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## Creator

[Norimasa Nishiyama](https://orcid.org/0000-0001-6214-0875), Yoshinori Tange, Takashi Sawahata, Masafumi Matsushita, Kazuya Tokuda, Kosuke Tominaga, Takashi Sekiya, Koji Kuramochi, Keita Sasaki, Fumihiro Wakai, Zenji Horita, Yutaka Kobayashi, Akio Fujimura

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The version of record of this article, first published in Scientific Reports, is available online at Publisher’s website: https://doi.org/10.1038/s41598-025-33037-z[Creative Commons BY-NC-ND Attribution-NonCommercial-NoDerivs 4.0 International](https://creativecommons.org/licenses/by-nc-nd/4.0/)

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[High tensile strength and transformation-induced plasticity in bulk polycrystalline omega titanium](https://mdr.nims.go.jp/datasets/5b83daea-aa0e-435d-8f5e-ec0ff778a75c)

## Fulltext

1  High tensile strength and transformation-induced plasticity in bulk polycrystalline omega 1 titanium 2  3  4 Norimasa Nishiyama1,2,3*, Yoshinori Tange1, Takashi Sawahata2, Masafumi Matsushita4, 5 Kazuya Tokuda5, Kosuke Tominaga5, Takashi Sekiya5, Koji Kuramochi5, Keita Sasaki1, 6 Fumihiro Wakai2, Zenji Horita6,7,8, Yutaka Kobayashi1, Akio Fujimura1 7  8  9 1: Advanced Materials Laboratory, Sumitomo Electric Industries, Ltd, 1-1-1 Koyakita, Itami 10 664-0016, Japan 11 2: Laboratory for Materials and Structures, Tokyo Institute of Technology, R3-22, 4259 12 Nagatsuta-cho, Midori-ku, Yokohama 226-8503, Japan 13 3: Research Center for Materials Nanoarchitectonics, National Institute for Materials Science, 14 1-1 Namiki, Tsukuba 305-0044, Japan 15 4: Department of Mechanical Engineering, Ehime University, 3 Bunkyo-cho, Matsuyama 16 790-8577, Japan 17 5: Analysis Technology Research Center, Sumitomo Electric Industries, Ltd, 1-1-1 Koyakita, 18 Itami 664-0016, Japan 19 2  6: School of Engineering, Kyushu Institute of Technology, 1-1 Sensui-cho, Tobata-ku, 1 Kitakyushu 804-8550, Japan  2 7: Magnesium Research Center, Kumamoto University, 2-39-1 Kurokami, Chuo-ku 3 Kumamoto 860-8555, Japan 4 8: Synchrotron Light Application Center, Saga University, 1 Honjo-machi, Saga, 840-8502, 5 Japan 6  7 *Corresponding author: Norimasa Nishiyama 8 Present affiliation: Research Center for Materials Nanoarchitectonics, National Institute for 9 Materials Science, 1-1 Namiki, Tsukuba 305-0044, Japan 10 Email: nishiyama.norimasa@nims.go.jp 11 3  Abstract 1 Titanium and its alloys exhibit advantageous ductility and strength-to-weight ratios, which 2 makes them suitable for use as structural materials in numerous industrial applications. The 3  phase has been observed to precipitate during the aging process of titanium alloys, 4 resulting in a loss of ductility. Here we report tensile behavior of bulk polycrystalline -5 titanium with a chemical composition of commercially pure titanium grade 4 and an average 6 grain size of 3.4 m. We observed that stress-induced  →  martensitic phase 7 transformation occurs exclusively in the plastic regime. As plastic deformation proceeds, the 8 volume fraction of -phase increases. The 0.2% offset yield strength, tensile strength, and 9 elongation to failure were determined to be 1130 ± 30 MPa, 1220 ± 30 MPa, and 16 ± 2 %, 10 respectively. The present study revealed transformation-induced plasticity in this material. 11 The mechanical properties of this material with the pure titanium composition are 12 comparable to those of a titanium alloy Ti-6Al-4V. The bulk polycrystalline -titanium can 13 potentially be utilized for biomedical applications, such as dental implants. 14  15 Keywords: 16 Titanium;  phase; high-pressure and temperature; tensile strength; martensitic 17 transformation; transformation-induced plasticity  18 4  Introduction 1 Titanium exhibits three major phases, designated as , , and  phases, which vary 2 according to pressure and temperature1. At ambient conditions, the substance stabilizes in a 3 hexagonal closed-packed (hcp) structure, designated as the  phase. However, at ambient 4 pressure and temperatures greater than 1155 K, the  phase undergoes a phase transition to a 5 body-centered cubic (bcc) structure, which is designated as the  phase. Moreover, under 6 conditions of high pressure, the  phase transforms into a simple hexagonal structure, 7 referred to as the  phase2. The  phase is stable at least up to 100 GPa and it further 8 transforms into high-pressure phases, called  and  phases3,4. 9 The occurrence of metastable  phase in -Ti alloys has been observed at ambient 10 pressure after heat treatment and aging process5,6. The presence of metastable  phase has 11 been shown to have a significant impact on the mechanical properties of these alloys, causing 12 substantial increases in strength and significant reductions in ductility7, or, in some cases, 13 leading to a complete loss of ductility8,9. The embrittlement phenomenon associated with the 14  phase10 has posed a significant challenge to the utilization of these alloys in safety-critical 15 applications such as medical devices. Consequently, numerous studies have been conducted 16 to suppress the formation of  phase11,12. Recently, new approaches have emerged to 17 overcome this problem by controlling the growth of the  phase during aging and by 18 preserving the  matrix stability. These attempts have enhanced the strength through 19 5  nanoprecipitation of the  phase while maintaining good ductility by preserving the 1 TRIP/TWIP effect, particularly in Ti-12Mo alloy13,14. Understanding the basic properties of 2 the -phase is crucial for improving the mechanical properties of Ti alloys. 