [111] zone axis for each region presented in the inset; and

FIG. 3D illustrates conventional BF images of Ni_50.3Ti_40.7Hf₉ (at. %) at 50 nm after aging at 550° C. for 3.5 hours without the pre-aging treatment, with corresponding SAED patterns along [111] zone axis for each region presented in the inset.

DETAILED DESCRIPTION

The present invention generally relates to a two-step aging treatment of metal alloys, and the material formed with this treatment. Zirconium behaves similarly as Hf in Ni-rich NiTi-based alloys forming the same H-phase precipitates and consequently ternary and quaternary alloys with compositions in the range of between about 50 and about 53 atomic percent Ni, between about 1 and about 10 atomic percent Hf+Zr (where zirconium (Zr) can range from about 0 to about 10 atomic percent), and the balance being Ti and any incidental material (i.e. O, C, Fe, Mn, Cr, V, Nb, Mo, Ta, or W, and combinations thereof), would behave in a similar manner as that described above, thus also benefitting from the pre-aging treatment disclosed.

The benefit of the pre-aging treatment is clearly demonstrated by the data in Table 1. Table 1 includes mechanical data for various Ni-rich, NiTiHf alloys with and without the pre-aging treatment. The pre-aged samples exhibited greater recoverable (and superelastic) strains when loaded to about 2 GPa in compression and also did not undergo any permanent (plastic) strain when loaded to that level, indicating that the yield strength of the material was greater than about 2 GPa. Table 1 illustrates the benefits of the pre-aging treatment to enhance the strength, and improve recoverable and superelastic strains of the various alloys compared to the same material without the pre-aging treatment. Each value in Table 1 is approximate.

TABLE 1
		Recoverable	Permanent
		strain at	strain at 2 GPa
Composition (at. %)	Heat treatment	2 GPa (%)	(%)
Ni51.5Ti42.5Hf6	With pre-aging	4.3	0
	Without pre aging	1.8	0
Ni51Ti41Hf8	With pre-aging	4.75	0
	Without pre-aging	2.5	0
Ni50.3Ti43.7Hf6	With pre-aging	3.8	0
	Without pre aging	2.4	5
Ni50.3Ti41.7Hf8	With pre-aging	4	0
	Without pre aging	3.4	3
Ni50.3Ti41.2Hf8.5	With pre-aging	4	0
	Without pre aging	4	2

One aspect of the present invention is a method to produce a high-strength NiTiHf alloy with homogeneous distribution of H-phase precipitates, where the alloy comprises greater than or equal to about 4% recoverable compression strain. The method includes processing an alloy by at least one process of casting, additive manufacturing, drawing, forging, extrusion, powder metallurgy, or other typical metallurgical process, or combinations thereof. The composition of the alloy includes between about 50 and about 53 atomic percent Ni, between about 1 and about 10 atomic percent of Hf, less than about 1 atomic percent of incidental materials (e.g. O, C, Fe, Mn, Cr, V, Nb, Mo, Ta, or W, or combinations thereof), and the balance being Ti. Following the processing step, the alloy is solution treated at 700° C. and about 1150° C. followed by pre-aging at a temperature between about 100° C. and about 400° C. for at least about 10 min.

In some embodiments, the temperature of the pre-aging step can be between about 300° C. and about 400° C., or about 200° C., about 250° C., about 300° C., about 350° C., about 400° C., about 450° C. The alloy can be maintained at the temperature for between about 10 min and 24 hours, in some embodiments between about 3 hours and about 18 hours, between about 5 hours and about 12 hours, or between about 7 hours and about 10 hours. In some embodiments, the alloy can be maintained at the pre-aging temperature for about 10, min, 30 min, one hour, about two hours, about three hours, about four hours, about five hours, about six hours, about seven hours, about eight hours, about nine hours, about 10 hours, about 11 hours, about 12 hours, about 13 hours, about 14 hours, about 15 hours, about 16 hours, about 17 hours, about 18 hours, about 19 hours, about 20 hours, about 21 hours, about 22 hours, about 23 hours, or about 24 hours.

