Polymer-like ultrahigh-strength metal alloy

Abstract

Disclosed in the present disclosure is a polymer-like ultrahigh-strength metal alloy, wherein the alloy has the following levels of strength and Young's modulus over a wide temperature range of 20° C. to 80° C. or −80° C. to 80° C.: a ultra-high yield strength σy of greater than or equal to 1 GPa, and a Young's modulus E of about 7-20 GPa. Thus, a novel polymer-like ultrahigh-strength metal alloy with a polymer-like ultralow Young's modulus is achieved in the present disclosure.

Description

CROSS-REFERENCE TO RELATED APPLICATIONS

[0001] This application claims priorities from U.S. provisional application 63 / 692,329 filed Sep. 9, 2024, Chinese patent applications 2024112595717 filed Sep. 9, 2024, and 202511286615X filed Sep. 9, 2025, the content of which are incorporated herein in the entirety by reference.TECHNICAL FIELD

[0002] The present disclosure relates to the technical field of new materials, in particular to a polymer-like ultrahigh-strength metal alloy.BACKGROUND

[0003] Dissimilar metal alloys showing steel-like ultrahigh strength and polymer-like ultrahigh flexibility simultaneously have long been desired for many emerging technologies such as morphing aircrafts and superman-type artificial muscles. Such an unconventional combination of “strong yet flexible” properties (comparable to a red region in FIG. 1a) will enable a large recoverable shape change under a small driving force and simultaneously provide strong resistance against fracture or yielding failure under large loads. Thus, such alloys can be used in a morphing wing in an aircraft or a super-strong artificial muscle in a humanoid robot, and the like.

[0004] However, achieving such unconventional alloys has remained challenging because of a trade-off relationship between strength and flexibility (the latter being measured by elastic compliance or an inverse of an elastic modulus), as shown in a grey band in which all the ellipses are covered (referred to as an Ashby plot) in FIG. 1a. The figure reveals a common phenomenon that steels can be made very strong (with a yield strength σy>1 GPa) but they are stiff (with a Young's modulus E of about 200 GPa), whereas organic materials such as fiber-reinforced polymers (FRPs) have the opposite property of being flexible (with a typical low E of about 10 GPa) but weak (σy<0.3 GPa). The strength-flexibility trade-off stems from the fact that the two properties are oppositely correlated to the bonding strength of a material, which is represented roughly by the elastic modulus. Therefore, this inevitable trade-off relationship makes it impossible for a 1 GPa-class steel-like high-strength alloy to achieve a 10 GPa-class polymer-like low modulus.

[0005] Over the past decades, efforts have been made to seek metal alloys with high strength and low modulus at the same time, but an alloy showing both steel-like high yield strength (σy>1 GPa) and polymer-like low modulus (E≈10 GPa) still remains unattainable. So far, several alloys based on shape memory alloys (SMAs) have been reported to have a 1 GPa-class high strength and a moderately low modulus E of about 30 GPa, and a conventional Mg—Sc strain glass alloy has recently been shown to have a lower modulus E of about 20 GPa but with a lower strength σy of about 0.3 GPa. Despite these efforts, existing alloys still fall into the conventional grey band of FIG. 1a, and the desired polymer-like ultrahigh-strength property (comparable to the red region in the upper right corner of FIG. 1a, where a plurality of five-pointed stars are illustrated) still remains unattainable.

[0006] Therefore, there is an urgent need to research and develop a novel polymer-like ultrahigh-strength metal alloy with a polymer-like ultralow Young's modulus.SUMMARY

[0007] In order to achieve the above object, the present disclosure provides the following technical solutions:

[0008] provided is a polymer-like ultrahigh-strength metal alloy, wherein,

[0009] the alloy has the following levels of strength and Young's modulus over a wide temperature range of −80° C. to 80° C.:

[0010] a ultra-high yield strength σy of greater than or equal to 1.3 GPa, and

[0011] a Young's modulus E of about 10-16 GPa.

[0012] Preferably,

[0013] the wide temperature range is selected to be in: an aviation ambient temperature range of −80° C. to 70° C.

[0014] Preferably,

[0015] the alloy includes a strain glass matrix embedded with a small number of R and B19′ martensite seeds, i.e., martensite dual seeds.

[0016] Preferably,

[0017] the ultra-high yield strength σy is about 1.8 GPa.

[0018] Preferably,

[0019] the ultra-high yield strength σy is about 1.3-1.8 GPa.

[0020] Preferably,

[0021] the R martensite seeds have a size of 50-100 nm, P=3.

[0022] Preferably,

[0023] the B19′ martensite seeds have a size of 50-100 nm, P=2.

[0024] Preferably,

[0025] the alloy has a dual-seed strain glass (DS-STG) microstructure.

[0026] Preferably,

[0027] a stress-induced nucleation-free transition is present between dual-seed strain glass (DS-STG) and R and B19′ martensites.

[0028] Preferably,

[0029] the ultra-high yield strength σy is about 1.8 GPa, and the Young's modulus E is about 16 GPa at −80° C.;

[0030] the ultra-high yield strength σy is about 1.8 GPa, and the Young's modulus E is about 10 GPa at 20° C.; and

[0031] the ultra-high yield strength σy is about 1.8 GPa, and the Young's modulus E is about 15 GPa at 80° C.

[0032] Provided is a polymer-like ultrahigh-strength metal alloy, wherein,

[0033] the alloy has the following levels of strength and Young's modulus over a wide temperature range of −80° C. to 80° C.:

[0034] a ultra-high yield strength σy of greater than or equal to 1 GPa, and

[0035] a Young's modulus E of about 7-20 GPa.

[0036] Preferably,

[0037] the ultra-high yield strength σy is about 1-1.8 GPa, and

[0038] the Young's modulus E is about 7-20 GPa.

[0039] Preferably,

[0040] the ultra-high yield strength σy is about 1.2-1.8 GPa.

[0041] Preferably,

[0042] the ultra-high yield strength σy is greater than or equal to 1.2 GPa; and

[0043] the Young's modulus E is about 7-9 GPa.

[0044] Preferably,

[0045] the ultra-high yield strength σy is about 1.8 GPa, and the Young's modulus E is about 11 GPa at 20° C.

[0046] Preferably,

[0047] the ultra-high yield strength σy is about 1.6 GPa, and the Young's modulus E is about 11 GPa at 20° C.

[0048] Preferably,

[0049] the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 10.5 GPa at 20° C.

[0050] Preferably,

[0051] the ultra-high yield strength σy is about 1.7-1.9 GPa, and the Young's modulus E is about 8-13 GPa at 20° C.

[0052] Preferably,

[0053] the ultra-high yield strength σy is about 1.2-1.3 GPa, and the Young's modulus E is about 9-13 GPa at 20° C.

[0054] Preferably,

[0055] the ultra-high yield strength σy is about 1.3-1.5 GPa, and the Young's modulus E is about 12-15 GPa at 80° C.

[0056] Preferably,

[0057] the ultra-high yield strength σy is about 1.3-1.5 GPa, and the Young's modulus E is about 9-11 GPa at 80° C.

[0058] Preferably,

[0059] the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 18 GPa at 20° C.;

[0060] the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 9 GPa at 40° C.;

[0061] the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 9 GPa at 60° C.;

[0062] the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 10 GPa at 80° C.;

[0063] the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 13 GPa at 100° C.; and

[0064] the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 15 GPa at 120° C.

[0065] Preferably,

[0066] the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 7-10 GPa at 0° C.

[0067] In addition, the present disclosure also discloses a method for manufacturing any one of the alloys described above, including the steps of:

[0068] forming deformation-stabilized martensites;

[0069] forming a dual-crossover strain glass (DC-STG) state; and

[0070] forming a dual-seed strain glass (DS-STG) state to obtain the alloy.

[0071] Preferably,

[0072] room temperature cold working is adopted when forming the deformation-stabilized martensites or forming the dual-seed strain glass (DS-STG) state; and

[0073] heat treatment is adopted when forming the dual-crossover strain glass (DC-STG) state.

[0074] Preferably,

[0075] the alloy has the following levels of strength and Young's modulus over a wide temperature range of 20° C. to 80° C.:

[0076] the ultra-high yield strength σy is about 1.2 GPa, and the Young's modulus E is about 17-42 GPa at 20-80° C.; and

[0077] the alloy is an off dual-seed strain glass (off DS-STG) alloy.

