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Nature
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0028-0836
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10.1038/s41586-024-07900-4
Article
A polymer-like ultrahigh-strength metal alloy
Xu Zhizhi 1
http://orcid.org/0000-0002-3267-1605
Ji Yuanchao jyc.xjtu@xjtu.edu.cn

1
http://orcid.org/0000-0003-3314-5508
Liu Chang 12
He Liqiang 1
Zhao Hui 13
Yuan Ye 14
Qian Yu 1
Cui Jin 1
Xiao Andong 1
Wang Wenjia 1
Yang Yang 1
http://orcid.org/0000-0003-4637-0667
Ma Tianyu matianyu@xjtu.edu.cn

1
http://orcid.org/0000-0002-4973-2486
Ren Xiaobing ren.xiaobing@nims.go.jp

15
1 https://ror.org/017zhmm22 grid.43169.39 0000 0001 0599 1243 Multi-disciplinary Materials Research Center, Frontier Institute of Science and Technology, and State Key Laboratory for Mechanical Behaviour of Materials, Xi’an Jiaotong University, Xi’an, China
2 https://ror.org/011ashp19 grid.13291.38 0000 0001 0807 1581 School of Materials Science and Engineering, Sichuan University, Chengdu, China
3 Central R&D Institute, LONGi Green Energy Technology Co. Ltd, Xi’an, China
4 grid.9227.e 0000000119573309 Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, Beijing, China
5 https://ror.org/026v1ze26 grid.21941.3f 0000 0001 0789 6880 Center for Functional Materials, National Institute for Materials Science, Tsukuba, Japan
4 9 2024
4 9 2024
2024
633 8030 575581
22 11 2023
1 8 2024
© The Author(s) 2024
2024
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Futuristic technologies such as morphing aircrafts and super-strong artificial muscles depend on metal alloys being as strong as ultrahigh-strength steel yet as flexible as a polymer1–3. However, achieving such ‘strong yet flexible’ alloys has proven challenging4–9 because of the inevitable trade-off between strength and flexibility5,8,10. Here we report a Ti–50.8 at.% Ni strain glass alloy showing a combination of ultrahigh yield strength of σy ≈ 1.8 GPa and polymer-like ultralow elastic modulus of E ≈ 10.5 GPa, together with super-large rubber-like elastic strain of approximately 8%. As a result, it possesses a high flexibility figure of merit of σy/E ≈ 0.17 compared with existing structural materials. In addition, it can maintain such properties over a wide temperature range of −80 °C to +80 °C and demonstrates excellent fatigue resistance at high strain. The alloy was fabricated by a simple three-step thermomechanical treatment that is scalable to industrial lines, which leads not only to ultrahigh strength because of deformation strengthening, but also to ultralow modulus by the formation of a unique ‘dual-seed strain glass’ microstructure, composed of a strain glass matrix embedded with a small number of aligned R and B19′ martensite ‘seeds’. In situ X-ray diffractometry shows that the polymer-like deformation behaviour of the alloy originates from a nucleation-free reversible transition between strain glass and R and B19′ martensite during loading and unloading. This exotic alloy with the potential for mass producibility may open a new horizon for many futuristic technologies, such as morphing aerospace vehicles, superman-type artificial muscles and artificial organs.

A polymer-like ultrahigh-strength TiNi alloy was fabricated by a simple three-step process to give a combination of a polymer-like ultralow elastic modulus and a steel-like ultrahigh yield strength over a wide temperature range, and such exotic properties arise from a unique strain glass state.

Subject terms

Metals and alloys
Mechanical properties
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Exotic metal alloys showing steel-like ultrahigh strength and polymer-like ultrahigh flexibility simultaneously have long been desired for many emerging technologies such as morphing aircrafts1,2 and superman-type artificial muscles3. Such an unconventional ‘strong yet flexible’ property combination (the red circle of Fig. 1a) will enable a large recoverable shape change at low actuation cost and simultaneously provide strong resistance against fracture or yielding failure under large loads. As a result, such alloys would make possible a morphing wing in an aircraft or a super-strong artificial muscle in a humanoid robot, to name just a few possibilities.Fig. 1 A polymer-like ultrahigh-strength DS-STG alloy compared with typical metal alloys and organic materials.

a, The DS-STG alloy exhibits a combination of steel-like high strength and polymer-like ultralow Young’s modulus. This ‘strong yet flexible’ property overcomes the strength–flexibility trade-off (grey band, known as an Ashby plot4), which has rendered existing structural materials to be either ‘strong and stiff’ or ‘weak and flexible’. b, Rubber-like, J-shaped pseudoelastic stress–strain curves of the DS-STG alloy (plate and wire samples) compared with those of typical metal alloys (quenched 65Mn spring steel, Ti-6Al-4V ELI Ti alloy, 7021-T62 Al alloy and AZ91D Mg alloy) and organic materials (bamboo, FRP, PEEK and polyphenylene sulfide (PPS)) within their elasticity limit. The dotted lines represent the plastic deformation behaviour of each material beyond the elasticity limit. The higher yield strength of the wire sample (1.8 GPa) compared with the plate sample (1.3 GPa) arises from the stronger cold deformation experienced in the former (see Methods for details). The reproducibility of the result and statistical analysis of the data are presented in Extended Data Fig. 1. c, The flexibility figure of merit σy/E of DS-STG alloys far exceeds that of existing structural materials. Young’s modulus is determined by the initial slope of the stress–strain curve (inset). d, Visual evidence for the unconventional strong yet flexible property of the DS-STG alloy, which contrasts with that of two conventional materials, one being a strong and stiff spring steel and the other being a weak and flexible FRP with Young’s modulus of 11 GPa. See also Supplementary Video 1. A visual comparison of the flexibility up to the elastic limit for these three materials is given in Extended Data Fig. 2 and Supplementary Video 2.