The NiTiHf alloy can comprise several compositions. In some embodiments, the alloy comprises Ni_50.3-53Ti_49.7-xHf_x wherein x is between 6 and 9, 3 and 6, or 1 and 3. Incidental materials can be included in the NiTiHf alloy, which can include up to about 1 atomic percent of any transition metals (individually) or up to about 1 atomic % of a nonmetal such as oxygen or carbon, and combinations thereof. The maximum amount of the incidental materials can be about up to about 5 atomic percentage. In some embodiments, the amount of incidental materials can be between about 1 atomic percentage and about 5 atomic percentage. In some embodiments, the Hf can be a combinations of Hf and Zr, wherein the sum of Hf+Zr is between about 1 and about 10 atomic percentage and wherein the contribution of Zr can be between about 0.1 and about 10 atomic percentage (with 10 atomic percentage resulting in Ni_50.3-53Ti_49.7-xZrx).

The method can further include an annealing step prior to the pre-aging step. The annealing temperature is between about 700° C. and about 1150° C. In some embodiments, the annealing temperature can be between about 750° C. and about 1100° C., about 800° C. and about 1050° C., about 850° C. and about 1000° C., or about 900° C. and about 950° C. In some embodiments, the annealing temperature can be about 700° C., about 750° C., about 800° C., about 850° C., about 900° C., about 950° C., about 1000° C., about 1050° C., about 1100° C., or about 1100° C. The alloy can be maintained at the annealing temperature for between about 30 minutes and about 144 hours. In some embodiments, the annealing duration can be about 72 hours for cast material. Following annealing, the alloy can be quenched at a temperature between about −100° C. and about 200° C.

In some embodiments, the method can further include a heat treatment following the pre-aging step. The pre-aged alloy can be heat treated at a temperature between about 400° C. and about 600° C. for between about 10 minutes and about 24 hours. In some embodiments, the heat treatment temperature can be between about 500° C. and about 550° C.

An aspect of the invention is a method to produce a NiTiHf alloy. The method includes processing an alloy by at least one method of casting, additive manufacturing, drawing, forging, extrusion, powder metallurgy, or other typical metallurgical process, or combinations thereof. The composition of the NiTiHf alloy is between about 50 and about 53 atomic percent or Ni, between about 1 and about 10 atomic percent of Hf, no greater than about 2% atomic percent (total) of incidental materials, and the balance being Ti. The NiTiHf alloy is solution annealed at a temperature between about 700° C. and about 1150° C. to produce a solution annealed NiTiHf alloy. The solution annealed NiTiHf alloy is pre-aged at a pre-aging temperature, which is between about 100° C. and about 400° C. for a time period between about 1 hour and about 24 hours to produce a pre-aged NiTiHf workpiece. The pre-aged NiTiHf workpiece is heat treated at a heat treatment temperature between about 400° C. and about 600° C. to produce a high density of homogeneously distributed H-phase precipitates in the NiTiHf part or form.

In some embodiments, the temperature of the pre-aging step can be between about 300° C. and about 600° C., between about 400° C. and about 550° C., or about 200° C., about 250° C., about 300° C., about 350° C., about 400° C., about 450° C., about 500° C., about 600° C., about 650° C., or about 700° C. The alloy can be maintained at the temperature for between about one hour and 24 hours, in some embodiments between about 3 hours and about 18 hours, between about 5 hours and about 12 hours, or between about 7 hours and about 10 hours. In some embodiments, the alloy can be maintained at the pre-aging temperature for about one hour, about two hours, about three hours, about four hours, about five hours, about six hours, about seven hours, about eight hours, about nine hours, about 10 hours, about 11 hours, about 12 hours, about 13 hours, about 14 hours, about 15 hours, about 16 hours, about 17 hours, about 18 hours, about 19 hours, about 20 hours, about 21 hours, about 22 hours, about 23 hours, or about 24 hours.