[0078] The present disclosure has the following advantages:

[0079] The polymer-like ultrahigh-strength metal alloy is obtained in the present disclosure, and exhibits an unprecedented combination of a polymer-like ultra-low elastic modulus (e.g., about 10.5-20 GPa) and a steel-like ultra-high yield strength (e.g., about 1.3-1.8 GPa), and a large J-shaped elastic strain of about 7-8%. This unconventional alloy overcomes the long-standing trade-off between strength and flexibility. This polymer-like ultrahigh-strength alloy may open the door for a wide range of applications in emerging technologies such as morphing aerospace vehicles, superman-type humanoid robots and advanced biomedical devices. In addition, the manufacturing method disclosed in the present disclosure essentially involves a simple three-step thermomechanical processing route, which means that the industrial-grade scalable manufacturing route can achieve large-scale production of such alloys. In addition to this, among the existing flexible materials such as shape memory alloys and engineering polymers, the polymer-like ultrahigh-strength metal alloy has good temperature stability and excellent high-strain fatigue resistance.BRIEF DESCRIPTION OF FIGURES

[0080] FIG. 1 is a schematic diagram comparing the present disclosure with the prior art, wherein,

[0081] a is a schematic diagram comparing the present disclosure with those in the prior art in terms of strength and flexibility, wherein an abscissa represents flexibility and an ordinate represents strength,

[0082] b is a schematic diagram comparing the present disclosure with the prior art in terms of tensile properties, wherein an abscissa represents strain and an ordinate represents stress,

[0083] c is a schematic diagram comparing the present disclosure with the prior art in terms of a flexibility figure of merit σy / E, and

[0084] d is a schematic diagram of the present disclosure in terms of ultra-high strength and ultra-low Young's modulus;

[0085] FIG. 2 is a schematic diagram of a manufacturing method disclosed in the present disclosure and a microstructure of an alloy obtained by the manufacturing method; wherein,

[0086] a is a schematic diagram of a process flow of the manufacturing method disclosed in the present disclosure;

[0087] b is a schematic diagram of a phase transition behavior involved in Step 2, i.e., formation of a dual-crossover strain glass (DC-STG) state; and

[0088] c is a schematic diagram of the microstructure involved in Step 3, i.e., formation of a dual-seed strain glass (DS-STG) state;

[0089] FIG. 3 is a schematic diagram of a nucleation-free reversible transition of the obtained alloy of the present disclosure;

[0090] FIG. 4 is a schematic diagram showing a relationship between the annealing temperature in Step 2 in the manufacturing method disclosed in the present disclosure and the microstructure and properties, wherein

[0091] a is a schematic diagram showing a relationship between the annealing temperature in Step 2 and Young's modulus and recoverable strain; and

[0092] b is a schematic diagram showing a relationship between the annealing temperature in Step 2 and a phase transition behavior;

[0093] FIG. 5 is a schematic diagram showing the mechanical properties and fatigue resistance of the obtained alloy of the present disclosure over a wide temperature range, wherein,

[0094] a is a schematic diagram of characterization of strength and Young's modulus and recoverable strain; and

[0095] b is a schematic diagram of characterization of high-strain fatigue resistance at room temperature;

[0096] FIG. 6 is a schematic diagram showing the microstructure and mechanical properties of a Ti-50.8Ni alloy in an original solid solution state at room temperature according to the present disclosure, wherein

[0097] a is a schematic diagram of the microstructure, and

[0098] b shows a tensile stress-strain curve;

[0099] FIG. 7 is a schematic diagram showing the microstructure and mechanical properties of a Ti-50.8Ni alloy formed after processing in Step 1 of the present disclosure, wherein,

[0100] A is a microstructure schematic,

[0101] B is the tensile stress strain curve;

[0102] FIG. 8 is a schematic diagram of a tensile stress-strain curve of a dual-crossover strain glass (DC-STG) alloy formed after processing in Step 2 according to the present disclosure;

[0103] FIG. 9 is a schematic diagram of noise reduction processing of a high-resolution atomic image of a dual-seed strain glass alloy DS-STG according to the present disclosure;

[0104] FIG. 10 is a schematic diagram showing the annealing temperature in Step 2 versus a phase transition behavior of the alloy of the present disclosure, wherein,

[0105] a is a schematic diagram of an annealing temperature versus heat flow-temperature curve,

[0106] b is a schematic diagram of an annealing temperature versus storage modulus-temperature curve, and

[0107] c is a schematic diagram of an annealing temperature versus normalized resistivity-temperature curve;

[0108] FIG. 11 is a schematic diagram showing the annealing temperature in Step 2 versus an in situ cooling X-ray diffraction spectrum of the alloy according to the present disclosure, wherein,

[0109] a is a schematic diagram of an in situ cooling X-ray diffraction spectrum of mono-crossover strain glass (MC-STG) formed at an annealing temperature of 473 K in Step 2;

[0110] b is a schematic diagram of an in situ cooling X-ray diffraction spectrum of dual-crossover strain glass (DC-STG) formed at an annealing temperature of 573 K in Step 2; and

[0111] c is a schematic diagram of an in situ cooling X-ray diffraction spectrum of normal R / B19′ martensite formed at an annealing temperature of 773 K in Step 2;

[0112] FIGS. 12 to 16, 17 to 19, and 20 are schematic diagrams of tensile stress-strain curves of dual-seed strain glass alloys, i.e., DS-STG alloys manufactured in various embodiments of the present disclosure;

[0113] FIG. 21 is a schematic diagram of a tensile stress-strain curve of an Off DS-STG alloy manufactured by cold rolling according to the present disclosure;

[0114] FIG. 22 is a schematic diagram of a microstructure of the Off DS-STG alloy according to the present disclosure;

[0115] FIG. 23 is a schematic diagram of a quasi-linear elastic tensile behavior with room temperature high flexibility, high strength, and large recoverable strain of the Off DS-STG alloy according to the present disclosure; and

[0116] FIG. 24 is a schematic diagram showing the change of the superelasticity and the Young's modulus during loading of the Off DS-STG alloy as a function of strain in the present disclosure.DETAILED DESCRIPTION OF THE INVENTION

[0117] In one embodiment, the present disclosure discloses a method for manufacturing a polymer-like ultrahigh-strength metal alloy, including the steps of:

[0118] forming deformation-stabilized martensites;

[0119] forming a dual-crossover strain glass (DC-STG) state; and

[0120] forming a dual-seed strain glass (DS-STG) state to obtain the alloy.

[0121] In another embodiment,

[0122] room temperature cold working is adopted when forming the deformation-stabilized martensites or forming the dual-seed strain glass (DS-STG) state; and

[0123] heat treatment is adopted when forming the dual-crossover strain glass (DC-STG) state.

[0124] For example, a three-step thermomechanical manufacturing route of the alloy of the present disclosure, i.e., a dual-seed strain glass (DS-STG) alloy and analysis of the corresponding microstructural features after each step are achieved by the following description:

[0125] as shown in FIG. 2a, the DS-STG alloy is manufactured by a three-step thermomechanical processing route from a starting solution-treated Ti-50.8 at. % Ni alloy. This processing method is very simple and scalable to industrial production lines. The corresponding microstructure after each step is described below.

[0126] Step 1, processing: formation of deformation-stabilized martensites

[0127] The starting solution-treated Ti-50.8 at. % Ni alloy (FIG. 2a(i)) is in a B2 parent state (strictly speaking an unfrozen strain glass state (FIG. 6a)) at room temperature because its subzero martensitic transformation temperature Ms is about 250 K. Such B2 state exhibits a very high elastic modulus E of about 71 GPa, and incomplete super-elasticity, with a large residual strain and hysteresis during loading / unloading even at a moderate stress of about 0.3 GPa (FIG. 6b). This is a common behavior of Ti—Ni alloys when dislocation slip occurs during a stress-induced martensitic transformation. Thus, the starting B2 alloy is neither a strong alloy nor a polymer-like alloy.

[0128] During processing in Step 1, the B2 alloy undergoes severe deformation with a tensile elongation of 50% (FIG. 2(ii)). This cold working process results in a deformation-stabilized B19′ martensite, which remains unchanged at room temperature (FIG. 7a). As shown in FIG. 7b, this cold working strengthened B19′ martensite has a quasi-linear elastic behavior with a yield stress of up to about 1.3 GPa and an elastic modulus E of about 37 GPa. Such mechanical properties are consistent with previous reports on tensile-deformed titanium-nickel alloys; thus, the deformation-stabilized B19′ martensite alloy is a high-strength alloy, but not an ultralow-modulus alloy.

[0129] Step 2, processing: formation of a dual-crossover strain glass (DC-STG) state

[0130] During processing in Step 2, the deformation-stabilized B19′ martensite alloy is annealed at 573 K for 10 min (FIG. 2(iii)). Annealing at such a temperature will not only completely annihilate the transferable B19′ martensite produced during processing in Step 1, but will also produce a unique “dual-crossover strain glass” (DC-STG) state. The DC-STG is an unfrozen R strain glass (R-STG) at room temperature and undergoes a strain glass transition around Tg≈251 K, its characteristics being the same as those of an ordinary strain glass transition, including a frequency (ω) dependent Tg following the Vogel-Fulcher law (FIG. 2b) and invariance of average B2 structure throughout Tg, and appearance of nanoscale R strain domains. But unlike conventional R strain glass that maintains R nanostrain domains at 0 K or below, the DC-STG is able to transition to R and B19′ dual martensites at lower temperatures, as evidenced by in situ transmission electron microscopy (TEM) and X-ray diffractometry (XRD) (FIG. 2b). It is worth noting that the crossover transitions are incomplete and a significant amount of strain glass is still present even at the lowest temperature tested. This unique R strain glass can cross over into R and B19′ dual martensites and is thus named “dual-crossover strain glass” or DC-STG.

[0131] It is worth noting that this DC-STG state, albeit a forerunner state to the ultralow-modulus DS-STG state, is not a low modulus state by itself. It shows a nonlinear and hysteretic super-elastic behavior, with a medium initial elastic modulus of about 13 GPa and a high yield strength of about 1 GPa (FIG. 8).