However, achieving such unconventional alloys has remained challenging because of the inevitable trade-off between strength and flexibility (the latter being measured by elastic compliance or inverse of elastic modulus)5,8,10, as manifested by the grey band (known as the Ashby plot4) in Fig. 1a. The figure reveals a common observation4,10,11 that steels can be made very strong (with a high yield strength σy > 1 GPa) but they are stiff (with high Young’s modulus E ≈ 200 GPa), whereas organic materials such as fibre-reinforced polymers (FRPs) have the opposite property of being flexible (with a typical low E ≈ 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 modulus5,8,10. Therefore, this inevitable trade-off precludes a 1 GPa-class, steel-like, high-strength alloy from having a 10 GPa-class, polymer-like, low modulus.

Over the past decades, significant efforts have been made to seek metal alloys with simultaneously high strength and low modulus4–9,12–17, 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)5,16,17 have been reported with a 1 GPa-class high strength and a moderately low modulus of E ≈ 30 GPa, and a conventional Mg–Sc strain glass alloy7 has recently been shown to possess a lower modulus of E ≈ 20 GPa but with a lower strength of σy ≈ 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 (the red circle of Fig. 1a) still remains elusive.

Here we show that a Ti–50.8 at.% Ni dual-seed strain glass (DS-STG) alloy can overcome the strength–flexibility trade-off and demonstrate simultaneously a steel-like ultrahigh yield strength of σy ≈ 1.8 GPa and polymer-like ultralow Young’s modulus of E ≈ 10.5 GPa (Fig. 1a), together with a super-large, rubber-like, J-shaped elastic strain of around 8% (Fig. 1b). As a result, it exhibits a flexibility figure of merit of σy/E ≈ 0.17, which is very high compared with that of existing structural materials (Fig. 1c), and the alloy behaves like a super-strong polymer (Fig. 1d). In addition, the polymer-like ultrahigh-strength behaviour can persist over a wide ambient temperature range from 193 K (−80 °C) to 353 K (+80 °C) (Fig. 2a), which would enable its application in aerospace technologies, and the alloy demonstrates excellent fatigue resistance compared with other flexible materials (Fig. 2b).Fig. 2 DS-STG alloy shows polymer-like ultrahigh strength over a wide temperature range together with excellent high-strain fatigue resistance.

a, Temperature dependence of elastic modulus and yield strength over a wide temperature range from 123 K to 423 K, in which the polymer-like ultrahigh-strength behaviour of the DS-STG alloy can persist from 193 K to 353 K. The stress–strain curves of the DS-STG at various temperatures are shown in the insets. The shallow U-shaped dependence of Young’s modulus on temperature is a characteristic feature of all strain glasses, with the modulus bottom temperature Tg being the strain glass transition temperature (here Tg ≈ 280 K). The U-shaped dependence originates from the parent lattice softening at T > Tg, and the gradual dominance of elastically hardened STG nanodomains when T < Tg (ref. 34). The combination of stable ultrahigh yield strength and shallow U-shaped low modulus results in a broad peak in the flexibility figure of merit σy/E with a peak value of approximately 0.17 around room temperature, and it maintains a high value exceeding 0.06 over a wide temperature range of around 125–400 K, as shown in Extended Data Fig. 3. b, Excellent high-strain (more than 1.5%) fatigue resistance of the DS-STG compared with conventional flexible materials, including TiNi SMAs and engineering polymers such as PEEK and FRP. The upper left inset shows a comparison of the strain–life fatigue curves at a high-strain level (that is, maximum fatigue strain >1.5%) for the DS-STG and existing metal alloys (including TiNi SMAs and conventional metal alloys) and engineering polymers. The upper right inset shows that the DS-STG alloy does not fatigue even up to 5 million cycles at a sizable cyclic strain of around 1.6–1.8%. High-strain fatigue is defined by a fatigue test with the maximum cyclic strain εm > 1.5%, and the corresponding maximum stress is σm. The data source in the figure is given in the Supplementary Information.

In the following, we first demonstrate the unique mechanical properties of the alloy, then show the three-step thermomechanical fabrication route towards its unique DS-STG microstructure and finally discuss the origin of the mechanical properties in relation to the DS-STG microstructure.

Unique mechanical properties of DS-STG alloy

Figure 1a shows that the DS-STG alloy overcomes the well-observed strength–flexibility trade-off of existing structural materials (shown as the grey band of Fig. 1a) and exhibits a combination of steel-like ultrahigh yield strength of σy ≈ 1.8 GPa and polymer-like ultralow Young’s modulus E ≈ 10.5 GPa (corresponding to ultrahigh elastic compliance of S = 1/E ≈ 0.1 GPa−1). Such a property combination places the DS-STG alloy into the ‘strong yet flexible’ regime (the red circle in Fig. 1a).

Figure 1b shows that the DS-STG alloy demonstrates a rubber-like, J-shaped stress–strain curve with an ultralow initial modulus E ≈ 10.5 GPa and a super-large recoverable strain εre ≈ 8%, before reaching its elasticity limit at σy ≈ 1.8 GPa for the wire sample and σy ≈ 1.3 GPa for the plate sample. Also the loading–unloading curves have a narrow hysteresis of only around 15%. It is noted that the 1.8 GPa (or 1.3 GPa) yield strength of DS-STG alloy exceeds that of ultrahigh-strength steels such as quenched spring steel, and the 8% pseudoelastic strain not only far exceeds the small elastic strain εre ≈ 0.2–1% for typical metal alloys18–20 but also surpasses that (εre ≈ 1–5%) of many soft organic structural materials, such as wood, bamboo, bones, polymer and FRP21–24. The large J-shaped pseudoelasticity behaviour associated with a narrow stress hysteresis resembles that of ultra-flexible materials such as rubbers25, and it is very different from the superelasticity of conventional SMAs, which is characterized by a stress plateau with a much higher initial modulus of E ≈ 30–75 GPa and a large hysteresis (approximately 50–90%)26–28. The strongly hysteretic behaviour of SMAs is known to be undesirable for many applications28–30.