3 In order to measure the fundamental properties of the  phase, studies have been 4 conducted on the fabrication of polycrystalline titanium materials consisting primarily of the 5  phase.  Pioneering work in this field was performed by Todaka et al15, who employed high-6 pressure torsion (HPT) processing at 5 GPa and room temperature (30˚C) to severely deform 7 a pure -titanium specimen. Consequently, the authors obtained 90 mass% polycrystalline  8 phase with a grain size measured in hundreds nanometers. The Vickers hardness of this 9 material was reported to be 3.7 GPa, a value that exceeds that of -titanium. Tane et al.16 10 employed the same technique to produce polycrystalline pure titanium specimens consisting 11 exclusively of the  phase. A complete set of elastic stiffness of the  phase was measured, 12 and the reported values were found to be consistent with those predicted by theoretical 13 calculations17. Two other studies employed a static high-pressure technique to fabricate bulk 14 polycrystalline specimens consisting of a single-phase of the  phase in pure titanium. Li et 15 al.18 fabricated a specimen with dimensions of 3 × 3 × 4 mm3 under 12.2 GPa and room 16 temperature. The authors reported the Vickers hardness to be 2.5 GPa and the measured 17 elastic moduli are consistent with those reported by Tane et al16. Sawahata et al19 fabricated 18 rod-shaped specimens with a diameter of 4 mm and a height of 3 mm under simultaneous 19 6  high pressure (12 GPa) and high temperature (400˚C). The uniaxial compressive stress-strain 1 relationship was obtained and the 0.2% offset yield strength was determined to be 913 MPa. 2 The bulk polycrystalline -Ti material was reported to be more than twice as strong as the  3 phase19. However, the tensile behavior of this material has never been observed. 4 In the present study, tensile tests were performed on bulk polycrystalline  titanium 5 materials that were fabricated in the thermodynamic stability field under high-pressure and 6 high-temperature (high-PT) conditions. The stress-strain curves were obtained, and the data 7 were analyzed to determine yield strength, tensile strength, and elongation to failure. Time-8 resolved X-ray diffraction measurements were also carried out during tensile tests using 9 synchrotron radiation. We observed that the  phase transforms into the  phase during 10 plastic deformation. 11  12 Results and discussion 13 In this study, four types of pure titanium materials (-Ti) were utilized as starting 14 materials to fabricate bulk polycrystalline  titanium materials (-Ti materials) under high-15 PT: titanium with purity of 99.999 wt% metals basis (5N) (0.02 wt% oxygen); commercially 16 pure titanium grade 2 (CP-Ti Grade 2) (0.12 wt% oxygen); CP-Ti Grade 4 (0.38 wt% 17 oxygen); cold-worked CP-Ti Grade 4 (0.37 wt% oxygen). The minor element contents of 18 these CP-Ti materials are given in Supplementary Table 1. In each case, the presence of 19 7  single phase  was confirmed by X-ray diffraction measurements (XRD) after the high-PT 1 treatment (Figs. 1a-d). It is noteworthy that despite the literature on the blocking of the 2 martensitic  →  transformation by certain impurities15,20-22, CP-Ti Grade 4 readily 3 transformed to the  phase in the high-PT treatment. The microstructures of the -Ti 4 materials were examined using field-emission scanning electron microscopy (FESEM) with 5 an electron backscatter diffraction (EBSD) detector. The obtained images demonstrate a 6 decrease in grain size with increasing impurity content (Figs. 1e-g). The -Ti material 7 exhibiting the smallest grain size (Fig. 1h) was fabricated by utilizing cold-worked CP-Ti 8 Grade 4 as the starting material and by minimizing the duration of heating under high 9 pressure (Supplementary Figure 1). 10 8   1 The mechanical properties of the -Ti materials were evaluated using miniature tensile 2 test specimens. The specimens were prepared from the rod-shaped samples (see the inset in 3 Fig 1d) using wire electrical discharge machining (EDM). After wire EDM cutting, a 4 chemical polishing procedure was conducted using a commercially available liquid 5 containing a nitric-hydrofluoric acid mixture to remove the damaged surface area19. After 6 Figure 1.  The characterization of polycrystalline -Ti materials with pure titanium compositions fabricated under high pressure and temperature conditions is presented. Four different starting materials (all  phase) were employed for the high-PT treatment: Ti 99.999 wt% metals basis (5N); CP-Ti Grade2 (CP2); CP-Ti Grade 4 (CP4); cold-worked CP-Ti Grade4 (CP4-cw). The X-ray diffraction patterns of the -Ti materials fabricated from 5N, CP2, CP4, and CP4-cw are shown in a) – d), respectively. It is evident that all the materials consist solely of the  phase. The inset of d) presents a photograph of the titanium materials: the bottom, a starting material (CP4-cw); the middle, an -Ti material after high-PT treatment; the top, an -Ti material fabricated from CP4-cw after surface machining by a centerless grinder. The diameter of this object is 4.9 mm and its length is 17.7 mm. The samples were placed on a 5 mm grid paper. e) – h) were inverse pole figure (IPF) maps superimposed on band contrast (BC) images of the -Ti materials fabricated from 5N, CP2, CP4, and CP4-cw, respectively. The inset of e) shows the color coding for the orientations in e) – h). It is clear that grain size decreases from e) to h) (see text). 