[0132] Step 3, processing: formation of a dual-seed strain glass (DS-STG) state

[0133] During processing in Step 3, the DC-STG alloy undergoes a moderate tensile elongation of about 12% (FIG. 2a(iv)). As shown in an HRTEM image (FIG. 2c, top), this process introduces a small amount of R and B19′ martensite “dual seeds” (50-100 nm in size) into the room-temperature DC-STG matrix (manifested as nanosized R domains). The resulting strain glass state containing R and B19′ dual seeds is thus named “dual-seed strain glass” (DS-STG). Clearly, the realization of this unique DS-STG state is due to its forerunner state DC-STG being able to crossover into both R and B19′ martensites, as shown in FIG. 2b. This unique DS-STG state gives this alloy a polymer-like ultra-low modulus E of about 10.5 GPa as well as a J-shaped ultra-large pseudo-elastic strain εre of about 8% and a small hysteresis of about 15% (FIG. 1b). The mechanism of this unconventional elastic behavior will be revealed in FIG. 3.

[0134] A low-magnification image of the DS-STG alloy (FIG. 2c, bottom) shows that the DS-STG alloy contains high-density dislocations and a large amount of B2 mechanical twins, which are formed by severe cold working in the three-step process. These deformation hardening characteristics are characteristics of severely deformed Ti—Ni alloys and are also responsible for the ultra-high yield strength of the DS-STG alloy (FIG. 1b).

[0135] To confirm the general applicability of the three-step thermomechanical processing process in obtaining polymer-like ultrahigh-strength properties, we also tested a wire sample of the Ti-50.8 at. % Ni alloy by employing the same three-step process, but the severe deformation of the sample was done with a wire drawing machine rather than a tensile machine. As can be seen from FIG. 1b, the wire sample shows the same ultra-low modulus E of about 10.5 GPa as that of a plate sample, but the yield strength σy is even as high as about 1.8 GPa due to the advantageous effect of cold drawing. Thus, the current three-step processing route is suitable for wire and plate samples, with deformation in Step 1 being tensile elongation and cold drawing.Origin of the Polymer-Like Ultra-High Strength of the DS-STG Alloy

[0136] The in situ tensile XRD experiment (FIG. 3) shows that the DS-STG alloy with the above microstructural feature of the “dual seeds in the STG matrix” (FIG. 2c) undergoes a unique nucleation-free reversible transition between STG and R and B19′ martensites in a stress loading / unloading cycle up to 1.3 GPa, and this process leads to the observed ultra-low elastic modulus. Before loading (i.e., σ=0 GPa), the DS-STG, as a DC-STG matrix embedded with a small fraction of R and B19′ martensite seeds, shows one broadened B2 peak in its XRD profile (representing the average structure 38 of random STG nanodomains) with a long tail covering the R and B19′ positions. With increasing stress, the R and B19′ martensite peaks grow steadily without a critical stress, and the entire sample is almost completely transformed into B19′ martensite at σ=1.3 GPa. Apparently, this is a result of the R and B19′ martensite seeds bypassing the nucleation barrier of the stress-induced transition, resulting in the polymer-like ultra-low elastic moduli (E≈10.5 GPa). Upon unloading from 1.3 GPa, a smooth reverse transition occurs from B19′ to a mixture of R and B19′, and then to the original state of DS-STG. Thus, the original state of DS-STG is recovered after unloading. The large pseudo-elastic strain (εre≈8%) is a natural result of this reversible transformation because the maximum transformation strain to form the B19′ martensite in titanium-nickel alloys is about 8-9%. The ultra-low elastic modulus originates from the unique DS-STG state that enables a nucleation-free reversible transition between strain glass and R and B19′ martensites, and the ultra-high strength results from the deformation-induced strengthening effect.

[0137] Another important consequence of the nucleation-free DS-STG to R and B19′ transition is the J-shape and narrow hysteresis feature in a stress-strain curve of the DS-STG alloy, since R and B19′ martensite seeds can bypass the nucleation barrier during the stress-induced transition. This results in an ultra-low initial slope of the stress-strain curve, i.e., the “J-shape”, with narrow hysteresis. The narrow hysteresis is also reflected by the similarity of the XRD profiles between loading and unloading at the same stress level. The smooth, “J”-shaped and narrow-hysteretic behavior of DS-STG is in sharp contrast to that of the normal stress-induced transformation of STG to martensite (FIG. 8) or the traditional stress-induced martensitic transformation (FIG. 6b), in which the inevitable nucleation results not only in a higher elastic modulus E of about 30-70 GPa, but also in a strongly non-linear and hysteretic stress-strain behavior with a stress plateau.

[0138] FIG. 4a shows that the DS-STG state is critical for obtaining the polymer-like ultra-low modulus. When the annealing temperature Ta in Step 2 deviates from the optimum temperature of 573 K, the alloy deviates from the DS-STG state, resulting in an increase in elastic modulus and a decrease in recoverable strain.

[0139] FIG. 4b explains why the DS-STG state will only occur when the annealing temperature Ta in Step 1 is around 573 K, this is because the forerunner state DC-STG of DS-STG will only occur around Ta≈573 K, as shown in a phase diagram of FIG. 4b (determined from data of FIG. 2b, and FIGS. 10 and 11). When Ta deviates significantly from 573 K, the forerunner DC-STG will not be present anymore, so the R+B19′ dual-seed strain glass cannot be achieved after the processing in Step 3 (i.e., the elongation of 12%).

[0140] The phase diagram (FIG. 4b) shows that a significant decrease in the Ta value (e.g., Ta≈473 K) will result in mono-crossover strain glass (MC-STG) after the processing in Step 2, i.e., R strain glass that can cross (partially) over into one martensite B19′ at a low temperature. After the processing in Step 3 (the elongation of 12%), the mono-crossover strain glass will be converted to mono-seed strain glass (MS-STG, i.e., with B19′ seeds only) with a higher modulus of about 17 GPa (about 1.5 times that of DS-STG) and a lower recoverable strain of about 6%.

[0141] At a significantly higher Ta≈773 K (a common annealing temperature for titanium-nickel alloys), the strain glass transition no longer exists, and the alloy is characterized by a normal two-step martensitic transformation, i.e., transformation to R and B19′ martensites. After the processing in Step 3 (the elongation of 12%), the alloy will be transformed to aligned B19′ martensite with a higher modulus of about 30 GPa and a small recoverable strain of about 4%.

[0142] The elastic modulus of the DS-STG state is lower compared with the non-DS-STG state because of the phase instability of the system in the vicinity of the three different phases. This results in a nearly flat free energy landscape, which facilitates the transition from one phase to the other phase, making a lattice the softest.

[0143] In another more specific embodiment, the manufacturing method is described in detail by:Sample Preparation

[0144] Commercial Ti-50.8 at. % Ni alloys from SaiTe Inc. Ltd. were used for manufacturing plate and wire samples. All samples were solution treated in a homogeneous B2 state, followed by quenching in water. The plate samples were sealed in a vacuum quartz tube, and solution treated at a temperature of 1273 K for 90 min; the wire samples were solution treated in an automatic furnace at 1023 K, with a heating duration of 10 min, with moving wires entering and exiting the automatic furnace before quenching. After the solution treatment, the oxidation-affected surface layers of all samples were removed by a chemical etchant.

[0145] Both the plate and wire samples underwent a similar three-step process as described below, as shown in FIG. 2a, with a difference only in the deformation mode in Step 1. During the processing in Step 1, the plate samples were severely deformed by 50% cold elongation and the wire samples were severely deformed by 50% area reduction of cold elongation. The severe deformation in Step 1 transforms the starting B2 phase into a transferable martensite microstructure stabilized by high-density dislocations.

[0146] During the processing in Step 2, the severely deformed samples were moderately annealed at 573 K for 10 min, followed by quenching in water. This process annihilates the transferable martensite formed during the processing in Step 1 and produces a unique “dual-crossover strain glass” (DC-STG) state as described in the text. It is worth noting that the annealing temperature of 573 K is very important for the ultra-low modulus after the processing in Step 3, and this annealing temperature is lower than the conventional annealing temperature for titanium-nickel alloys (typically between 673 K and 773 K).

[0147] During the processing in Step 3, both the plate and wire samples processed in Step 2 undergo moderate deformation at room temperature with an elongation of 12%, slightly above the transformation strain of B19′ (about 8-9%). This step produces a small fraction of R and B19′ martensite “dual seeds” in the DC-STG matrix, and such a state is referred to as “dual-seed strain glass” (DS-STG).Mechanical Property Characterization

[0148] The stress-strain curves of all the samples were measured at room temperature with a Shimadzu AG-IS tensile machine at a strain rate of 2×10−3 s−1. An extensometer (Epsilon Tech.) was used to measure the strain during the tension. Young's modulus (E) was calculated from the initial slope of the stress-strain curve. Yield stress (σy) and critical stress (σc) to induce martensitic transformation were determined by the conventional double tangent method. Recoverable strain (εre) was determined from the difference between the maximum strain and the minimum strain shown in an unloading curve.