The combination of ultrahigh strength and ultralow modulus provides the DS-STG alloy with a super-high flexibility figure of merit of σy/E ≈ 0.17, which is much higher than that of existing structural materials including typical metal alloys and organic materials, as shown in Fig. 1c and Extended Data Fig. 3. The flexibility figure of merit σy/E is a measure of ultimate material flexibility before yielding failure10, and existing structural materials usually have a lower σy/E ≈ 0.01–0.06.

Figure 1d and Supplementary Video 1 show visual evidence for the unconventional strong yet flexible property of the DS-STG alloy, which contrasts with the conventional behaviour of two reference materials: a ‘strong and stiff’ spring steel (with high modulus E ≈ 190 GPa and high strength σy ≈ 1.36 GPa) and a ‘weak and flexible’ FRP (with low modulus E ≈ 11 GPa and low strength σy ≈ 0.13 GPa). Thus, the DS-STG alloy combines the high strength of a steel with the high flexibility of a polymer, and it behaves like a super-strong organic material, which has long been desired by many emerging technologies.

The polymer-like ultrahigh-strength behaviour of DS-STG alloy can persist over a wide temperature range from 193 K (−80 °C) to 353 K (80 °C) (Fig. 2a), which covers the aviation ambient temperature range (−80 °C to 70 °C)31 that a morphing aircraft is expected to experience. Over this wide temperature range, Young’s modulus maintains a polymer-like ultralow level of 10 to 16 GPa, which is similar to its room-temperature value of 10.5 GPa; moreover, the yield strength maintains its ultrahigh-strength level of 1.8 GPa. The recoverable strain also maintains a high value above 6.3% over this wide temperature range.

The DS-STG alloy exhibits excellent high-strain (defined by a maximum cyclic strain exceeding 1.5%) fatigue resistance as compared with other flexible materials, such as conventional TiNi SMAs and engineering polymers (for example, polyether ether ketone (PEEK) and FRP4) (Fig. 2b and upper left inset), which are candidate materials for morphing airplanes and artificial muscle fibres in which high-strain cyclic deformation is common. The upper right inset of Fig. 2b shows that the DS-STG alloy demonstrates a notable fatigue-free behaviour (with fatigue life of more than 5×106 cycles) at a sizable strain of 1.6%, which is high compared with existing metal alloys including SMAs and exceeds that of engineering polymers. Such a level of fatigue-free strain would meet the requirement of morphing wings to produce a sufficiently large camber shape change (of more than 6%)32.

Three-step thermomechanical fabrication process

The DS-STG alloy was fabricated through a three-step thermomechanical processing route from a starting solution-treated Ti–50.8 at.%Ni alloy, as shown in Fig. 3a. The simplicity of this processing method makes it scalable to industrial lines. The corresponding microstructure after each step is described below.Fig. 3 A three-step thermomechanical processing route to achieve the DS-STG alloy and the microstructure of the sample after each step.

a, Processing route and the corresponding phase state after each step. The microstructure of the starting B2 parent phase (an unfrozen strain glass) is shown in Extended Data Fig. 4a, and that of the deformation-stabilized B19′ martensite obtained after step 1 processing is shown in Extended Data Fig. 4b. b, After step 2 processing, the sample becomes a DC-STG, an unfrozen R-STG at room temperature that first undergoes a strain glass transition into a frozen strain glass around Tg ≈ 251 K, followed by crossover transitions into R and B19′ dual martensites, as revealed by Vogel–Fulcher (V–F)-type frequency (ω)-dependent Tg(ω) in the storage modulus in combination with in situ TEM and in situ XRD characterizations. Scale bars, main images, 50 nm; enlargements, 25 nm. c, After step 3 processing, the sample becomes a DS-STG, as revealed by the high-resolution TEM image (top) at room temperature. Scale bar, 10 nm. The DS-STG is characterized by aligned R and B19′ martensite seeds embedded in a DC-STG matrix. Noise-reduction processing of the image is shown in Extended Data Fig. 6. The enlarged images (middle) show the lattice fringes of R-STG nanodomains with a lattice spacing of 3d110B2, or periodicity P = 3, R martensite seed (with P = 3) and B19′ martensite seed (with P = 2). Scale bars, 2 nm. A low-magnification bright-field image (bottom) of the DS-STG reveals high-density dislocations and B2 deformation twins (evidenced by the twin diffraction pattern shown in the inset). Scale bar, 100 nm.

The starting solution-treated Ti–50.8 at.%Ni alloy (Fig. 3a(i)) is in its B2 parent state (strictly speaking an unfrozen strain glass state33–35 (Extended Data Fig. 4a)) at room temperature as it has a subzero martensitic transformation temperature of Ms ≈ 250 K. Such a B2 state exhibits a high elastic modulus of E ≈ 71 GPa and an incomplete super-elasticity with a large remanent strain of 6% and a large stress hysteresis of around 90% (Extended Data Fig. 4a) during loading and unloading even at a moderate stress of approximately 0.3 GPa. This is a well-observed behaviour of TiNi alloys when dislocation slip occurs during a stress-induced martensitic transformation14. Thus, the starting B2 alloy is neither a strong alloy nor a polymer-like alloy.