9  chemical polishing, we conducted XRD measurements on the tensile specimens. The XRD 1 profiles were essentially identical to those shown in Figures 1a-d, and the presence of a single 2 phase of the  phase in the tensile specimens was confirmed. As both the starting ( phase) 3 and -Ti materials were tested, eight different stress-strain curves were obtained, and these 4 curves are displayed in Figure 2. It is evident that, for each pure titanium composition, both 5 the -Ti and -Ti materials exhibited ductile behavior. We can also see that, for each 6 composition, tensile strength of the -Ti material is greater than that of the -Ti material. In 7 addition, the impurity content exhibited a direct correlation with the yield strength and tensile 8 strength of both the  and  phases, with an increase observed from 5N to CP-Ti Grade 4 9 (see Figs. 2a-c). The impurity effect observed in the -Ti material has been extensively 10 documented in the literature23,24 and is attributed to the presence of interstitial oxygen 11 atoms22-25, which have been shown to enhance the mechanical properties of the material. The 12 present data indicate that the same strengthening mechanism is activated in the -Ti materials. 13 Furthermore, as illustrated in Figures 2c and d, the strengths of the -Ti material increase 14 with decreasing grain size (Figs. 1g and h). The strengthening of the -Ti material in pure 15 titanium due to grain size reduction is well documented26 and is explained by the Hall-Petch 16 effect27,28. The observed strengthening of the -Ti material can be attributed to this effect. 17 10   1 The 0.2% offset yield strength (0.2%), tensile strength (TS), and elongation to failure (f) 2 were determined for all eight materials. The results are summarized in Table 1. For the -Ti 3 material, which exhibited an average grain size of 3.4 ± 1.1 m (Fig. 1h), fabricated from 4 cold-worked CP-Ti grade 4, the following values were obtained. The yield strength, 0.2%, 5 Figure 2.  Stress-strain curves of the starting () and -Ti materials with pure titanium compositions obtained by miniature tensile tests. a) 5N (blue) and the -Ti material fabricated from 5N (red); b) CP2 (blue) and the -Ti material (red); c) CP4 (blue) and the -Ti material (red); d) CP4-cw (blue) and the -Ti material (red). The inset of d) shows a photograph of a miniature tensile specimen. In Fig. 2a, “” refers to a stress-strain curve obtained from a bulk polycrystalline -titanium material; “” refers to a stress-strain curve obtained from a bulk polycrystalline -titanium material (see text). 11  was determined to be 1130 ± 30 MPa, while the tensile strength, TS, was found to be 1220 ± 1 30 MPa. The elongation to failure, f, was measured to be 16 ± 2 %. The tensile strength and 2 elongation to failure of this material are comparable to those of the most commonly used 3 titanium alloy in industry, Ti-6Al-4V (Table 1 and Supplementary Figure 2). Figure 2 further 4 illustrates that the -Ti materials demonstrate notable ductility, as exemplified by the 5 elongation to failure of the -Ti material with CP-Ti grade 2 composition, which was 6 determined to be 34 ± 1 %.  7  8 Table 1. Summary of mechanical properties of pure titanium materials determined by miniature tensile tests. 9 Phase   +    Material 5N  CP2  CP4  CP4-cw  Ti64   5N  CP2  CP4  CP4-cw Grain size (m) -   -   -   -    -    65(16) 12(4) 5.2(17) 3.4(11) 0.2% (MPa) 118(18) 257(15) 640(50) 973(5)  990(40)  264(9) 634(13) 990(20) 1130(30) TS (MPa) 284(11) 574(14) 930(20) 1121(7)  1190(70)  527(6) 833(11) 1090(30) 1220(30) f (%) 64(8) 40(2) 28(2) 22(1)  25(3)  38(2) 34(1) 30(2) 16(2) The averages of the three tensile test runs for each material are shown with one standard deviation in parentheses. 10 Phase: a phase or phases that are present in the tensile specimens before starting the tensile tests. 11 5N: 99.999% Ti; CP2: CP-Ti Grade2; CP4: CP-Ti Grade4; CP4-cw; cold-worked CP-Ti Grade4; Ti64: Ti-6Al-4V. 12 0.2%: 0.2% offset yield strength; TS: tensile strength; f: elongation to failure. 13  14 In order to identify the phases present in a fractured specimen of the -Ti materials, a 15 specimen with CP-Ti grade 2 composition was selected and X-ray diffraction (XRD) 16 measurements were performed using transmission mode with high-energy synchrotron 17 radiation. The XRD profiles were collected at 0.5 mm intervals along the tensile axis of the 18 specimen. The value of D was defined based on the distance from the edge of grip section 19 12  toward the fracture surface (Fig. 3a). The fracture surface was positioned at D = 5.0 mm. The 1 profiles collected at distances between D = 2.0 and 5.0 mm are displayed in Figure 3b. The 2 XRD profiles collected at D = 2.0, 2.5, and 3.0 mm (from the grip section to the shoulder of 3 the specimen) were found to be similar to those obtained prior to the tensile test (Fig. 1b). 