[0149] The stress-strain curves of typical structural materials (shown in FIG. 1b) were also tested by using the same tensile machine for comparison with our DS-STG alloy. These structural materials include metal alloys and organic-based materials: spring steel (65Mn), titanium alloy (Ti-6A1-4V ELI), aluminum alloy (7021-T62), magnesium alloy (AZ91D), bamboo (dry), fiber reinforced polymers (FRP, PPS GF40), polyether ether ketone (PEEK) and polyphenylene sulfide (PPS).In Situ Structural and Microstructural Characterization

[0150] Crystal structure characterization at different temperatures was carried out by an X-ray diffractometer (Shimadzu XRD-7000) equipped with a temperature sample holder. Cu-Kα radiation was used as an incident X-ray source. In situ X-ray diffraction measurements under tensile deformation were performed with a self-made tension holder that enabled in situ tension of a small XRD sample with a size of 40×2×0.05 mm3.

[0151] Microscopic characterization was carried out by using a JEM 2100F TEM equipped with a field emission gun and a cooling holder. Data was recorded by using a GATAN CCD slow scan camera and analyzed by DigitalMicrograph software. TEM samples were mechanically thinned to about 0.1 mm, followed by twin-jet polishing with a solution of 90% methanol and 10% sulfuric acid.

[0152] Characterization of martensitic transformation and strain glass transition by using a differential scanning calorimeter (DSC), a dynamical mechanical analyzer (DMA) and a resistivity

[0153] The occurrence of the martensitic transformation was monitored by monitoring a latent heat peak using a differential scanning calorimeter (DSC-Q200) produced by TA Instruments, with a heating / cooling rate of 10 K / min. Since the strain glass transition does not produce a significant differential scanning calorimeter peak, a dynamic mechanical analyzer (DMA-Q850) from TA Instruments was used to monitor a key feature of the strain glass transition: a Vogel-Fulcher-type frequency-dependent anomaly in a storage modulus curve. The measurement was done in a single cantilever mode with an amplitude of 15 μm at a frequency of 1, 2, 5, 10, 20 Hz. The cooling / heating rate was 1 K / min. The R and B19′ martensitic transformation and strain glass transition were monitored by using a four-probe resistivity measurement at a constant current of 100 mA and a cooling / heating rate of 1 K / min.Unique Mechanical Properties of the DS-STG Alloy

[0154] FIG. 1a shows that the DS-STG alloy overcomes the well-established trade-off relationship between strength and flexibility in existing structural materials (as shown by a gray band in FIG. 1a), and exhibits an unprecedented combination of ultrahigh-strength-steel grade yield strength σy of about 1.8 GPa and polymer-like ultra-low Young's modulus E of about 10.5 GPa (corresponding to ultra-high elastic compliance S=1 / E≈0.1 GPa−1). This combination of properties brings the DS-STG alloy into a “strong yet flexible” state (a red state) heretofore difficult to attempt in FIG. 1a.

[0155] FIG. 1b shows that the stress-strain curve of the DS-STG alloy is in the “J” shape with an excessively large pseudo-elastic strain εre of about 8% before reaching the elasticity limit σy of about 1.8 GPa for the wire sample and σy of about 1.3 GPa for the plate sample. The 1.8 GPa yield strength of the DS-STG alloy even exceeds that of most ultrahigh-strength steels (such as quenched spring steels), and the pseudo-elastic strain value of 8% not only far exceeds the small elastic strain εre of about 0.2-1% of typical metal alloys, but is also greater than the pseudo-elastic strain (εre of about 1-5%) of many organic structural materials (such as wood, bamboo, bone, polymers and fiber reinforced polymers). In distinct contrast to the superelasticity of conventional shape memory alloys, which are characterized by a stress plateau with a larger initial modulus (E≈30-75 GPa) and a huge hysteresis (about 50-90%). It is well known that the strong hysteresis behavior of shape memory alloys is undesirable in many applications.

[0156] As shown in FIG. 1c, the combination of ultra-high strength and ultra-low modulus provides the DS-STG alloy with ultra-high flexibility σy / E of about 0.17, far exceeding that of the existing structural materials including typical metal alloys and organic materials. The flexibility figure of merit σy / E is a measure of the ultimate flexibility of the material before yielding failure, whereas σy / E of the existing structural materials is typically much lower, and is about 0.01-0.06.

[0157] FIG. 1d show visual evidence for the unconventional “strong yet flexible” (i.e., polymer-like ultra-high strength) properties of the DS-STG alloy, which is in sharp contrast to the conventional behavior of two reference materials: a “strong and stiff” spring steel (with a high modulus E of about 190 GPa and a high strength σy of about 1.36 GPa) and a “weak and flexible” fiber reinforced plastic (with a low modulus E of about 11 GPa and a low strength σy of about 0.13 GPa). Thus, the DS-STG alloy combines the high strength of steels and the high flexibility of a polymer and behaves like a super-strong organic material, which has long been desired by many emerging technologies.

[0158] A further detailed explanation of FIGS. 1 to 11 is made below:

[0159] Referring to FIG. 1, the DS-STG alloy has polymer-like ultra-high strength compared with typical metal alloys and organic materials. a, the DS-STG alloy exhibits an unprecedented combination of steel-like high strength σy of about 1.3-1.8 GPa and polymer-like ultra-low Young's modulus E of about 10.5 GPa (or ultra-high elastic compliance S of about 0.1 GPa−1). This “strong yet flexible” characteristic overcomes the trade-off between strength and flexibility (the gray band, i.e., an Ashby plot), which makes the existing structural materials either “strong and stiff” or “weak and flexible”. NICSMA, PEEK and PPS are abbreviations for in situ composites of nanowires and shape memory alloys, polyether ether ketone and polyphenylene sulfide, respectively. b, J-shape pseudo-elastic stress-strain curves of the DS-STG alloy compared with stress-strain curves of typical metal alloys (quenched 65Mn spring steel, Ti-6A1-4V ELI titanium alloy, 7021-T62 aluminum alloy and AZ91D magnesium alloy) and organic materials (bamboo, fiber reinforced plastic, PEEK and PPS) within their elastic limits. The DS-STG alloy achieved an ultra-large pseudo-elastic strain εre of about 8%. c, The flexibility figure of merit σy / E of the DS-STG alloy far exceeds that of existing structural materials, including metal alloys and organic materials. Young's modulus is determined by the initial slope of the stress-strain curve. Visual evidence for the untraditional “strong yet flexible” property of the DS-STG alloy is in sharp contrast to that of two traditional materials, one being a “strong and stiff” spring steel and the other being a “weak and flexible” fiber reinforced polymer.

[0160] Referring to FIG. 2, a three-step thermomechanical processing route to achieve the polymer-like ultrahigh-strength DS-STG alloy and microstructure characteristics of the sample after each step are shown. a, the three-step thermomechanical processing route and the corresponding state of the alloy after each step. FIG. 6 shows the microstructure of the starting B2 parent phase (actually unfrozen strain glass), and FIG. 7 shows the microstructure of the deformation-stabilized B19′ martensite obtained after the processing in Step 1. b, the dual-crossover strain glass (DC-STG) obtained after the processing in Step 1 was revealed by DMA results combined with in situ TEM and XRD characterization. At room temperature, the DS-STG is unfrozen R strain glass that first undergoes a strain glass transition into frozen strain glass around Tg≈251 K, followed by crossover transitions into R and B19′ dual martensite at a low temperature. The strain glass transition is manifested by a frequency (w)-dependent Tg(w) versus Vogel-Fulcher (V-F) relationship, invariance of an average B2 structure (XRD), presence of nanoscale (5-20 nm) R domains and corresponding asymmetric and diffuse ⅓(110)B2 diffraction spots (labeled R-STG in a TEM dark field image (DFI) and a selected diffraction pattern). The dual-crossover transition from strain glass to R and B19′ martensites is manifested by the appearance of large R and B19′ domains (50-100 nm) in the DFI, sharp ⅓(110)B2 spots (labeled R) and sharp ½(110)B2 spots (labeled B19′) appeared in a selected area diffraction (SAD) pattern, and R and B19′ martensite peaks also appeared in the XRD pattern at 123 K and 183 K. c, An HRTEM image (top) at room temperature of the DS-STG alloy obtained after the processing in Step 1, and the DS-STG alloy is characterized by R and B19′ martensite seeds embedded in the R-STG matrix. Noise reduction processing of the image is shown in FIG. 9. The enlarged images (middle) show lattice fringes of R strain glass nanodomains (with a size of 5 -15 nm and a lattice spacing 3d110B2, or P=3), R martensite seeds (with a size of 50-100 nm, P=3) and B19′ martensite seeds (with a size of 50-100 nm, P=2). A low-magnetic-field bright-field image (bottom) of DS-STG shows high-density dislocations and B2 mechanical twins, which are responsible for the ultrahigh-strength of the alloy.

[0161] Referring to FIG. 3, the polymer-like elasticity of the DS-STG alloy originates from the stress-induced nucleation-free transition between DS-STG and R and B19′ martensite, as evidenced by in-situ XRD during the stress loading / unloading cycle. XRD profiles during loading (left) from 0 GPa and 1.3 GPa and unloading (right) from 1.3 GPa and 0 GPa are shown. A smooth transition from the STG (containing R and B19′ seeds) to the R+B19′ mixture and then to B19′ martensite occurs during loading, while a smooth reverse transition occurs during unloading. The smooth evolution of the R and B19′ peaks indicates that the transition can begin at almost zero stress, and that no nucleation barrier exist during the transition. The top inset shows X-ray goniometry under tensile loading.