The annealing temperature is generally between about 700° C. and about 1150° C. In some embodiments, the annealing temperature can be between about 750° C. and about 1100° C., about 800° C. and about 1050° C., about 850° C. and about 1000° C., or about 900° C. and about 950° C. In some embodiments, the annealing temperature can be about 700° C., about 750° C., about 800° C., about 850° C., about 900° C., about 950° C., about 1000° C., about 1050° C., about 1100° C., or about 1100° C. The alloy can be maintained at the annealing temperature for between about 30 minutes and about 144 hours, in some embodiments about 72 hours. Following annealing, the alloy can be quenched at a temperature between about −100° C. and about 200° C.

The pre-aged alloy can be heat treated at a temperature between about 400° C. and about 600° C. for between about 30 minutes and about 24 hours. In some embodiments, the heat treatment temperature can be between about 500° C. and about 550° C.

The NiTiHf alloy can comprise several compositions. In some embodiments, the alloy comprises Ni_50.3-53Ti_49.7-xHf_x wherein x is between 6 and 9, 3 and 6, or 1 and 3. Incidental materials can be included in the NiTiHf alloy, which can include up to about 1 atomic percent of any transition metals (individually) or up to about 1 atomic % of a nonmetal such as oxygen or carbon and combinations thereof. The maximum amount of the incidental materials can be about 5 atomic percentage. In some embodiments, the amount of incidental materials can be between about 1 atomic percentage and about 5 atomic percentage. In some embodiments, the composition can be Ni_50.3Ti_49.7-xHf_x, where x can be between 6 and 9, in some embodiments x can be 6, 8, 8.5 or 9.

An aspect of the invention is a method to produce a NiTiHfZr alloy. The method includes processing an alloy workpiece by a method from the group consisting of casting, additive manufacturing, drawing, forging, extrusion, powder metallurgy, or other typical metallurgical process or combinations thereof. The NiTiHfZr alloy comprises between about 50 and about 53 atomic percent of Ni, between about 1 and 10 atomic percent Hf+Zr (where Zr can range from about 0 to about 10 atomic percent), less than about 2 atomic percentage of incidental materials, and the balance being Ti. The NiTiHfZr workpiece is solution annealed at a temperature of between about 700° C. and about 1150° C. to produce a solution annealed NiTiHfZr workpiece. The annealed NiTiHfZr alloy is pre-aged at a pre-aging temperature of between about 100° C. and about 400° C. for between about 1 hour and about 24 hours to produce a pre-aged NiTiHfZr alloy. The pre-aged NiTiHf Zr alloy is heat treated at a temperature between about 400° C. and about 600° C. to produce a high density of homogeneously distributed H-phase precipitates in the NiTiHfZr alloy.

In some embodiments, the temperature of the pre-aging step can be between about 100° C. and about 300° C., between about 300° C. and about 350° C., or about 200° C., about 250° C., about 300° C., about 350° C., or about 400° C. The alloy can be maintained at the temperature for between about one hour and 24 hours, in some embodiments between about 3 hours and about 18 hours, between about 5 hours and about 12 hours, or between about 7 hours and about 10 hours. In some embodiments, the alloy can be maintained at the pre-aging temperature for about one hour, about two hours, about three hours, about four hours, about five hours, about six hours, about seven hours, about eight hours, about nine hours, about 10 hours, about 11 hours, about 12 hours, about 13 hours, about 14 hours, about 15 hours, about 16 hours, about 17 hours, about 18 hours, about 19 hours, about 20 hours, about 21 hours, about 22 hours, about 23 hours, or about 24 hours.