4 The XRD peaks of the  phase were solely observed. At D = 3.5 mm, XRD peaks of the  5 phase appear and the intensity of these peaks increases with the distance D (4.0 and 4.5 mm). 6 At the position including the fracture surface (D = 5.0 mm), XRD peaks of the  phase 7 disappear and those of the  phase are exclusively observed. It is noteworthy that the 8 appearance of the  phase was also detected in fractured tensile specimens in other pure 9 titanium compositions (Supplementary Figure 3). The XRD measurements of the fractured -10 Ti materials demonstrated that the → transformation occurs in the gauge sections during 11 the tensile tests. This observation indicates that the stress-strain curves of the -Ti materials 12 are influenced by this transformation. It should therefore be noted that the values reported in 13 Table 1 are not representative of those for the -phase itself. 14 13   1 An investigation into the microstructure of a fractured -Ti material (CP-Ti Grade2 2 composition) was conducted through the utilization of the FESEM-EBSD technique. As can 3 be seen in Figure 4a, a cross section of the fractured tensile specimen was prepared. The 4 observations were performed at four different positions where the  phase is present: D = 3.5, 5 4.0, 4.5 and 5.0 mm (Fig. 4b). Backscattered electron images (BEIs, Fig. 4c) demonstrate that 6 grain size varies as a function of the distance D. At the position of D = 3.5 mm, the average 7 grain size of the  phase is comparable to that of an undeformed -Ti material (Fig. 1f). It is 8 evident that the grain size decreases with the distance D. Furthermore, the presence of 9 micropores near the fracture surface (D = 5.0 mm) was observed. These observations suggest 10 Figure 3. X-ray diffraction profiles obtained from a fractured -Ti tensile specimen. The -Ti material was fabricated from CP2. The fractured specimens were placed on a 5 mm grid paper. a) A photograph of fractured tensile specimens and definition of the distance from the edge of grip section, D. The beam size (1.0 mm × 0.5 mm) is shown by the dashed squares. b) X-ray diffraction profiles collected at different D positions from D = 2.0 mm (in the grip section) to D = 5.0 mm (the position including fracture surface). XRD peaks of the  phase can be seen in the gauge section between D = 3.5 and 5.0 mm. 14  that plastic deformation leads to a reduction in grain size29 and the formation of pores30. As 1 can be seen from the band contrast (BC) images (Fig. 4d), the proportion of darker areas (i.e., 2 weaker intensities of Kikuchi bands) increases with distance D. These observations suggest 3 that plastic deformation leads to the formation of lattice defects31. The EBSD mappings were 4 also used for phase analysis (Fig. 4e). At the position of D = 3.5 mm, the presence of small 5 particles of the  phase was observed at grain boundaries of the  phase. Moreover, at D = 6 4.0 mm, the presence of lamellar  phase was observed within the  phase matrix. The 7 proportion of the  phase increases at D = 4.5 mm and only the  phase was observed at the 8 position including the fracture surface (D = 5.0 mm), which is consistent with the XRD data 9 (Fig. 3b). These microstructure observations as well as the XRD data (Fig. 3) demonstrate 10 that stress-induced  →  transformation occurs during tensile deformation. 11 One of the unknown factors is the effect of hydrogen on the -Ti materials. In the present 12 study, miniature tensile specimens were prepared using wire EDM cutting followed by 13 chemical polishing. During these processes, hydrogen might be absorbed by the specimens, 14 affecting their tensile behavior, including the stress-induced  →  transformation. It is well 15 known that hydrogen decreases the ductility and creep resistance of CP-Ti materials32. Future 16 studies are needed to measure the hydrogen content of the -Ti materials and to examine the 17 effect of hydrogen on the mechanical properties of these materials. 18 15   1  2  3  4 Figure 4. Microstructure observations of a fractured -Ti tensile specimen. The -Ti material was fabricated from CP2. a) Photograph of a fractured specimen. The dashed line indicates the position of the cross section for the observations. The cross section is perpendicular to the flat surface of the tensile specimen. b) Overview back-scattered electron image (BEI) of the investigated area. Detailed observations were conducted at positions where  →  transformations occur: D = 3.5, 4.0, 4.5 and 5.0 mm. BEIs (c), band contrast (BC) images (d), and phase maps (e) are shown at these positions. In c), the red arrows indicate the locations of pores. At D = 4.0mm, the lamellar  phase was observed within the  phase matrix. 16  It has been well established that martensitic  →  transformation occurs in titanium33-35. 1 The martensitic transformation induces two types of orientation relationships between the  2 and  phases. They are referred to as Variant I (OR I) and Variant II (OR II) following 3 Usikov and Zilbershtein’s notation33: 4 OR I: (0001)𝛼 ∥ (0111)𝜔,  [1120]𝛼 ∥ [1011]𝜔, 5 ORII: (0001)𝛼 ∥ (1120)𝜔, [1120]𝛼 ∥ [0001]𝜔. 6 Given the fact that the  phase is metastable at ambient conditions, the reverse process,  → 7  transformation, has scarcely been studied. Zahili et al.36 conducted molecular dynamics 8 simulations and theoretically investigated the deformation process of -Ti. The authors 9 predicted that the  →  transformation also follows either OR I or OR II and demonstrated 10 that OR I is energetically more favorable than OR II (the energy barrier of OR I is 11 approximately one-quarter of that of OR II). In order to investigate the orientation 12 relationships between the  and  phases in fractured specimens obtained in the present study, 13 more detailed microstructure analysis was performed using the FESEM-EBSD technique. 