In step 1 processing, the B2 alloy is severely deformed by 50% tensile elongation (Fig. 3a(ii)). This cold-working process results in a deformation-stabilized B19′ martensite, which is preserved at room temperature (Extended Data Fig. 4b). The cold-working-strengthened B19′ martensite shows a quasilinear elastic behaviour with a high yield stress of around 1.3 GPa and a moderate elastic modulus of E ≈ 37 GPa, as shown in Extended Data Fig. 4b. Such mechanical properties are consistent with previous reports on tensile-deformed TiNi alloys12,36; thus, the deformation-stabilized B19′ martensite alloy is a high-strength alloy but not an ultralow-modulus alloy.

In step 2 processing (formation of a dual-crossover strain glass state), the deformation-stabilized B19′ martensite alloy is annealed or conditioned at 573 K for 10 min (Fig. 3a(iii)). Annealing at such a temperature fully annihilates the metastable B19′ martensite produced in step 1 processing, but is unable to annihilate or recrystallize the high-density dislocations and B2 deformation twins (to be shown below). As a result, the sample seems to be like a ‘parent B2 phase’ with a high yield strength of approximately 1 GPa (Extended Data Fig. 4c).

However, this parent B2 phase is actually a unique dual-crossover strain glass (DC-STG) with average B2 structure. The DC-STG is an unfrozen R strain glass (R-STG) at room temperature and it undergoes a strain glass transition around Tg ≈ 251 K. This is evidenced by the same signatures being shown as those of a normal strain glass transition reported in the literature33–35, including a frequency-dependent Tg following the Vogel–Fulcher law (inset of Fig. 3b) and the invariance of average B2 structure across Tg, together with the appearance of nano sized R strain domains (5–20 nm). However, being different from a conventional R-STG, which keeps R nanodomains down to 0 K, the DC-STG is able to crossover to R and B19′ dual martensites (50–100 nm in size) at lower temperatures, as evidenced by in situ transmission electron microscopy (TEM) and X-ray diffractometry (XRD) (Fig. 3b). It is noted that the crossover transitions are incomplete and a significant amount of strain glass remains even at the lowest temperature tested. This unique R-STG that can cross over into R and B19′ dual martensites is thus named ‘dual-crossover strain glass’.

It should be noted 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 behaviour, with a moderate initial elastic modulus of around 33 GPa and a high yield strength of around 1 GPa (Extended Data Fig. 4c).

In step 3 processing (formation of a DS-STG state; Extended Data Fig. 5), the DC-STG alloy undergoes a moderate tensile elongation of approximately 12% (Fig. 3a(iv)). This process introduces a small amount of aligned R and B19′ martensite dual seeds (50–100 nm in size) into the room-temperature DC-STG matrix (manifested as nanosized R domains), as revealed by a high-resolution TEM image (Fig. 3c, top). The resultant strain glass state containing R and B19′ dual seeds is thus named ‘dual-seed strain glass’. Clearly, the realization of this unique DS-STG state is due to its forerunner state DC-STG being able to crossover into two martensites R and B19′, as revealed in Fig. 3b. Such a unique DS-STG state gives this alloy a polymer-like ultralow modulus of E ≈ 10.5 GPa, together with a rubber-like, J-shaped, super-large recoverable strain of εre ≈ 8% and a narrow hysteresis of around 15% (Fig. 1b). The mechanism of such unconventional elastic behaviour is shown in Fig. 4.Fig. 4 Origin of the polymer-like elasticity of the DS-STG alloy as revealed by in situ XRD.

In situ XRD profiles during loading (left) and unloading (right) from 0 GPa and 1.3 GPa. The variation of the volume fraction of the constituent phases with applied stress during loading and unloading is shown in the pie charts. A smooth transition from the STG (with R and B19′ seeds) to the R + B19′ mixture and then to B19′ martensite occurs during loading and a smooth reverse transition occurs during unloading. The smooth evolution of the R and B19′ peaks indicates that the transition can start at nearly zero stress and that there is no nucleation barrier associated with the transition. X-ray goniometry under tensile loading is illustrated in the top inset. The possibility of internal stress as the origin of the ultralow modulus of the DS-STG is excluded in the experiment shown in Extended Data Fig. 7. a.u., arbitrary units.

The low-magnification image of the DS-STG alloy (Fig. 3c, bottom) shows that it contains high-density dislocations and a large number of B2 mechanical twins, which are formed by the severe cold-working during the thermomechanical processing. These deformation-strengthening features are characteristic of severely deformed TiNi alloys16,37–42, and can account for the ultrahigh yield strength of the DS-STG alloy (Fig. 1b).

To confirm the general applicability of the three-step thermomechanical processing in achieving polymer-like ultrahigh-strength properties, we also tested a wire sample of the Ti–50.8 at.%Ni alloy using the same three-step procedure, but the severe deformation of the sample (that is, step 1 processing) was done with a wire-drawing machine to produce a 50% area reduction instead of using a tensile machine to produce 50% elongation for the plate sample. As can be seen in Fig. 1b, the wire sample shows the same ultralow modulus of E ≈ 10.5 GPa as that of the plate sample, but with an even higher yield strength of σy ≈ 1.8 GPa due to the stronger cold deformation associated with the 50% area reduction for the cold-drawn sample. Therefore, the present three-step processing route is applicable to both wire and plate samples with step 1 deformation being either tensile elongation or cold-drawing.