[0162] Referring to FIG. 4, the polymer-like ultra-low modulus and high recoverable strain are consistent with those of the DS-STG state which arises from the forerunner state DC-STG formed upon annealing in Step 2 at Ta≈573 K. a, The polymer-like low modulus (E≈10.5 GPa) and maximum recoverable strain of 8% are consistent with those of the DS-STG state that is achieved by the processing in Step 3 (the elongation of 12%) of the forerunner DC-STG sample. b, a phase diagram of severely deformed Ti-50.8Ni alloy as a function of the annealing temperature Ta in Step 2, where the DC-STG state is only achieved around Ta≈573 K, and low temperature annealing (e.g., Ta=473 K) will produce mono-crossover strain glass (MC-STG) that is transformed into mono-seed strain glass (MS-STG) after the processing in Step 3; high temperature annealing (e.g., Ta=673-773 K) will produce normal R / B19′ martensitic transformation, and the alloy is transformed into aligned B19′ martensite after the processing in Step 3. In the phase diagram, Tnd is the nanodomain starting temperature, Rs and Ms are the martensitic transformation starting temperature for R and B19′ martensite, respectively, Tg is the STG transition temperature, and Ts is the spontaneous STG-R transition starting temperature. These temperatures are determined by experiments in FIGS. 10 and 11.

[0163] It should be noted that with further reference to FIG. 5, the DS-STG alloy has polymer-like ultra-high strength over a wide temperature range from 193 K (−80° C.) to 353 K (80° C.), which covers the ambient temperature range around a morphing aircraft. b. DS-STG has excellent high-strain fatigue resistance compared with traditional flexible materials (including Ti—Ni shape memory alloys and engineering polymers such as PEEK and fiber reinforced plastic), which are candidate materials for morphing aircrafts and artificial muscle fibers. The upper left inset shows the fatigue behavior characteristics of DS-STG in terms of strain-life curves compared with existing metal alloys (including Ti—Ni shape memory alloys, and traditional metal alloys such as steel, titanium alloys, aluminum alloys, and magnesium alloys) and engineering polymers. The upper right inset shows that the DS-STG alloy does not fatigue even after 5 million cycles under a considerable elastic strain of 1.6-1.8%. High-strain fatigue refers to a fatigue test with a maximum cyclic strain em>1.5%; the corresponding maximum stress is sm. The alloy can maintain this excellent performance over a wide temperature range from 193 K (−80° C.) to 353 K (80° C.) and exhibits excellent high-strain fatigue resistance among the existing flexible materials such as shape memory alloys and engineering polymers, making it a viable candidate material for morphing aircrafts and artificial muscles. In situ X-ray diffractometry has shown that DS-STG enables nucleation-free reversible transitions between the strain glass and R and B19′ martensites during stress loading / unloading, resulting in ultra-low modulus and large recoverable strain over a wide temperature domain, with a narrow hysteresis. The polymer-like ultrahigh-strength behavior can persist over a wide ambient temperature range from 193 K to 353 K (FIG. 5a) and therefore can be applied in the aerospace field, and this alloy has excellent fatigue resistance compared with other artificial muscle fiber candidate materials (FIG. 5b). The DS-STG alloy has polymer-like ultra-high strength over a wide temperature range from 193 K (−80° C.) to 353 K (80° C.) (FIG. 5a), which covers the aviation ambient temperature range (−80° C. to 70° C.) that the morphing aircraft is expected to experience. Over this wide temperature range, the Young's modulus is maintained at a polymer-like ultra-low level of 10-16 GPa, which is similar to its room temperature value of 10.5 GPa; the yield strength is maintained at a ultrahigh-strength level of 1.8 GPa. The recoverable strain is also maintained at a high value of 6.3% or more over this wide temperature range. The DS-STG alloy has excellent high-strain fatigue resistance (FIG. 5b) compared with other flexible materials (such as traditional Ti—Ni shape memory alloys and engineering polymers such as polyether ether ketone and fiber reinforced plastic), which are candidate materials for morphing aircrafts and artificial muscle fibers. High-strain fatigue resistance is critical for morphing aircraft and artificial muscle applications, as high-strain cyclic deformation is common in these applications. The upper left inset of FIG. 5b shows that the DS-STG alloy has the best high-strain fatigue resistance not only among Ti—Ni shape memory alloys, and conventional metal alloys such as steels, titanium alloys, and aluminum alloys, but also among engineering polymers such as PEEK and fiber reinforced plastic. The upper right inset of FIG. 5b shows that the DS-STG alloy exhibits remarkable fatigue-free characteristics (with a fatigue life of more than 5×106 cycles) at a larger strain amplitude of 1.8%, which is the highest in the existing metal alloys including shape memory alloys, and also exceeds that of engineering polymers. Such a fatigue-free strain amplitude level may meet the requirement of morphing wings to produce a sufficiently large camber shape change (>6%). The temperature dependence of the mechanical properties was characterized by the same tensile machine (Shimadzu AG-IS) equipped with a temperature chamber (Shimadzu TCLN-220P). An extensometer (Epsilon Tech. 3442-006M-050M-LHT) was used to ensure accurate strain measurements. The stress-strain curve is the result of the second cycle of the 1.8 GPa tensile test. The fatigue life (N)-maximum cyclic stress sm curve was measured by using a fatigue tester (TA 3330 Series III) at room temperature in a load-controlled tension-tension mode between 0 and the maximum stress sm. The test frequency was 0.5-10 Hz. The corresponding maximum cyclic strain em was obtained from the experimental stress-strain relationship at 20° C. (an inset in FIG. 5a), except for the fatigue-free samples (5×106 cycles), for which the strain em was measured directly by an extensometer (as shown in a right inset in FIG. 5b).

[0164] Referring further to FIG. 6, which illustrates the microstructure and mechanical properties of the starting solution-treated Ti-50.8Ni alloy at room temperature, wherein a. the alloy exhibits a B2 average structure but contains nanoscale R domains manifested as ⅓(110)B2 superlattices in the diffraction pattern and lattice fringes with a spacing of 3d110B2 or an atomic arrangement period P=3 in HRTEM and noise-filtered inverse fast Fourier transform (IFFT) images; and b. imperfect superelasticity with low yield strength (about 0.32 GPa) and high Young's modulus (about 71 GPa) associated with the starting solution-treated Ti-50.8Ni alloy which has a martensitic transformation temperature Ms of about 250 K (see FIG. 10).

[0165] Referring further to FIG. 7, which illustrates the microstructure and mechanical properties of a severely deformed Ti-50.8Ni alloy (i.e., after step-1 processing), wherein a. the Ti-50.8Ni alloy after step-1 processing exhibits high density dislocations and deformation-stabilized B19′ martensitic laths at room temperature; and b. the severely deformed Ti-50.8Ni alloy shows a quasi-linear elastic behavior with a high strength of about 1.3 GPa and a moderate elastic modulus of about 37 GPa.

[0166] Referring further to FIG. 8, which illustrates that a dual-crossover strain glass (DC-STG) sample (i.e., the sample after step 2-processing) exhibits superelasticity at room temperature, and has a moderate strength of about 1.0 GPa and a moderate initial modulus of about 32 GPa, as well as large non-linearity and hysteresis.

[0167] Referring further to FIG. 9, which illustrates a room-temperature high-resolution transmission electron microscopy (HRTEM) image (top left) and a noise-filtered inverse fast Fourier transform (IFFT) image (bottom right) of the DS-STG alloy, showing R and B19′ martensite seeds embedded in a R strain glass matrix. The corresponding diffraction patterns show sharp R and B19′ superlattice spots from the R and B19′ seeds. The IFFT image is generated by a fast Fourier transform (FFT) pattern (bottom left).