The annealing temperature is between about 700° C. and about 1150° C. In some embodiments, the annealing temperature can be between about 750° C. and about 1100° C., about 800° C. and about 1050° C., about 850° C. and about 1000° C., or about 900° C. and about 950° C. In some embodiments, the annealing temperature can be about 700° C., about 750° C., about 800° C., about 850° C., about 900° C., about 950° C., about 1000° C., about 1050° C., about 1100° C., or about 1100° C. The alloy can be maintained at the annealing temperature for between about 30 minutes and about 144 hours, in some embodiments about 72 hours. Following annealing, the alloy can be quenched at a temperature between about −100° C. and about 200° C.

The pre-aged alloy can be heat treated at a temperature between about 400° C. and about 600° C. for between about 1 minutes and about 24hours. In some embodiments, the heat treatment temperature can be between about 500° C. and about 550° C.

An aspect of the invention is a NiTiHf alloy, NiTiZr alloy or a NiTiHfZr alloy, wherein the alloy comprises a homogenously distributed H-phase in the precipitate.

In some embodiments, the compression yield strength of the NiTiHf alloy can be greater than about 1.5 GPa. In some embodiments, the compressive yield strength can be between about and about 8% recoverable compression strain. In some embodiments, the compression yield strength can be greater than about 2 GPa when tested at about room temperature. In some embodiments, the compressive strength can be up to about 4 GPa, in some embodiments between about 1.5 GPa and 4 GPa. The recoverable compression strain can be between about 2% and about 8%, in some embodiments about 4%. The tensile strength can be between about 60% and 70% of the compressive strength.

An aspect of the invention is a method for forming a biomedical implant. The method includes providing a NiTiHfZr alloy, a NiTiZr alloy or a NiTiHf alloy; and producing the biomedical implant by additive manufacturing.

An aspect of the invention is a biomedical implant comprising a NiTiHfZr alloy, NiTiZr alloy or NiTiHf alloy.

EXAMPLES

Example 1

NiTiHf alloys with target compositions of Ni_50.3Ti_50-xHf_x, with x=6, 8, 8.5, and 9 at. % were made by induction-melting high-purity elemental constituents using a graphite crucible and casting into a copper mold. The ingots were homogenized in vacuum at 1050° C. for 72 hours and then extruded at about 900° C. at a 7:1 area reduction ratio. The extruded rods were sectioned into samples that were initially solution-annealed at 1050° C. for 30 min, water quenched, and then pre-aged at 300° C. for 12 h and air-cooled, and finally aged a second time at 550° C. for 3.5 hours and air-cooled. To isolate the effect of pre-aging on the functional properties of NiTiHf alloys, other test samples were directly aged at 550° C. for 3.5 hours after the solution-anneal treatment (without pre-aging at 300° C. for 12 hours).

Mechanical compression tests were performed on an MTS servo-hydraulic load-frame equipped with an MTS 661.20 load cell. Compression samples were cylindrical with a diameter of 5 mm and a length of 10 mm. Five compression cycles were applied to the samples using a maximum load of 40 kN and a minimum load of 250 N, corresponding to 2 GPa and 13 MPa engineering stress limits. A cross-head speed of 0.1 mm/min was used, corresponding to an approximate strain rate of 10-4 s⁻¹. The surfaces of the samples were speckled using an airbrush to deposit sequential layers of alumina powders (≤10 μm) and Brother TB450 carbon black toner powder. Digital images were acquired during loading using a Basler ACA2500 camera, and the Ncorr digital image correlation (DIC) software was used to analyze the displacements of the particles. From these displacement fields, the software calculated the surface strains during deformation. Before each test, eight images of the undeformed sample were acquired and analyzed to establish the strain noise for each pattern, which fell between 10⁻⁴ to 10⁻⁵ for the data in FIG. 1. Because these figures are compared to each other, they are provided as a single figure with sub-figures.