Origin of the polymer-like ultrahigh strength

The in situ tensile XRD experiment (Fig. 4) shows that the DS-STG alloy with the above microstructural feature (Fig. 3c) undergoes a unique nucleation-free reversible transition between STG and R and B19′ martensites during a stress loading–unloading cycle up to 1.3 GPa. This process leads to the observed ultralow elastic modulus. Before loading (that is, σ = 0 GPa), the DS-STG appears as a DC-STG matrix embedded with a small fraction of aligned R and B19′ martensite seeds, as evidenced in the XRD profile by a broadened B2 peak (representing the average structure of the random STG nanodomains43) with a long asymmetrical tail covering R and B19′ high-angle peak positions only. With increasing stress, unlike the case in a normal stress-induced martensitic transformation in which the inevitable nucleation barrier requires a critical stress to induce martensite, in DS-STG the R and B19′ martensite peaks grow smoothly without needing a critical stress and the whole sample transforms almost fully into B19′ martensite at σ = 1.3 GPa. Clearly, this is a result of R and B19′ martensite seeds that bypass the nucleation barrier of the stress-induced transition, consequently leading to the polymer-like ultralow elastic modulus (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 DS-STG state is recovered upon unloading. The large recoverable strain (εre ≈ 8%) is a natural result of this reversible transformation for polycrystalline TiNi samples with some contribution from the elastic strain of martensite26.

Another important consequence of the nucleation-free DS-STG to R and B19′ transition is that it causes a J-shape and narrow hysteresis feature in the stress–strain curve of the DS-STG alloy, which arises from R and B19′ martensite seeds bypassing the nucleation barrier during a stress-induced transition. This results in an ultralow initial slope in the stress–strain curve, that is, the J-shape, and a narrow stress hysteresis of around 15%. 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 behaviour of DS-STG contrasts with that of the normal stress-induced STG to martensite transition (Extended Data Fig. 4c) or a conventional stress-induced martensitic transformation (Extended Data Fig. 4a), in which the inevitable nucleation leads not only to a much higher elastic modulus of E ≈ 30–70 GPa but also to a strongly nonlinear and hysteretic stress–strain behaviour with a stress plateau14,16.

Figure 5a shows that a DS-STG state is essential to achieve a polymer-like ultralow modulus. Such a state is obtained by a three-step processing with step 2 annealing at a temperature of Ta ≈ 573 K. When the step 2 annealing temperature Ta deviates from the optimal 573 K, the alloy deviates from a DS-STG state and this results in an increase in elastic modulus and a decrease in recoverable strain.Fig. 5 Polymer-like ultralow modulus and high recoverable strain of the DS-STG state occurring at Ta ≈ 573 K.

a, Young’s modulus and recoverable strain of the three-step processed Ti–50.8 at.%Ni alloy as a function of step 2 annealing temperature Ta. Polymer-like low modulus (E ≈ 10.5 GPa) and maximum recoverable strain of 8% of the DS-STG state arise from its forerunner DC-STG state obtained at Ta ≈ 573 K. b, Phase diagram of annealed samples as a function of step 2 annealing temperature Ta, in which the DC-STG state is achieved around 573 K. In the phase diagram Tnd is the nanodomain starting temperature, and 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 temperature46. These temperatures are determined by experiments shown in Extended Data Figs. 8 and 9. The alloy shows a mono-crossover strain glass (MC-STG) transition when Ta is much lower than 573 K and shows a conventional two-step martensitic transformation when Ta is much higher than 573 K. MS-STG, mono-seed strain glass.

Figure 5b explains why the DS-STG state is obtained only when the step 2 annealing temperature Ta is around 573 K. This is because the DS-STG’s forerunner state DC-STG is achieved only around Ta ≈ 573 K, as shown in the phase diagram (established from data from Fig. 3b and Extended Data Figs. 8 and 9). When Ta deviates significantly from 573 K, the forerunner DC-STG will no longer exist and consequently a R + B19′ DS-STG cannot be achieved after step 3 processing (that is, 12% elongation).

The phase diagram (Fig. 5b) shows that a significantly lower Ta (for example, Ta ≈ 473 K) will lead to a mono-crossover strain glass after step 2 processing, which is a R-STG that can cross over (partially) into one martensite B19′ at low temperature. After step 3 processing (12% elongation) it will change into a mono-seed strain glass (that is, with B19′ seeds only), which has a higher modulus of around 17 GPa (approximately 1.5 times higher than that of DS-STG) and a smaller recoverable strain of around 6%.

At a significantly higher Ta ≈ 773 K, which is a common annealing temperature for TiNi alloys26, the strain glass transition no longer exists, and the alloy is characterized by a normal two-step martensitic transformation into R and B19′ martensite. After step 3 processing (12% elongation) the alloy will change into an aligned B19′ martensite, which has an even higher modulus of around 30 GPa (ref. 44) and an even smaller recoverable strain of around 4%. This is very different from the case of DS-STG, in which R and B19′ martensite seeds, albeit aligned, make up only a small volume fraction in the DC-STG matrix, as shown in Figs. 3 and 4. On loading, these R and B19′ seeds in DS-STG grow into R and B19′ martensites at the expense of the matrix DC-STG at nearly zero stress owing to the absence of a nucleation barrier, thus resulting in ultralow Young’s modulus of around 10.5 GPa and an 8% recoverable strain, as observed in Fig. 4.

The lower elastic modulus of the DS-STG state as compared with the non-DS-STG states can be explained by 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 another, and consequently the lattice becomes ultrasoft. Such an effect has been observed in many ferroelectrics and relaxor materials associated with multiple ferroelectric phases45.

It should be noted that the thermal stability of the ultralow modulus of the DS-STG alloy over a wide temperature range (Fig. 2a) is a natural result of the DS-STG microstructure, in which R and B19′ martensite seeds can persist over a wide temperature range and they lead to the observed thermal stability of the ultralow modulus. The excellent fatigue resistance of the DS-STG alloy arises from the strengthening effect of the thermomechanical processing of the alloy that prohibits the motion and multiplication of dislocations; as a result, fatigue damage accumulation is significantly delayed or avoided.