[0168] Referring further to FIG. 10, the transition temperatures of Tnd, Rs, Ms, Tg, and Ts in the phase diagram of FIG. 4b are determined by combining DSC (differential scanning calorimetry), resistivity, and DMA (dynamic mechanical analysis) measurements, where Tnd, Rs, Ms, Tg, and Ts represent the nanodomain starting temperature, the R martensitic transformation starting temperature, the B19′ martensitic transformation starting temperature, the strain glass transition temperature, and the spontaneous transition temperature from strain glass to R martensite, respectively. a1, b1, and c1 show that the starting solution-treated Ti-50.8Ni alloy undergoes transitions from the B2 parent phase to the unfrozen strain glass (at Tnd) and then to the B19′ martensite (at Ms). a2, b2, c2, a3, b3, and c3, the low annealing temperatures Ta=473 K and Ta=523 K in Step 2 lead to a mono-crossover strain glass (MC-STG) state, which first shows a strain glass transition into a frozen R strain glass, followed by a crossover transition to a single B19′ martensite. The strain glass transition is evidenced by a nearly vanishing DSC peak, and a Vogel-Fulcher type frequency-dependent elastic modulus dip at Tg of about 219 K in DMA curves; the mono-crossover transition to B19′ is evidenced by a decreasing resistivity below Ms and the existence of large transformation hysteresis characteristic, as well as the appearance of B19′ peaks at low temperature in XRD. a4, b4, and c4, the moderate annealing temperature Ta=573 K in Step 2 leads to a dual-crossover strain glass (DC-STG) state, which first shows a strain glass transition into a frozen R strain glass, followed by a dual-crossover transition to R martensite and then to B19′ martensite. The strain glass transition is evidenced by a Vogel-Fulcher type frequency-dependent elastic modulus dip at Tg of about 251 K in DMA curves; the crossover transition into R and B19′ martensites is evidenced by the appearance of R and B19′ peaks at low temperatures in XRD, which is a key measurement for a crossover strain glass transition. The second crossover transition to B19′ martensite is also evidenced by a decreasing resistivity below Ms and the existence of large transformation hysteresis. a5, b5, c5, a6, b6, c6, a7, b7, and c7, the high annealing temperatures Ta=623, 673, and 773 K in Step 2, which are the standard annealing temperatures for Ti—Ni alloys, lead to a familiar two-step martensitic transition, first to R martensite, and then to B19′ martensite. As evidenced by two prominent DSC peaks, and frequency-independent elastic modulus dip in DMA curves. The martensitic transition to R martensite is evidenced by the sharp increase in resistivity below Rs and the existence of about 5 K transition hysteresis characteristics of R transition. The martensitic transition to B19′ martensite is evidenced by the decrease in resistivity below Ms and the existence of large transition hysteresis, being consistent with the literatures 43 and 44. Direct evidence from XRD confirms this transition sequence.

[0169] Referring further to FIG. 11, which illustrates in situ XRD (X-ray diffraction) analysis of (a) a mono-crossover strain glass (MC-STG) sample (Ta=473 K), (b) a dual-crossover strain glass (DC-STG) sample (Ta=573 K), and (c) a conventional R / B19′ martensite sample (Ta=773 K) during cooling. a. The MC-STG sample maintains an average B2 structure below the strain glass transition temperature Tg, and transforms to single B19′ martensite below Ms, but the transition is incomplete. b. The DC-STG sample maintains an average B2 structure down to the strain glass transition temperature Tg, then transforms to a R phase first and then to B19′ martensite, but the transitions are also incomplete. c. The conventional R / B19′ martensite sample (Ta=773 K) undergoes a spontaneous transformation from unfrozen strain glass (FIG. 6, a R local structure with an average B2 structure) to the R phase and finally to B19′ martensite.

[0170] Referring to FIG. 12, in another embodiment, the present disclosure discloses a DS-STG alloy manufactured by cold drawing, the DS-STG alloy uses Ti-50.8Ni as a master alloy, and a manufacturing method includes the following steps:

[0171] Step 1: deformation-stabilized martensites are formed, specifically,

[0172] cold deformation: Ti-50.8Ni is drawn with a cross-sectional shrinkage of 50% at room temperature to manufacture the deformation-stabilized martensites;

[0173] Step 2: a dual-crossover strain glass (DC-STG) state is formed, specifically,

[0174] annealing: heat preservation is performed at 300° C. for 15 min to form dual-crossover strain glass; and

[0175] Step 3: a dual-seed strain glass (DS-STG) state is formed, specifically,

[0176] cold deformation: uniaxial stretching with an elongation of 12% is performed at room temperature to form dual-seed strain glass.

[0177] Referring further to FIG. 12, the DS-STG alloy manufactured by cold drawing described above has the following yield strength and Young's modulus:

[0178] the ultra-high yield strength σy is about 1.8 GPa and the Young's modulus E is about 11 GPa at 20° C.

[0179] Referring to FIG. 13, in another embodiment, the present disclosure discloses a DS-STG alloy manufactured by cold rolling, the DS-STG alloy uses Ti-50.8Ni as a master alloy, and a manufacturing method includes the following steps:

[0180] Step 1: dislocation-induced strain glass is formed, specifically,

[0181] cold deformation: Ti-50.8Ni is rolled with a thickness reduction of 50% at room temperature to manufacture the dislocation-induced strain glass;

[0182] Step 2: a dual-crossover strain glass (DC-STG) state is formed, specifically,

[0183] annealing: heat preservation is performed at 300° C. for 15 min to form dual-crossover strain glass; and

[0184] Step 3: a dual-seed strain glass (DS-STG) state is formed, specifically,

[0185] cold deformation: uniaxial stretching with an elongation of 12% is performed at room temperature to form dual-seed strain glass.

[0186] Referring further to FIG. 13, the DS-STG alloy manufactured by cold rolling described above has the following yield strength and Young's modulus:

[0187] the ultra-high yield strength σy is about 1.6 GPa and the Young's modulus E is about 11 GPa at 20° C.

[0188] Referring to FIG. 14, in another embodiment, the present disclosure discloses a DS-STG alloy manufactured by cold tension, the DS-STG alloy uses Ti-50.8Ni as a master alloy, and a manufacturing method includes the following steps:

[0189] Step 1: deformation-stabilized martensites are formed, specifically,

[0190] cold deformation: Ti-50.8Ni is stretched with an elongation of 50% at room temperature to manufacture the deformation-stabilized martensites;

[0191] Step 2: a dual-crossover strain glass (DC-STG) state is formed, specifically,

[0192] annealing: heat preservation is performed at 300° C. for 15 min to form dual-crossover strain glass; and

[0193] Step 3: a dual-seed strain glass (DS-STG) state is formed, specifically,

[0194] cold deformation: uniaxial stretching with an elongation of 12% is performed at room temperature to form dual-seed strain glass.

[0195] Referring further to FIG. 14, the DS-STG alloy manufactured by cold tension described above has the following yield strength and Young's modulus:

[0196] the ultra-high yield strength σy is about 1.3 GPa and the Young's modulus E is about 10.5 GPa at 20° C.

[0197] Referring further to FIG. 15, in further embodiments, the DS-STG alloy manufactured by cold drawing uses Ti-50.8Ni as a master alloy, and a manufacturing method includes the following steps:

[0198] Step 1: deformation-stabilized martensites are formed, specifically,

[0199] cold deformation: Ti-50.8Ni is drawn with a cross-sectional shrinkage of 50% at room temperature to manufacture the deformation-stabilized martensites;

[0200] Step 2: a dual-crossover strain glass (DC-STG) state is formed, specifically,

[0201] annealing: heat preservation is performed at 300° C. for 15 min to form dual-crossover strain glass; and

[0202] Step 3: a dual-seed strain glass (DS-STG) state is formed, specifically,

[0203] cold deformation: uniaxial stretching with an elongation of about 12% is performed at room temperature to form dual-seed strain glass.

[0204] Referring further to FIG. 15, the DS-STG alloy manufactured by cold drawing described above has the following yield strength and Young's modulus at 20° C.:Sample 1:a Young's modulus E of about 10 GPa

[0206] a yield strength σy of about 1.75 GPa;Sample 2:a Young's modulus E of about 10 GPa

[0208] a yield strength σy of about 1.76 GPa;Sample 3:a Young's modulus E of about 11 GPa

[0210] a yield strength σy of about 1.82 GPa;Sample 4:a Young's modulus E of about 10 GPa

[0212] a yield strength σy of about 1.76 GPa;Sample 5:a Young's modulus E of about 11 GPa

[0214] a yield strength σy of about 1.81 GPa;Sample 6:a Young's modulus E of about 8 GPa

[0216] a yield strength σy of about 1.77 GPa;Sample 7:a Young's modulus E of about 13 GPa

[0218] a yield strength σy of about 1.75 GPa;Sample 8:a Young's modulus E of about 11 GPa

[0220] a yield strength σy of about 1.82 GPa;

[0221] that is, the ultra-high yield strength σy is about 1.75-1.82 GPa and the Young's modulus E is about 8-13 GPa at 20° C.

[0222] Referring further to FIG. 16, in further embodiments, the DS-STG alloy manufactured by cold tension uses Ti-50.8Ni as a master alloy, and a manufacturing method includes the following steps:

[0223] Step 1: deformation-stabilized martensites are formed, specifically,

[0224] cold deformation: Ti-50.8Ni is stretched with an elongation of 50% at room temperature to manufacture the deformation-stabilized martensites;

[0225] Step 2: a dual-crossover strain glass (DC-STG) state is formed, specifically,

[0226] annealing: heat preservation is performed at 300° C. for 15 min to form dual-crossover strain glass; and

[0227] Step 3: a dual-seed strain glass (DS-STG) state is formed, specifically,

[0228] cold deformation: uniaxial stretching with an elongation of about 12% is performed at room temperature to form dual-seed strain glass.

[0229] Referring further to FIG. 16, the DS-STG alloy manufactured by cold tension described above has the following yield strength and Young's modulus at 20° C.:Sample 1:a Young's modulus E of about 9 GPa

[0231] a yield strength σy of about 1.28 GPa;Sample 2:a Young's modulus E of about 10 GPa

[0233] a yield strength σy of about 1.25 GPa;Sample 3:a Young's modulus E of about 10 GPa

[0235] a yield strength σy of about 1.27 GPa;Sample 4:a Young's modulus E of about 10 GPa

[0237] a yield strength σy of about 1.28 GPa;Sample 5:a Young's modulus E of about 9 GPa

[0239] a yield strength σy of about 1.24 GPa;Sample 6:a Young's modulus E of about 12 GPa

[0241] a yield strength σy of about 1.28 GPa;Sample 7:a Young's modulus E of about 12 GPa

[0243] a yield strength σy of about 1.27 GPa;Sample 8:a Young's modulus E of about 13 GPa

[0245] a yield strength σy of about 1.28 GPa;

[0246] that is, the ultra-high yield strength σy is about 1.2-1.3 GPa and the Young's modulus E is about 9-13 GPa at 20° C.