Conventional and high-resolution transmission electron microscopy (HRTEM or TEM) of aged NiTiHf samples was carried out using an FEI Talos TEM (FEG, 200 kV). The TEM samples were prepared by grinding slices to a thickness of 90-100 μm; a mechanical punch was then used to create discs with a diameter of 3 mm. A Fischione automatic twin-jet electropolisher (model 120) at 13 V and an electrolyte of 30% HNO₃ in methanol (by volume) at around −35° C. was then used to further thin the TEM foils. To measure the size of H-phase precipitates and interparticle distance (the distance of a single precipitate from its closest precipitate), several HRTEM images taken from various regions, were used. This measurement was repeated for almost 100 precipitates on each sample and average precipitate size, average interparticle distance and their corresponding standard error is reported.

FIG. 1 illustrates the compression responses of the NiTiHf alloys aged at 550° C. for 3.5 hours, with (FIGS. 1A-1D) and without (FIGS. 1E-1H) an initial pre-aging treatment of 300° C. for 12 hours. The compression tests were performed at room temperature (about 23° C.), and all the samples were initially austenite. The mechanical loading stress-induced martensite that did not recover upon unloading at room temperature. Heating the samples to about 150° C. and measuring the recovered strains assessed how much of the deformation was due to stress induced martensite vs. plasticity, in a phenomenological sense. Complete recovery of the strain (almost 4%) was observed for all the NiTiHf samples that were pre-aged at 300° C. for 12 hour (solid arrows in FIG. 1A-1D). For the other samples with 6 to 8.5 at. % Hf (FIGS. 1E-1F), only a portion of the strain was recovered upon heating, indicating a stronger presence of plasticity in the mechanical responses sans pre-aging. For the Ni_50.3Ti_40.7Hf₉ sample (FIG. 1G), however, the strain was fully recovered sans pre-aging, as well as in the pre-aged case. While superelastic behavior was not directly observed, these data do suggest that the pre-aged samples and the un-pre-aged 9Hf sample would likely exhibit stable superelastic responses at higher test temperatures.

FIG. 2A illustrates a representative bright field (BF) TEM image of the Ni_50.3Ti₄L₂Hf_8.5 sample after pre-aging at 300° C. for 12 hours. Fully formed H-phase precipitates were not observed. The corresponding selected area electron diffraction pattern (SAED) of the B2 zone (upper right inset in FIG. 2A) confirms this observation, as the characteristic super-reflection of H-phase precipitates along <110> directions were not detected; however, a clear structured pattern of diffuse intensity is present. This pattern of diffuse intensity was also observed for other samples with different Hf content after pre-aging at 300° C. for 12 hours. A HRTEM study of the sample further corroborated the lack of fully formed H-phase precipitates after pre-aging at 300° C. for 12 hour (as demonstrated in the lower left inset in FIG. 2A).

These diffuse intensities with periodical character in reciprocal space indicate the existence of short-range order in the real-space lattice. Such diffuse intensities have been reported in binary Ni_50.6T_149.4 after low-temperature aging and were attributed to the existence of microdomains in the form of clusters of Ni atoms as precursors to full formation of Ni₄Ti₃ nanoprecipitates. Analysis of H-phase precipitates in NiTiHf alloys indicates that the Hf content of the H-phase is higher than the Hf content in the matrix while the Ni content is also slightly higher in the precipitate compared with the matrix. It is therefore expected that Hf and/or Ni atom clusters form after low-temperature aging (pre-aging treatment) as a precursor to H-phase precipitation upon further aging at higher temperatures, similar to the sequence of mechanisms that give rise to Ni₄Ti₃ nano-precipitation in binary NiTi alloys.

FIG. 2B is a HRTEM micrograph of the Ni_50.3Ti₄L₂Hf_8.5 alloy pre-aged at 300° C. for 12 hours and subsequently aged at 550° C. for 3.5 h. The corresponding fast Fourier transform (FFT) is illustrated in the upper left inset. The primary spots in the FFT result from the B2 cubic austenite structure, and the super reflections at ⅓ positions along <110> B2, are reflections from uniquely oriented H-phase precipitates (depicted in FIG. 2B with the letter “P”). The H-phase precipitate dimensions for Ni50.3Ti49.7-xHfx alloys (where x is 6, 8, 8.5, and 9 atomic percentage, respectively) with and without pre-aging at 300° C. for 12 hours followed by normal aging at 550° C. for 3.5 hours are depicted in Table 2. For the latter cases (without pre-aging), characteristic precipitate morphologies are described for both regions at grain boundaries (GB) as well as within the grain interior (IG). All values in Table 2 are approximate.