Conclusion

A polymer-like ultrahigh-strength metal alloy DS-STG was fabricated by a simple three-step thermomechanical processing route. It has a combination of a polymer-like ultralow elastic modulus of around 10.5 GPa and a steel-like ultrahigh yield strength of around 1.3–1.8 GPa, together with a large pseudoelastic strain of approximately 8% and good temperature stability and excellent high-strain fatigue resistance. This unconventional alloy overcomes the long-standing strength–flexibility trade-off that has precluded such a long-sought-after property from being achieved. The ultralow 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 ultrahigh strength results from the deformation-strengthening effect. 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 artificial organs. The industry-scalable fabrication route may enable mass production of such alloys.

Methods

Sample preparation

Commercial Ti–50.8 at.%Ni alloy specimens from Xi’an Saite Co., Ltd and Beijing GEE SMA Technology Co., Ltd were used to fabricate all the samples. Two forms of the specimens were fabricated; one was a plate shape, which was spark-cut into a dog-bone shape with an initial gauge section dimension of 50 × 4.5 × 1 mm3, and the other was a wire shape with an initial diameter of 1 mm.

All samples were solution treated, followed by quenching in water. The plate samples were sealed in a vacuumized quartz tube and solution treated at 1,273 K for 90 min; the wire samples were solution treated in an automatic furnace kept at 1,023 K, which enabled the moving wires to enter and exit with a heating duration of 10 min before being quenched. After the solution treatment, the oxidation affected surface layer of all samples was removed by a chemical etchant.

Both plate and wire samples underwent similar three-step processing as described below and also shown in Fig. 3a, with a difference only in the deformation mode in step 1. In step 1 processing, the plate samples were severely deformed by 50% cold elongation and the wire samples were severely deformed by 50% area reduction to a final approximately 0.7 mm diameter through cold-drawing. The severe deformation in step 1 changes the starting B2 phase into a metastable martensite microstructure stabilized by high-density dislocations.

In step 2 processing, the severely deformed samples were moderately annealed at 573 K for 10 min, followed by quenching in water. This process annihilates the metastable martensite formed during step 1 processing and yields a unique DC-STG state as described in the main text. It is noted that this 573 K annealing is important to produce the ultralow modulus after step 3 processing, and this is a lower annealing temperature than the conventional annealing temperatures of TiNi alloys, which are usually between 673 K and 773 K.

In step 3 processing, both the plate and wire samples after step 2 processing were moderately deformed at room temperature by a 12% elongation, which is slightly above the transformation strain of B19′ martensite (approximately 8–9%). This step creates a small fraction of R and B19′ martensite dual seeds out of the DC-STG matrix, and such a state is called ‘dual-seed strain glass’.

Mechanical property characterization

The stress–strain curves of all the samples were measured at room temperature with a Shimadzu AG-IS tensile machine operated at a strain rate of 2 × 10−3 s−1. An extensometer (Epsilon Tech. 3442-006M-050M-LHT) was used to ensure accurate strain reading. Young’s modulus (E) was calculated from the initial slope of the stress–strain curves. The yield stress (σy) was determined to be the stress that produces 0.2% plastic deformation after unloading47. Recoverable strain (εre) was determined from the difference between the maximum and minimum strain shown in the unloading curve48.

The stress–strain curves of typical structural materials (shown in Fig. 1b) were also tested using the same tensile machine for comparison with our DS-STG alloy of wire and plate shaped samples. These structural materials include metal alloys and organic-based materials: spring steel (65Mn), Ti alloy (Ti-6Al-4V ELI), Al alloy (7021-T62), Mg alloy (AZ91D), bamboo (dry), FRP (PPS GF40), PEEK and PPS. It is noted that the DS-STG wire sample had a higher yield strength (approximately 1.8 GPa) as compared with the DS-STG plate sample (approximately 1.3 GPa). This difference stems from the stronger work-hardening in the wire sample, because this sample (with 50% area reduction in step 1 processing) experienced stronger work-hardening than the plate sample (with 50% elongation, equivalent to a smaller area reduction of about 33%).

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 curves were the second cycle of a tensile test to 1.8 GPa.

The fatigue life N versus maximum cyclic stress σm curve was measured using a fatigue testing machine (TA 3330 Series III) at room temperature in load-controlled tension–tension mode between 0 to the maximum stress σm. The testing frequency was 0.5–10 Hz. The corresponding maximum cyclic strain εm was obtained from the experimental stress–strain relation at 20 °C (Fig. 2a, upper-middle inset), except for the fatigue-free samples (with 5 × 106 cycles), for which the strain εm was directly measured using an extensometer during its first cycle and the runout cycle (Fig. 2b, right inset).

In situ structural and microstructural characterizations

Crystal structure characterization at different temperatures was carried out with an X-ray diffractometer (Shimadzu XRD-7000) equipped with a temperature sample holder. Cu Kα radiation was used as the incident X-ray source. In situ XRD under tensile deformation was performed with a self-made tension holder, which enabled in situ tension of a small XRD sample with a gauge dimension of 40 × 2 × 0.05 mm3.

Microscopic characterization was carried out with a JEM 2100 F TEM equipped with a field-emission gun and a cooling holder. The data were recorded using a GATAN CCD slow scan camera and analysed by DigitalMicrograph software. The TEM specimens were made by mechanical thinning down to about 0.1 mm followed by twin-jet polishing with a solution of 90% methanol and 10% sulfuric acid.