[0247] Referring to FIG. 17, in another embodiment, the present disclosure discloses a DS-STG alloy manufactured by cold rolling, the DS-STG alloy uses Ti-50Ni as a master alloy, and a manufacturing method includes the following steps:

[0248] Step 1: dislocation-induced strain glass is formed, specifically,

[0249] cold deformation: Ti-50Ni is rolled with a thickness reduction of 50% at room temperature to manufacture the dislocation-induced strain glass;

[0250] Step 2: a dual-crossover strain glass (DC-STG) state is formed, specifically,

[0251] annealing: heat preservation is performed at 300° C. for 15 min to form dual-crossover strain glass; and

[0252] Step 3: a dual-seed strain glass (DS-STG) state is formed, specifically,

[0253] cold deformation: uniaxial stretching with an elongation of 8% is performed at 80° C. to form dual-seed strain glass.

[0254] Referring further to FIG. 17, the DS-STG alloy manufactured by cold rolling described above has the following yield strength and Young's modulus at 80° C.:Sample 1:a Young's modulus E of about 15 GPa

[0256] a yield strength σy of about 1.3 GPa;Sample 2:a Young's modulus E of about 12 GPa

[0258] a yield strength σy of about 1.5 GPa;

[0259] that is, the ultra-high yield strength σy is about 1.3-1.5 GPa and the Young's modulus E is about 12-15 GPa at 80° C.

[0260] Referring to FIG. 18, in another embodiment, the present disclosure discloses a DS-STG alloy manufactured by cold rolling, the DS-STG alloy uses Ti-50Ni as a master alloy, and a manufacturing method includes the following steps:

[0261] Step 1: dislocation-induced strain glass is formed, specifically,

[0262] cold deformation: Ti-50Ni is rolled with a thickness reduction of 40% at room temperature to manufacture the dislocation-induced strain glass;

[0263] Step 2: a dual-crossover strain glass (DC-STG) state is formed, specifically,

[0264] annealing: heat preservation is performed at 300° C. for 15 min to form dual-crossover strain glass; and

[0265] Step 3: a dual-seed strain glass (DS-STG) state is formed, specifically,

[0266] cold deformation: uniaxial stretching with an elongation of 8% is performed at 80° C. to form dual-seed strain glass.

[0267] Referring further to FIG. 18, the DS-STG alloy manufactured by cold rolling described above has the following yield strength and Young's modulus at 80° C.:Sample 1:a Young's modulus E of about 11 GPa

[0269] a yield strength σy of about 1.3 GPa;Sample 2:a Young's modulus E of about 9 GPa

[0271] a yield strength σy of about 1.5 GPa;

[0272] that is, the ultra-high yield strength σy is about 1.3-1.5 GPa and the Young's modulus E is about 9-11 GPa at 80° C.

[0273] Referring to FIG. 19, in another embodiment, the present disclosure discloses a DS-STG alloy manufactured by cold rolling, the DS-STG alloy uses Ti-50Ni as a master alloy, and a manufacturing method includes the following steps:

[0274] Step 1: dislocation-induced strain glass is formed, specifically,

[0275] cold deformation: Ti-50Ni is rolled with a thickness reduction of 40% at room temperature to manufacture the dislocation-induced strain glass;

[0276] Step 2: a dual-crossover strain glass (DC-STG) state is formed, specifically,

[0277] annealing: heat preservation is performed at 300° C. for 15 min to form dual-crossover strain glass; and

[0278] Step 3: a dual-seed strain glass (DS-STG) state is formed, specifically,

[0279] cold deformation: uniaxial stretching with an elongation of 8% is performed at 80° C. to form dual-seed strain glass.

[0280] Referring further to FIG. 19, the DS-STG alloy manufactured by cold rolling described above has the following yield strength and Young's modulus:

[0281] at 20° C.:

[0282] a Young's modulus E of about 18 GPa

[0283] a yield strength σy of about 1.3 GPa;

[0284] at 40° C.:

[0285] a Young's modulus E of about 9 GPa

[0286] a yield strength σy of about 1.3 GPa;

[0287] at 60° C.:

[0288] a Young's modulus E of about 9 GPa

[0289] a yield strength σy of about 1.3 GPa;

[0290] at 80° C.:

[0291] a Young's modulus E of about 10 GPa

[0292] a yield strength σy of about 1.3 GPa;

[0293] at 100° C.:

[0294] a Young's modulus E of about 13 GPa

[0295] a yield strength σy of about 1.3 GPa;

[0296] at 120° C.:

[0297] a Young's modulus E of about 15 GPa

[0298] a yield strength σy of about 1.3 GPa.

[0299] Referring to FIG. 20, in another embodiment, the present disclosure discloses a DS-STG alloy manufactured by cold tension, the DS-STG alloy uses Ti-50.8Ni as a master alloy, and a manufacturing method includes the following steps:

[0300] Step 1: deformation-stabilized martensites are formed, specifically,

[0301] cold deformation: Ti-50.8Ni is stretched with an elongation of 35% at room temperature to manufacture the deformation-stabilized martensites;

[0302] Step 2: a dual-crossover strain glass (DC-STG) state is formed, specifically,

[0303] annealing: heat preservation is performed at 300° C. for 15 min to form dual-crossover strain glass; and

[0304] Step 3: a dual-seed strain glass (DS-STG) state is formed, specifically,

[0305] cold deformation: uniaxial stretching with an elongation of 10% is performed at room temperature to form dual-seed strain glass.

[0306] Referring further to FIG. 20, the DS-STG alloy manufactured by cold tension described above has the following yield strength and Young's modulus at 0° C.:Sample 1:a Young's modulus E of about 10 GPa

[0308] a yield strength σy of about 1.3 GPa;Sample 2:a Young's modulus E of about 8 GPa

[0310] a yield strength σy of about 1.3 GPa;Sample 3:a Young's modulus E of about 7 GPa

[0312] a yield strength σy of about 1.3 GPa;

[0313] that is, the ultra-high yield strength σy is about 1.3 GPa and the Young's modulus E is about 7-10 GPa at 0° C.

[0314] Taking together the embodiments, the present disclosure discloses a Ti-50.8 at. % Ni or Ti-50 at. % Ni “dual-seed strain glass” (DS-STG) alloy that can overcome the strength-flexibility tradeoff while exhibiting steel-like ultra-high yield strength (e.g., σy of greater than 1 GPa, e.g., σy of about 1.8 GPa) and polymer-like ultra-low Young's modulus (e.g., E of about 10.5 GPa, corresponding to ultra-high elastic compliance S=1 / E≈0.1 GPa−1, see FIG. 1a), and rubber-like ultra-large J-shaped elastic strain of about 8% (FIG. 1b). Thus, the DS-STG alloy exhibits an unprecedented flexibility (e.g., σy / E of about 0.17), far exceeding that of the existing structural materials (FIG. 1c), and the alloy behaves like a super-strong polymer (FIG. 1d).

[0315] Referring to FIG. 21, in another embodiment, the present disclosure discloses an off DS-STG alloy manufactured by cold rolling, the off DS-STG alloy uses Ti-50Ni as a master alloy, and a manufacturing method includes the following steps:

[0316] Step 1: dislocation-induced strain glass is formed, specifically,

[0317] cold deformation: Ti-50Ni is rolled with a thickness reduction of 50% at room temperature to manufacture the dislocation-induced strain glass;

[0318] Step 2: a dual-crossover strain glass (DC-STG) state is formed, specifically,

[0319] annealing: heat preservation is performed at 300° C. for 15 min to form dual-crossover strain glass; and

[0320] Step 3: an off dual-seed strain glass state is formed, specifically,

[0321] cold deformation: rolling with a thickness reduction of 10%, 20% and 30%, respectively is performed at room temperature to form a plurality of alloys deviating from dual-seed strain glass (i.e., a 10% sample, a 20% sample, and a 30% sample).

[0322] Since the alloy deviating from dual-seed strain glass is finally formed in Step 3, this alloy is referred to as an off DS-STG alloy in the present disclosure.