	TABLE 2
	Without pre-aging
With pre-aging	Ni50.3Ti42.7Hf6	Ni50.3Ti42.7Hf6	Ni50.3Ti41.2Hf8.5	Ni50.3Ti41.2Hf8.5	Ni50.3Ti40.7Hf9
All alloys	(GB)	(IG)	(GB)	(IG)	(GB & IG)
Length	14 ± 1	60 ± 4	110 ± 10	43 ± 1	93 ± 2	23 ± 1
(nm)
Width	7 ± 1	15 ± 2	57 ± 5	12 ± 1	25 ± 1	10 ± 1
(nm)
Inter-	8 ± 1	5 ± 1	95-700	5 ± 1	100 ± 4	12 ± 1
particle
distance
(nm)

For all the pre-aged samples, regardless of Hf content, the precipitates were ellipsoidal in shape with average dimensions of 14±1 nm (length) and 7±1 nm (width); the interparticle distance was 8±1 nm. Obvious variation in H-phase precipitate morphologies were not detected for any of the pre-aged NiTiHf alloys. Therefore, the fully reversible 4% deformation depicted in FIGS. 1A-1D is aided by the uniform distribution of finely spaced H-phase precipitates, which strengthen the matrix against plastic deformation, while not impeding the formation of martensite.

The microstructure of the Ni_50.3Ti_42.7Hf₆ alloy that was aged at 550° C. for 3.5 hours (without pre-aging) is illustrated in the BF-TEM micrographs of FIGS. 3A and 3B taken from the vicinity of a grain boundary (GB) and the interior of the austenite grain, respectively. It is apparent that aging of the Ni_50.3Ti_42.7Hf₆ alloy at 550° C. for 3.5 hours induces heterogeneous nucleation of spindle-like H-phase precipitates along GBs (with an aspect ratio of 4) and a sparse distribution of larger, longer, and widely spaced ellipsoidal H-phase precipitates (with an aspect ratio of 2) in the interior of the grains, though the major part of the grain shown in FIG. 3B is actually free of precipitates. GBs are well known locations for heterogeneous nucleation in solid state precipitation processes because they decrease the interfacial energy between the precipitate and the parent phase, reduce stress fields, and because of the chemical composition gradient that exists around GBs. However, because there is competitive growth on the GBs, the size of the grain boundary precipitates is small compared with the precipitates formed in the grain interiors.

With an increase in Hf content, the interparticle distance and size of the precipitates formed on the GB's and inside the grains decreases, as illustrated in FIG. 3C for the Ni_50.3Ti_41.2Hf_8.5 sample (Table 2). This suggests that there is transition from heterogeneous precipitation in the 550° C. aged NiTiHf alloys (for low Hf content) to homogeneous precipitation in alloys with higher Hf content for the same Ni:(Ti+Hf) ratio. The microstructure of Ni_50.3Ti_40.7Hf₉ after aging at 550° C. for 3.5 hours (without pre-aging) is illustrated in FIG. 3D.

Clearly, the H-phase precipitates are homogenously nucleated and distributed in grain interiors and regions near GBs; heterogeneous GB precipitation is not observed. Consistent with the microstructural observations, the amount of permanent strain observed during compression testing decreased as the Hf content increased from 5.7% (FIG. 1E for Ni_50.3Ti_42.7Hf₆) to 2% (FIG. 1G for Ni_50.3Ti_41.2Hf_8.5), and finally, full recovery without permanent deformation was observed for the Ni_50.3Ti_40.7Hf₉ alloy, where precipitation was homogeneous and finely spaced.