14 The specimen observed in Figure 4 was analyzed. 15 Observations were conducted at a position of D = 3.75 mm, where the area fraction of the 16 -phase is less than that at D =3.5 mm (see Fig. 4e). The presence of thin lamellar -phase is 17 clearly observable within  grains (Fig. 5a). Figures 5b, c, and d represent EBSD inverse 18 pole figure maps of -phase, -phase, and both phases, respectively. In each  grain with the 19 17  lamellar texture, the thin layered  grains that are arranged parallel to each other appear to be 1 oriented to the same crystallographic orientation. This result strongly indicates the presence 2 of an orientation relationship between the original  phase and the resultant  phase. Figure 3 5e represents an EBSD grain boundary map at the same position. Grain boundaries between 4 two grains with the same crystal structure (i.e., / and / grain boundaries) are 5 represented by black lines, whereas those between two different phases (i.e., /a grain 6 boundaries) are represented by colored lines. The red and blue grain boundaries satisfy the 7 orientation relationships, OR I and OR II, respectively. It has been demonstrated that the 8 majority of the long and straight / boundaries are explained by OR I, and that the OR II 9 grain boundaries are significantly shorter than those of OR I. Utilizing this image, the 10 boundary line lengths (BLLs) of these three types of grain boundaries were calculated. In the 11 specific position (D = 3.75 mm), characterized by an area fraction of 11% for the -phase, 12 the OR I grain boundary occupies 73% of the total / phase boundary. In contrast, the OR II 13 grain boundary only occupies 11%. This phenomenon can be attributed to the energetic 14 favorability of OR I (ref. 35). The same grain-boundary analysis was subsequently conducted 15 at other locations within the fractured specimen: D = 4.0, 4.5 and 4.65 mm (Figs. 5f, g, and h). 16 The results are summarized in Figure 5i. It has been demonstrated that the ratio of OR I grain 17 boundary decreases with D, while the ratio of OR II grain boundary increases slightly with D. 18 Since the OR I grain boundary is primarily observable in relatively larger  grains (Fig. 5a), 19 18  this tendency with D (i.e., the distance from the edge of grip section toward the fracture 1 surface along the tensile axis, see Fig. 3a) may be associated with the increase of -phase 2 fraction (i.e., the decrease of -phase fraction) and the grain size reduction observed in 3 Figure 4. Moreover, transmission electron microscope (TEM) observations were performed 4 on the fractured specimen. The presence of an orientation relationship that is identical to OR I 5 was also observed between the  and  phases (Supplementary Figure 5). The data obtained 6 by FESEM-EBSD and TEM demonstrate that stress-induced  →  martensitic 7 transformation occurs during tensile deformation. 8 Further TEM studies, especially high-resolution observations, are needed to elucidate the 9 phase transformation process from the  to  phase under deformation. Figure 5a clearly 10 shows the presence of thin -phase layers in  grains. During the initial stage of deformation, 11 -phase nucleation may occur on a nanoscale at the grain boundaries of  grains and/or 12 within -phase grains. Understanding the phase transformation process can be beneficial to 13 controlling the mechanical properties of the -Ti materials. 14 19   1  2  3  4  5  6 In order to investigate the process of this stress-induced transformation, time-resolved 7 synchrotron XRD measurements of the -Ti material with CP-Ti Grade 2 composition were 8 conducted during a tensile test. Figure 6 summarizes the results obtained from the 9 Figure 5. Crystal orientation maps by EBSD at a position of D = 3.75 mm (a-d) and grain boundary maps at different positions (e-h) within the fractured -Ti tensile specimen. a) A phase map illustrating the distribution of the  and  phases at a position of D = 3.75 mm. The presence of thin lamellar -phase is clearly observable within  grains. b-d) Inverse pole figure maps of , , and both phases at a position of D =3.75 mm, respectively. e-h) Grain boundary maps at D = 3.75, 4.0, 4.5, and 4.65 mm, respectively. The grain boundaries between the  and  phases are represented by colored lines. The red and blue lines represent the grain boundaries that satisfy the orientation relationships, OR I and OR II, respectively. i) The ratios of / boundary line length (BLL) that satisfy the orientation relationship of OR I at different D positions are represented by red symbols. The ratios of / BLL that satisfy OR II at different D positions are represented by blue symbols. The former exhibits a decrease, while the latter displays a slight increase with D.  20  measurements. A miniature tensile specimen was mounted on a portable mechanical testing 1 device and pulled to failure at a constant cross-head displacement