Characterization of martensitic transformation and strain glass transition by using a differential scanning calorimeter, dynamical mechanical analyser and electrical resistivity measurements

A differential scanning calorimeter (DSC-Q200) from TA Instruments was used to monitor the occurrence of the martensitic transformation through monitoring its latent heat peak, and the heating–cooling rate was 10 K min−1. As a strain glass transition does not produce a significant differential scanning calorimeter peak, a dynamical mechanical analyser (DMA-Q850) from TA Instruments was used to monitor the key signature of a strain glass transition: a Vogel–Fulcher-type frequency-dependent anomaly in the storage modulus curve. The measurement was done in a single cantilever mode with an amplitude of 15 μm at 1, 2, 5, 10 and 20 Hz. The cooling–heating rate was 2 K min−1. Four-probe electrical resistivity measurements were used to monitor the R and B19′ martensitic transformation and strain glass transition under a constant current 100 mA and a cooling–heating rate of 2 K min−1.

Online content

Any methods, additional references, Nature Portfolio reporting summaries, source data, extended data, supplementary information, acknowledgements, peer review information; details of author contributions and competing interests; and statements of data and code availability are available at 10.1038/s41586-024-07900-4.

Supplementary information

Supplementary Information Full descriptions of Supplementary Videos 1 and 2, and Tables 1–4.

Supplementary Video 1 A metal alloy showing polymer-like ultrahigh flexibility and steel-like ultrahigh strength simultaneously.

Supplementary Video 2 A visual comparison of flexibility up to the elastic limit among spring steel, a FRP with 11 GPa Young’s modulus and DS-STG alloy using an elastic bending test.

Peer Review File

Extended data figures and tables

Extended Data Fig. 1 Reproducibility of polymer-like ultrahigh-strength property of DS-STG wire and plate samples and statistical analysis of their mechanical properties.

Stress-strain curves of 8 DS-STG samples (both wire and plate) were tested to check reproducibility of the property. Statistical analysis of the property is summarized in the table.

Extended Data Fig. 2 A visual comparison of flexibility up to elastic limit among Spring steel, Fiber-reinforced polymer (FRP with ~11 GPa Young’s modulus), and DS-STG alloy by elastic bending test.

Material flexibility at elastic limit is measured by the maximum bending curvature (or minimum bending radius) without permanent deformation or fracture. The minimum elastic bending radius is ~85 mm, ~8 mm, and ~6 mm for spring steel, FRP, and DS-STG, respectively.

Extended Data Fig. 3 Flexibility figure of merit of DS-STG alloy as a function of temperature.

a, DS-STG alloy undergoes a strain glass transition with freezing temperature Tg ~ 280 K, which is the storage modulus dip temperature showing Vogel-Fulcher (V-F) type frequency-dependence (insets), the Vogel-Fulcher relation ω = ω0exp(-Ea/kB(Tg-T0)), where ω is the frequency, ω0 is the frequency prefactor, Ea is the activation energy, kB is the Boltzmann constant and T0 ~ 276 K is the ideal freezing temperature. b, Flexibility figure of merit σy/E of DS-STG alloy as a function of temperature, calculated from the measured yield strength σy and Young’s modulus E at each temperature from Fig. 2a. c, Flexibility figure of merit of DS-STG alloy at room temperature (293 K) as compared with existing structural materials.

Extended Data Fig. 4 The room-temperature microstructure and mechanical property of Ti-50.8Ni alloy in three different states.

a, Starting solution-treated state, which exhibits an average B2 structure but contains nano-sized R domains manifested as 1/3{110}B2 faint spots in the diffraction pattern and lattice fringes with a spacing of 3d110B2 or P = 3 shown in high-resolution TEM and noise-filtered inverse fast Fourier transform (IFFT) images. Such a state of B2 alloy has been identified to be an unfrozen strain glass33–35. Imperfect superelasticity with low yield strength (~0.32 GPa) and high Young’s modulus (~71 GPa) associated with the solution-treated Ti-50.8Ni alloy which has a sub-zero martensitic transformation temperature of Ms ~ 250 K (Extended Data Fig. 8). b, Severely-deformed state (i.e., after step-1 processing), which reveals high-density dislocations and deformation-stabilized B19’ martensite plates at room temperature. The metastable B19’ martensite is identified by XRD (inset) and in the selected area diffraction pattern (SADP). The severely-deformed Ti-50.8Ni alloy reveals a quasi-linear elastic behavior with a high strength ~1.3 GPa and a moderate elastic modulus of ~37 GPa. c, Severely-deformed and 573 K-annealed state (i.e., after step-1 plus step-2 processing, or DC-STG state), which exhibits an average B2 structure but contains nano-sized R domains manifested as 1/3{110}B2 diffuse spots in the diffraction pattern. The 573 K-annealing in step-2 processing to the deformed Ti-50.8Ni alloy changes the dislocation-stabilized B19’ martensite into an R-STG. High density dislocations (black patches) and B2 twins (with sharp boundaries) in the bright-field image are visible. The phase transition behavior of DC-STG is shown in Fig. 3b. DC-STG shows a nearly flag-shaped superelastic deformation behavior with large non-linearity and hysteresis; it demonstrates a moderately high strength of ~1.0 GPa and a moderate initial modulus of ~33 GPa. Therefore, DC-STG state is not an ultralow modulus state, unlike DS-STG state (dotted stress-strain curve).

Extended Data Fig. 5 Stress-strain curve of DC-STG alloy (Ta = 573 K) during step-3 processing (i.e., ~12% plastic elongation at room temperature).

The phase constituents of the sample before and after deformation were probed by XRD (bottom profiles). Before deformation the DC-STG is a R strain glass with average B2 structure, whereas after 12% elongation a small fraction of dislocation-stabilized R and B19’ martensite seeds (~10% in volume fraction) remain in the DC-STG matrix.