Moreover, the H-phase precipitate dimensions of Ni_50.3Ti_40.7Hf₉ alloy, with the application of pre-aging treatment (Table 2), are slightly smaller as is the interparticle distance compared to the same alloy only aged at 550° C. for 3.5 hours (without pre-aging). Therefore, the lower critical plateau stress of the Ni_50.3Ti_40.7Hf₉ sample without the pre-aging treatment (300 MPa in FIG. 1H) compared with the same material with pre-aging treatment (520 MPa in FIG. 1D) could be due to a larger interparticle distance where martensite can propagate more easily between particles. Other possibilities include a different H-phase precipitate chemistry between the two conditions, resulting in different matrix chemistries, hence different transformation temperatures, and/or the slightly larger coherent precipitates cause larger coherency strains, hence more stress and lower transformation stresses.

Finally, the stresses required to form martensite (the plateau stresses) for the samples that were pre-aged at 300° C. for 12 hours were relatively consistent, falling between about 410 and 510 MPa (FIGS. 1A-1D), yet in contrast, the martensite formation plateau stresses for the samples without the pre-aging treatment (FIG. 1E-1H) continuously decreased from about 810 MPa to about 300 MPa as Hf content increased from 6 to 9 at. %. In the former case, it is possible that the chemical diffusion mechanism that is promoted by pre-aging also promotes more available hafnium to go into H-phase precipitates as the overall Hf-content is increased for each different composition. In other words, the extra time for diffusion results in an increase in the Hf content that goes into the precipitates themselves, but a relatively uniform Hf and/or Ni content left in the matrix (or at least ratios that balance each other in terms of transformation temperature effects). Contrarily, the changes observed in the plateau stresses in the latter case (no pre-aging) are logically expected from the increase in overall Hf content of each alloy, given that for most of the compositions, H-phase precipitation is not uniform enough throughout the matrix to effect the transformation temperatures greatly. Increasing the Hf content in NiTiHf alloys increases the Martensite start (Ms) temperature, thus through the Clausius-Clapeyron relation, the martensite formation stresses would decrease given a constant test temperature. In the higher Hf contents where there is relatively uniform H-phase precipitation, the combined mechanisms discussed in the previous paragraph likely all interact to still cause a relative decrease in transformation temperatures relative to the lower Hf alloys.

Pre-aging Ni_50.3Ti_429.7-xHf_x alloys (x=6, 8, 8.5, 9 at. %) at about 300° C. for about 12 hours after solution annealing at about 1050° C. leads to a higher density of the H-phase precipitates during subsequent aging at 550° C., than aging alone. This results in greater strength and subsequently improved mechanical and functional performance with 4% recoverable compression strain. The results suggest that Hf or other atom clusters form during the pre-aging treatment resulting in nucleation of a uniform distribution of H-phase when subsequently aged at 550° C. Aging NiTiHf alloys at 550° C. directly after the solution annealing treatment (without pre-aging), leads to heterogeneous nucleation of H-phase precipitates on the grain boundaries in the lower Hf compositions (6 to 8 at. %). Most of the grain interior is free of the precipitates, which results in reduced strength and leads to poor mechanical behavior. However, as Hf content is increased, homogenous precipitation is achieved directly, without pre-aging. As a result, mechanical and functional behavior is improved compared to the compositions with lower Hf content and no pre-aging.

Ranges, for example temperature ranges, atomic percentages, and others, have been discussed and used within the forgoing description. One skilled in the art would understand that any sub-range within the stated range would be suitable, as would any number within the broad range, without deviating from the invention.