rate (0.5 m/s). XRD 2 profiles were collected at 2-second intervals, and the complete set of profiles obtained until 3 failure are shown in Figure 6a. The initial profile, which indicates the presence of a single 4 phase , is depicted in Figure 6b. The final profile, recorded just before the failure of the 5 tensile specimen, depicts the coexistence of the  and  phases (Fig. 6c). The relationship 6 between nominal stress and elapsed time (t) is illustrated in Figure 6d. Since the crosshead 7 speed is constant, this diagram corresponds to a nominal stress-strain curve of the -Ti 8 material. It is evident that the specimen undergoes elastic deformation up to t ~ 500 s, and no 9 transformation to the  phase occurs in the elastic regime. At t ~ 600 s, the XRD peaks of the 10  phase appeared, and the peak intensities increased with increasing elapsed time. The 11 volume fraction of the  phase was calculated using the XRD data with the reference 12 intensity ratio (RIR) method. The results are shown in Figure 6e. It is observed that, 13 following the appearance of the  phase, the volume fraction undergoes a rapid increase up to 14 t ~ 900 s within the region of uniform plastic deformation (Fig. 6d). Subsequent to reaching 15 the ultimate tensile strength (t ~ 1100 s), the volume fraction remains constant at 16 approximately 65% until fracture. As demonstrated in Figures 3b and 4e, the  phase 17 undergoes complete transformation into the  phase at the fracture position. This discrepancy 18 can be attributed to the misalignment between the X-ray beam position on the tensile 19 21  specimen and the fracture position. The data presented in Figures 3-6 demonstrate that stress-1 induced  →  martensitic phase transformation occurs exclusively in the plastic regime. 2 The present study revealed transformation-induced plasticity37 (TRIP) in the -Ti materials. 3  4  5  6 Figure 6. Results of time-resolved synchrotron XRD measurements of an -Ti material during a tensile test. The -Ti material was fabricated from CP2. a) All the XRD profiles obtained by the time-resolved data collection (every 2 s). The color contrast represents peak intensity. b) The profile recorded immediately after the crosshead was moved. c) The profile recorded instantly prior to the failure of the tensile specimen. d) Nominal stress versus elapsed time plot. This corresponds to the nominal stress-stain curve because the crosshead speed is constant (0.5 m/s). e) Calculated volume fraction of the  phase as a function of elapsed time. The stress-induced  →  transformation proceeds in the plastic regime. f) A schematic illustration of the experimental setup. 22  The tensile behavior of the -Ti materials may be explained by that of the mixture of the 1  and  phases. However, it should be noted that, as the plastic deformation proceeds, the 2 volume fraction of the -phase increases (Fig. 6a) and grain sizes of the both phases decrease 3 (Figs. 4 and 5). Moreover, measuring the mechanical properties of the -phase itself under 4 ambient conditions poses a substantial challenge, owing to the inherent thermodynamic 5 metastability. Tensile testing at cryogenic temperatures may be advantageous in observing the 6 tensile behavior of the -phase, as it can suppress the  →  phase transformation1. 7 The results of the present experimental study are consistent with those of previous 8 theoretical studies. Zong et al.38 conducted molecular dynamics (MD) simulations in pristine 9 single crystal -titanium. They found a highly anisotropic deformation behavior, brittle 10 fracture and superplastic deformation depending on the crystallographic orientation. In a 11 subsequent study, Zahiri et al.36 performed large-scale MD simulations in the  phase with 12 initial defects, which are inevitably contained in real titanium materials, particularly in the -13 phase. The authors found that stress-induced  →  martensitic transformation causes 14 extensive plasticity under various loading directions with respect to the crystallographic 15 orientation. The findings of these studies36,38 demonstrated that the  phase undergoes elastic 16 deformation at the initial stage of deformation and that the  →  martensitic transformation 17 occurs exclusively within the plastic regime. In the present experimental study, the majority 18 of the theoretically predicted phenomena that occur during deformation of -Ti were 19 23  successfully observed. A notable absence from the present study was the observation of phase 1 transformation sequence from  to  phase. Zahiri et al36 predicted that the  →  phase 2 transformation following OR I invariably involves a thin layer of the -phase at the / 3 interface. Further TEM observations are necessary to elucidate the nature of this 4 transformation sequence. 5 CP-Ti and Ti-6Al-4V are frequently utilized in biomedical applications due to their 6 exceptional corrosion resistance39. However, it should be noted that both materials possess 7 certain limitations40. CP-Ti exhibits comparatively lower mechanical strength, while Ti-6Al-8 4V has the potential to induce toxicity through the release of ions (Al and V). Since the 1960s, 9 CP-Ti has been the preferred material for dental implant devices to prevent cytotoxicity41. 