Extended Data Fig. 6 Noise reduction process for high-resolution TEM image.

Original high-resolution TEM image (top left) and corresponding diffraction pattern (top right, showing sharp R and B19’ spots from R and B19’ seeds). Fast Fourier transform (FFT) pattern (bottom left) is produced from the high-resolution image (top left). The filtered inverse FFT (IFFT) image is produced by using the circled spots in the FFT pattern.

Extended Data Fig. 7 Young’s modulus and residual strain of 3-step processed Ti-50.8Ni alloy as a function of Step-2 annealing time at 573 K.

Both ultralow modulus of ~10 GPa and complete reversibility are achieved at an optimal aging time of 10 min (DS-STG). It suggests that the ultralow modulus of DS-STG has a microstructural origin rather than an internal stress origin. This is because the internal stress, which is created during the deformed martensite transforming into parent phase upon heating to 573 K, would be strongest for short-time (i.e., 0.5 min) annealed sample, but this sample has the highest modulus of ~39 GPa. This can be explained only by a microstructural origin as shown in Fig. 4. Long annealing time reduces yield strength and thus leads to incomplete strain recovery due to dislocation slip.

Extended Data Fig. 8 Determination of phase diagram of Fig. 5b by a combination of DSC, electrical resistivity, and DMA measurements.

Tnd, Rs, Ms, Tg, and Ts represent nanodomain starting temperature, R martensitic transformation starting temperature, B19’ martensitic transformation starting temperature, strain glass transition temperature, and spontaneous transition temperature from STG to R martensite, respectively. a1, b1, c1, show the starting solution-treated Ti-50.8Ni alloy undergoes transitions from B2 parent phase to unfrozen STG (at Tnd) then to B19’ martensite (at Ms), being consistent with recent reports33–35. a2, b2, c2, a3, b3, c3, annealing at low Step-2 temperatures Ta = 473 K and Ta = 523 K leads to a mono-crossover strain glass (MC-STG) state, which first shows a strain glass transition into a frozen R strain glass and followed by a crossover transition to a single martensite B19’. The strain glass transition is evidenced by a nearly vanishing DSC peak, and a Vogel-Fulcher type frequency-dependent elastic modulus dip at Tg ~ 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 of B19’ formation, as well as the appearance of B19’ peaks at low temperature in XRD profiles (Extended Data Fig. 9a). a4, b4, c4 annealing at a moderate Step-2 annealing temperature Ta = 573 K leads to a dual-crossover strain glass (DC-STG) state, which first shows a strain glass transition into a frozen R strain glass and followed by a dual-crossover transition to R martensite then to B19’ martensite. The strain glass transition is evidenced by a Vogel-Fulcher type frequency-dependent elastic modulus dip at Tg ~ 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 profiles (Extended Data Fig. 9b), which is a key measurement for a crossover strain glass transition34. 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, c7 annealing at high Step-2 annealing temperatures Ta = 623, 673, 773 K, which is the standard annealing temperature for Ti-Ni alloys, leads 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 below Rs and the existence of a ~ 5 K transition hysteresis characteristic 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 literature49,50. Direct evidence from XRD (Extended Data Fig. 9c) confirms this transition sequence.

Extended Data Fig. 9 In-situ XRD profiles during cooling for (a) a mono-crossover strain glass (MC-STG) sample (with Ta = 473 K), (b) a dual-crossover strain glass (DC-STG) sample (with Ta = 573 K), and (c) a conventional R and B19’ martensite sample (with Ta = 773 K).

a, MC-STG sample maintains an average B2 structure below strain glass transition temperature Tg, and crossovers into single B19’ martensite below Ms, but the transition is incomplete. b, DC-STG sample maintains an average B2 structure down to strain glass transition temperature Tg, then transforms to R phase first and then to B19’ martensite, but the transitions are incomplete. c, Normal R/B19’ martensitic sample with high annealing temperature (Ta = 773 K) undergoes a spontaneous transformation from unfrozen STG (Extended Data Fig. 4a, R local structure with an average B2 structure33–35) to R phase and finally to B19’ martensite.

Supplementary information

The online version contains supplementary material available at 10.1038/s41586-024-07900-4.

Acknowledgements

This research was supported by the National Key R&D Program of China (grant no. 2022YFB3808700), the National Natural Science Foundation of China (grant nos. 51831006, 52071257, 12204368, 52301251 and U2241242), Key R&D Program of Shaanxi Province (grant no. 2021GXLH-Z-041) and the National 111 Project 2.0 (grant no. BP0618008). X.R. acknowledges support from JSPS Kakenhi (grant no. 23H01675).

Author contributions

Z.X. made the initial key finding. Y.J. and X.R. conceived the project. Z.X. and Y.J. fabricated the samples. Z.X. performed mechanical testing and property and phase-transition characterizations. Z.X., T.M., C.L. and L.H. performed the TEM observations. X.R. provided the physical interpretation. Z.X. and H.Z. carried out the in situ XRD testing. Z.X., Y. Yuan and Y. Yang analysed the in situ XRD tests. Z.X., Y.Q., A.X. and W.W. carried out the fatigue tests and analysed the results. Z.X. and J.C. tested the reproducibility of the results and performed the statistical analysis. Z.X. analysed all results with contributions from all the authors. Z.X., Y.J., T.M. and X.R. wrote the manuscript with input from all the authors.

Peer review

Peer review information

Nature thanks Jun Cui, Petr Sittner and Tong-Yi Zhang for their contribution to the peer review of this work. Peer reviewer reports are available.

Data availability

All data generated or analysed during this study can be accessed from the Science Data Bank (10.57760/sciencedb.12257).

Competing interests

The authors declare no competing interests.

Publisher’s note Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.
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