The foregoing description of the present invention related to NiTiHf alloys has been presented for purposes of illustration and description. Furthermore, the description is not intended to limit the invention to the form disclosed herein. Consequently, variations and modifications commensurate with the above teachings, and the skill or knowledge of the relevant art, are within the scope of the present invention. The embodiment described hereinabove is further intended to explain the best mode known for practicing the invention and to enable others skilled in the art to utilize the invention in such, or other, embodiments and with various modifications required by the particular applications or uses of the present invention. It is intended that the appended claims be construed to include alternative embodiments to the extent permitted by the prior art.

Claims

1. A method to produce a high-strength NiTi(Hf/Zr) alloy with homogeneous distribution of H-phase precipitates, with between about 2 and about 8% recoverable compression strain, comprising:

processing a NiTi(Hf/Zr) alloy by at least one method of a casting method, an additive manufacturing method, a drawing method, a forging method, an extrusion method, a powder metallurgy method, or combinations thereof, wherein the composition of the NiTi(Hf/Zr) alloy comprises between 50 and 53 atomic percent of Ni, between about 1 and 10 atomic percent of Hf, Zr, or combinations thereof, less than about 5 atomic percent total % of incidental materials, and the balance of the composition being Ti;

annealing the NiTi(Hf/Zr) workpiece to produce an annealed NiTi workpiece; and

pre-aging the annealed NiTi workpiece at a temperature between about 100° C. and about 400° C. for a pre-aging time period between about 1 and 24 hours to produce the high-strength NiTi alloy.

2. The method of claim 1, wherein the pre-aging temperature is about 300° C.

3. The method of claim 2, wherein the pre-aging time period is about 12 hours.

4. The method of claim 1, further comprising quenching the annealed NiTi workpiece at a temperature between about −100° C. and about 200° C. prior to pre-aging.

5. The method of claim 1, wherein a compressive yield strength of the high-strength NiTi alloy is between about 2 GPa and about 4 GPa.

6. The method of claim 1, further comprising a heat treatment at a heat treatment temperature of between about 400° C. and about 600° C. for between about 10 minutes and about 24 hours.

7. The method of claim 1, wherein the annealing is solution annealing.

8. The method of claim 7, wherein the solution annealing temperature is between about 700° C. and about 1150° C.

9. The method of claim 1, wherein a primary precipitate is not formed following the preaging step.

10. The method of claim 1, wherein the NiTi alloy is of the formula Ni50+x(Ti+Hf)49.9-1x wherein x is between 0.1 and 2.9.

11. The method of claim 10, wherein Hf in the NiTiHf alloy is between about 0.1 to 10%.

12. The method of claim 1, wherein the NiTi alloy is of the formula Ni50+x(Ti+Zr)49.9-1x wherein x is between 0.1 and 2.9.

13. The method of claim 10, wherein Zr in the NiTi alloy is between about 0.1 to 10%.

14. The method of claim 10, wherein the NiTi alloy further comprises Zr, wherein the Zr in the NiTi alloy is between about 0 and about 10 atomic percentage of a percentage of Hf in the NiTiHf alloy of a composition NiTiHfZr.

15. A NiTi alloy, wherein a recoverable yield of the alloy is between about 2 and about 8%, and wherein a compressive yield strength of the alloy is greater than about 1.5 GPa.

16. The NiTi alloy of claim 15, wherein a compressive yield strength of the high-strength NiTi alloy is between about 2 GPa and about 4 GPa.

17. The NiTi alloy of claim 15, wherein a composition of the NiTi alloy is Ni50.3-53.Ti49.7-xHfx wherein x is between 6 and 9, 3 and 6, or 1 and 3.

18. The NiTi alloy of claim 15, further comprising incidental materials of up to about 1 atomic percent of any individual transition metals, up to about 1 atomic % of a nonmetal, and combinations thereof, wherein a total maximum amount of the incidental materials is about 5 atomic percentage.

19. A biomedical implant, comprising a NiTi alloy, wherein a composition of the NiTi alloy is Ni50.3-53.Ti49.7-xHfx wherein x is between 6 and 9, 3 and 6, or 1 and 3.