10 The present study demonstrates that the -Ti material with CP-Ti Grade 4 composition 11 exhibits equivalent strength to that of Ti-6Al-4V (Table 1). This finding indicates that the  12 phase in pure titanium may have potential applications as a biomaterial, for instance in dental 13 and spinal implants. The role of passive film on titanium and its alloys is crucial for their 14 excellent biocompatibility42. Therefore, investigation of the passive film on the -Ti 15 materials may be beneficial for the evaluation of the biocompatibility. 16  17  18  19 24  Methods 1 Synthesis under high-pressure and temperature. High pressure and high temperature 2 synthesis runs were conducted using Kawai-type apparatus43, with tungsten carbide cubes 3 serving as the second stage anvils. The pressure transmitting medium was octahedral-shaped 4 semi-sintered magnesia doped with Cr2O3. A cylindrical graphite furnace was employed, with 5 a sample container made of hBN and MgO embedded into the furnace. The initial 6 compression of the sample was conducted at pressures ranging from 9 to 12 GPa at ambient 7 temperature. Subsequently, the temperature was elevated to a range of 400 to 500˚C. The 8 grain size of the -Ti materials was controlled through the implementation of various time-9 temperature profiles. Further details can be found in the supplementary information. 10  11 XRD measurements and microstructure observations of -Ti materials. The rod-shaped -12 Ti materials after high-PT treatment (see inset of Fig.1d) were chemically polished using a 13 commercially available liquid containing a nitric hydrofluoric acid mixture19. XRD 14 measurements on the flat surfaces of the rod-shaped specimens were performed by the 15 reflection geometry using a benchtop X-ray diffractometer (Miniflex600, Rigaku, Tokyo, 16 Japan) with Cu anode. For each -Ti material, a cross-section parallel to the rod-axis was 17 prepared and chemically polished to observe the microstructure. We used a FESEM 18 (Gemini450, Zeiss AG, Germany) equipped with an EBSD detector (Symmetry, Oxford 19 25  Instruments, UK). Typical operating parameters were 15 keV accelerating voltage and 10 nA 1 beam current. The step size utilized for the EBSD analysis was set at 100 nm.  2  3 Miniature tensile tests. Tensile specimens were extracted from the rod-shaped -Ti materials 4 (inset of Fig. 1d) through the implementation of wire EDM. The tensile axis of the specimen 5 is parallel to the rod axis. The fabrication process commonly yielded two specimens from a 6 single rod. After wire EDM cutting, the specimens underwent a chemical polishing procedure 7 to remove the damaged surface layer19. The total length and width of the specimen are 10 and 8 4.5 mm, respectively (Fig. 2d). The gauge length and gauge section are 1.5 mm and 0.7 × 0.7 9 mm2, respectively. Each tensile specimen was mounted horizontally on grips and pulled to 10 failure using a testing machine44 operating at a constant rate of cross-head displacement. The 11 initial strain rate of each test was 5.5 × 10-4 s-1. The stress and strain of the specimens were 12 calculated by utilizing the applied load and cross-head displacement, respectively, on the 13 assumption that the gauge section deforms uniformly and that its volume is constant 14 throughout the tensile tests. It is noteworthy that, in the present study, all the tensile tests 15 were conducted under identical testing conditions, thereby ensuring internal consistency. It is 16 also important to note that the present tensile specimens are relatively small. Therefore, when 17 comparing the present data with other results obtained by using tensile specimens with 18 different dimensions, caution should be taken45,46. Nevertheless, it should be noted that the 19 26  tensile strength and elongation to failure of Ti-6Al-4V determined in the present study (Table 1 1 and Supplementary Fig. 2) are reasonably consistent with those reported in a previous study 2 conducted by using larger tensile specimens47. 3  4 High energy XRD measurements for fractured specimens. The measurements were 5 conducted at BL16, SAGA Light Source, Japan. Synchrotron X-rays from a 4T 6 superconducting wiggler were monochromatized by a Si (111) double-crystal monochromator. 7 The higher harmonics were subsequently removed by utilizing a Pt-coated mirror. The X-ray 8 energy of the incident beam was 18 keV. The X-rays were then collimated by slits into beams 9 with a width of 1.0 mm and a height of 0.5 mm, which were irradiated on the specimens. The 10 diffracted X-rays were recorded by a 2D detector (Pilatus 100K-S, Dectris AG, Switzerland). 11 The X-ray energy and the sample-to-detector distance (152.5 mm) were calibrated using a 12 CeO2 standard (NIST SRM 674a). The typical exposure time was 10 seconds. 13  14 Microstructure observations for a fractured specimen. A fractured specimen was mounted 15 in epoxy resin and ground to expose a cross section parallel to the tensile axis and 16 perpendicular to the flat surface of the tensile specimen (Fig. 4a). The cross section was then 17 subjected to a polishing procedure using an Ar ion beam (IB-09020CP, JEOL Ltd., Japan). 18 27  The microstructure was then examined by utilizing the FESEM-EBSD technique. The 1 instruments and operating parameters were identical to those employed for the -Ti materials. 