[0323] Referring further to FIG. 21, different samples of the above off DS-STG alloy manufactured by cold rolling have the following yield strength and Young's modulus, wherein:

[0324] 1) for the 10% sample:

[0325] at 20° C.:

[0326] a Young's modulus E of about 32 GPa

[0327] a yield strength σy of about 1.2 GPa;

[0328] at 40° C.:

[0329] a Young's modulus E of about 29 GPa

[0330] a yield strength σy of about 1.2 GPa

[0331] at 60° C.:

[0332] a Young's modulus E of about 23 GPa

[0333] a yield strength σy of about 1.2 GPa

[0334] at 80° C.:

[0335] a Young's modulus E of about 42 GPa

[0336] a yield strength σy of about 1.2 GPa;

[0337] 2) for the 20% sample:

[0338] at 20° C.:

[0339] a Young's modulus E of about 30 GPa

[0340] a yield strength σy of about 1.2 GPa;

[0341] at 40° C.:

[0342] a Young's modulus E of about 26 GPa

[0343] a yield strength σy of about 1.2 GPa;

[0344] at 60° C.:

[0345] a Young's modulus E of about 26 GPa

[0346] a yield strength σy of about 1.2 GPa;

[0347] at 80° C.:

[0348] a Young's modulus E of about 24 GPa

[0349] a yield strength σy of about 1.2 GPa;

[0350] 3) for the 30% sample:

[0351] at 20° C.:

[0352] a Young's modulus E of about 50 GPa

[0353] a yield strength σy of about 1.2 GPa;

[0354] at 40° C.:

[0355] a Young's modulus E of about 44 GPa

[0356] a yield strength σy of about 1.2 GPa;

[0357] at 60° C.:

[0358] a Young's modulus E of about 43 GPa

[0359] a yield strength σy of about 1.2 GPa;

[0360] at 80° C.:

[0361] a Young's modulus E of about 42 GPa

[0362] a yield strength σy of about 1.2 GPa;

[0363] that is, the ultra-high yield strength σy is about 1.2 GPa and the Young's modulus E is about 23-42 GPa at 20-80° C.

[0364] It can be found that the above off DS-STG alloy breaks through the performance constraint of the Young's modulus of 30 GPa.

[0365] In another embodiment, the present disclosure discloses an off DS-STG alloy manufactured by cold tension, the off DS-STG alloy uses Ti-50.8Ni as a master alloy, and a manufacturing method includes the following steps:

[0366] Step 1: deformation-stabilized martensites are formed, specifically,

[0367] cold deformation: Ti-50.8Ni is stretched with an elongation of 50% at room temperature to manufacture the deformation-stabilized martensites;

[0368] Step 2: a mono-crossover strain glass (MC-STG) state is formed, specifically,

[0369] annealing: heat preservation is performed at 200° C. for 10 min to form mono-crossover strain glass; and

[0370] Step 3: an off dual-seed strain glass state is formed, specifically,

[0371] cold deformation: uniaxial stretching with an elongation of 12% is performed at room temperature to form the off DS-STG alloy.

[0372] The strain glass microstructure of the off DS-STG alloy is shown in FIG. 22,

[0373] wherein a bright field image in (a) shows an extremely high density of dislocations entangled within a band region, where the band region is a B2 austenite twin. Through observation along the direction [1-11] of a B2 lattice, bright ½{110} superlattice diffraction spots can be found, indicating that there is long-range ordered B19′ martensite formation in the region. Further, the observation of high-resolution TEM in (b) shows that there are indeed larger-sized B19′ martensitic variants with different orientations present within the R strain glass matrix.

[0374] Considering that the lattice strain of nano-martensite domains presents the characteristics of short-range order and long-range disorder, when the alloy is in this state, it is called a strain glass state, and then, if R strain glass is not regarded as a matrix, and the high-symmetry B2 parent phase is still used as the matrix, the off DS-STG alloy undergoes plastic deformation to form a polymorphic coexistence structure of B2 austenite+nano-R martensite domain+B19′ martensite.

[0375] Referring further to the room temperature tensile stress-strain curve of the off DS-STG alloy schematically illustrated in FIG. 23 and the change of the superelasticity and the Young's modulus during loading of the off DS-STG alloy as a function of strain schematically illustrated in FIG. 24:

[0376] The tensile curve of the off DS-STG alloy has no obvious stage boundary, and has only one stage: a coexistence stage of phase transformation and strain hardening. As the tensile strain increases, the Young's modulus of the alloy is maintained at a level of approximately 17 GPa and slowly rises, for example, to 23 GPa as shown in FIG. 24(b). In connection with the example of a yield strength σy of about 1.2 GPa and a Young's modulus E of 23 GPa shown in FIG. 21, it can be seen that in the present disclosure, the Young's modulus of the off DS-STG alloy can be as low as 17-20 GPa, so that the off DS-STG alloy can have both steel-like high yield strength (σy>1 GPa) and polymer-like low modulus (E of about 10 GPa) similar to that indicated in the background.

[0377] Through a combination of FIG. 23 and FIG. 24, it can be seen that at room temperature, the off DS-STG alloy shows quasi-linear elasticity, with a very low Young's modulus of 17 GPa, a high yield strength of 1.27 GPa and a large recoverable strain of 6.8%, achieving a superelasticity close to 7%. Although the off DS-STG alloy is at a disadvantage in terms of performance compared with the DS-STG alloy, the off DS-STG alloy still has extremely outstanding ultra-high flexibility, high strength, and quasi-linear elasticity compared with other typical structural materials.

[0378] Although the embodiments of the present disclosure have been described above by way of a plurality of examples, the present disclosure is not limited to the specific embodiments and application fields described above, and the specific embodiments are only illustrative, instructional, and not restrictive. Those of ordinary skill in the art, in light of the inspiration of the specification and without departing from the scope of the appended claims of the present disclosure, can make a variety of forms, which fall within the scope of protection of the present disclosure.

Claims

1. A polymer-like ultrahigh-strength metal alloy, wherein, the alloy has the following levels of strength and Young's modulus over a wide temperature range of −80° C. to 80° C.:a ultra-high yield strength σy of greater than or equal to 1.3 GPa, and a Young's modulus E of about 10-16 GPa.

2. The alloy according to claim 1, wherein the wide temperature range is selected to be in: an aviation ambient temperature range of −80° C. to 70° C.

3. The alloy according to claim 1, wherein, the alloy comprises a strain glass matrix embedded with a small number of R and B19′ martensite seeds, i.e., martensite dual seeds.

4. The alloy according to claim 3, wherein, the R martensite seeds have a size of 50-100 nm, P=3; and the B19′ martensite seeds have a size of 50-100 nm, P=2.

5. The alloy according to claim 1, wherein, the ultra-high yield strength σy is about 1.8 GPa, and the Young's modulus E is about 16 GPa at −80° C.;the ultra-high yield strength σy is about 1.8 GPa, and the Young's modulus E is about 10 GPa at 20° C.; andthe ultra-high yield strength σy is about 1.8 GPa, and the Young's modulus E is about 15 GPa at 80° C.

6. A polymer-like ultrahigh-strength metal alloy, wherein, the alloy has the following levels of strength and Young's modulus over a wide temperature range of −80° C. to 80° C.:a ultra-high yield strength σy of greater than or equal to 1 GPa, anda Young's modulus E of about 7-20 GPa.

7. The alloy according to claim 6, wherein, the ultra-high yield strength σy is greater than or equal to 1.2 GPa, and the Young's modulus E is about 7-9 GPa.

8. The alloy according to claim 6, wherein, the ultra-high yield strength σy is about 1.2-1.8 GPa.

9. The alloy according to claim 6, wherein, the ultra-high yield strength σy is about 1.8 GPa, and the Young's modulus E is about 11 GPa at 20° C.

10. The alloy according to claim 6, wherein, the ultra-high yield strength σy is about 1.6 GPa, and the Young's modulus E is about 11 GPa at 20° C.

11. The alloy according to claim 6, wherein, the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 10.5 GPa at 20° C.

12. The alloy according to claim 6, wherein, the ultra-high yield strength σy is about 1.7-1.9 GPa, and the Young's modulus E is about 8-13 GPa at 20° C.

13. The alloy according to claim 6, wherein, the ultra-high yield strength σy is about 1.2-1.3 GPa, and the Young's modulus E is about 9-13 GPa at 20° C.

14. The alloy according to claim 6, wherein, the ultra-high yield strength σy is about 1.3-1.5 GPa, and the Young's modulus E is about 12-15 GPa at 80° C.

15. The alloy according to claim 6, wherein, the ultra-high yield strength σy is about 1.3-1.5 GPa, and the Young's modulus E is about 9-11 GPa at 80° C.

16. The alloy according to claim 6, wherein, the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 18 GPa at 20° C.;the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 9 GPa at 40° C.;the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 9 GPa at 60° C.;the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 10 GPa at 80° C.;the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 13 GPa at 100° C.; andthe ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 15 GPa at 120° C.

17. The alloy according to claim 6, wherein, the ultra-high yield strength σy is about 1.3 GPa, and the Young's modulus E is about 7-10 GPa at 0° C.

18. The alloy according to claim 6, wherein, the alloy has the following levels of strength and Young's modulus over a wide temperature range of 20° C. to 80° C.:the ultra-high yield strength σy is about 1.2 GPa, and the Young's modulus E is about 17-42 GPa at 20-80° C.; andthe alloy is an off dual-seed strain glass (off DS-STG) alloy.

19. A method for manufacturing the alloy according to claim 1, comprising the steps of:forming deformation-stabilized martensites;forming a dual-crossover strain glass (DC-STG) state; andforming a dual-seed strain glass (DS-STG) state to obtain the alloy.

20. The method according to claim 19, wherein, room temperature cold working is adopted when forming the deformation-stabilized martensites or forming the dual-seed strain glass (DS-STG) state; and heat treatment is adopted when forming the dual-crossover strain glass (DC-STG) state.

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