2  3 Time-resolved XRD measurements during a tensile test. The measurements were conducted 4 at BL16, SAGA Light Source, Japan. The X-ray energy of the incident beam was 22 keV. The 5 X-rays were collimated by slits into 0.3 mm in width and 0.5 mm in height and irradiated on 6 the gauge section of the tensile specimen. A portable tensile testing device (TST350, Linkam 7 Scientific Instruments Ltd., UK) was employed, which is mounted on motorized two-axis 8 (perpendicular to the X-ray beam) translation stages to adjust the sample position48. The 9 experimental setup for the XRD measurements is similar to that for fractured specimens, with 10 the sample-to-detector distance set at 175.6 mm. 11  12 Data availability 13 The data that support the findings of this study are available from the corresponding author 14 upon reasonable request. 15  16 References 28  1. Hennig, R. G., Lenosky, T. J., Trinkle, D. R., Rudin, S. P. & Wilkins J. W. Classical potential describes martensitic phase transformations between , , and  titanium phases. Phys. Rev. B 78, 054121 (2008). 2. Jamieson, J. C. Crystal structures of titanium, zirconium, and hafnium at high pressures. 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Processing pure Ti by high-pressure torsion in wide ranges of pressures and strain. Metall. Mater. Trans. A 40A, 2079-2086 (2009). 46. Zhao, Y. H. Influence of specimen dimensions on the tensile behavior of ultrafine-grained Cu. Scripta Mater. 59, 627-630 (2008). 47. Hu, H., Xu, Z., Dou, W. & Huang, F. Effect of strain rate and stress state on mechanical properties of Ti-6Al-4V alloy. Int. J. Impact Eng. 145, 103689 (2020). 48. Tokuda, K., Goto, K., Iihara, J. & Adachi, H. Visualization of local deformation in metallic materials through in situ X-ray diffraction mapping. Mater. Trans. 66, 50-55 (2025).  Acknowledgements 1 34  We thank T. Hanawa, H. Hosoda, A. Holtzheid, T. Nakano, P. Wiedekehr, M. Pawelkiewicz-1 Koebel, H. Fujii, and M. Goto for discussion. We also thank H. Ichihara, K. Manabe, I. Sou, J. 2 Li, H. Muto, and H. Ikeda for technical assistance. Synchrotron XRD measurements were 3 conducted at BL16, SAGA-LS (project number: SEI2023C-012, SEI2024B-012). This 4 research was supported by Japan Society for the Promotion of Science, a Grant-in-Aid for 5 Scientific Research on the innovation area “Science of New-Class of Materials Based on 6 Elemental Multiplicity and Heterogeneity (Grant No. 18H05452)” partially to N.N. 7  8 Funding 9 Open access funding provided by Sumitomo Electric Industries, Ltd. 10  11 Author Contributions 12 N.N., K.S., F.W., and Z.H. designed the research. N.N., M.M., Y.K., A.F., and Z.H. led the 13 project. N.N., Y.T., and T.Sawahata performed high-pressure synthesis. Y.T. and T.Sawahata 14 performed XRD measurements. Y.T., T.Sawahata, and Z.H. performed tensile tests. K. 15 Tominaga and T.Sekiya performed FESEM-EBSD measurements. K.K. performed TEM 16 observations. N.N. and K.Tokuda performed synchrotron XRD measurements. N.N., M.M., 17 and Z.H. wrote the manuscript with contributions from other authors. 18  19 35  Figure 1.  The characterization of polycrystalline -Ti materials with pure titanium 1 compositions fabricated under high pressure and temperature conditions is presented. Four 2 different starting materials (all  phase) were employed for the high-PT treatment: Ti 99.999 3 wt% metals basis (5N); CP-Ti Grade2 (CP2); CP-Ti Grade 4 (CP4); cold-worked CP-Ti 4 Grade4 (CP4-cw). The X-ray diffraction patterns of the -Ti materials fabricated from 5N, 5 CP2, CP4, and CP4-cw are shown in a) – d), respectively. It is evident that all the materials 6 consist solely of the  phase. The inset of d) presents a photograph of the titanium materials: 7 the bottom, a starting material (CP4-cw); the middle, an -Ti material after high-PT 8 treatment; the top, an -Ti material fabricated from CP4-cw after surface machining by a 9 centerless grinder. The diameter of this object is 4.9 mm and its length is 17.7 mm. The 10 samples were placed on a 5 mm grid paper. e) – h) were inverse pole figure (IPF) maps 11 superimposed on band contrast (BC) images of the -Ti materials fabricated from 5N, CP2, 12 CP4, and CP4-cw, respectively. The inset of e) shows the color coding for the orientations in 13 e) – h). It is clear that grain size decreases from e) to h) (see text). 14  15 Figure 2.  Stress-strain curves of the starting () and -Ti materials with pure titanium 16 compositions obtained by miniature tensile tests. a) 5N (blue) and the -Ti material 17 fabricated from 5N (red); b) CP2 (blue) and the -Ti material (red); c) CP4 (blue) and the -18 Ti material (red); d) CP4-cw (blue) and the -Ti material (red). The inset of d) shows a 19 36  photograph of a miniature tensile specimen. In Fig. 2a, “” refers to a stress-strain curve 1 obtained from a bulk polycrystalline -titanium material; “” refers to a stress-strain curve 2 obtained from a bulk polycrystalline -titanium material (see text). 3  4 Figure 3. X-ray diffraction profiles obtained from a fractured -Ti tensile specimen. The -5 Ti material was fabricated from CP2. The fractured specimens were placed on a 5 mm grid 6 paper. a) A photograph of fractured tensile specimens and definition of the distance from the 7 edge of grip section, D. The beam size (1.0 mm × 0.5 mm) is shown by the dashed squares. b) 8 X-ray diffraction profiles collected at different D positions from D = 2.0 mm (in the grip 9 section) to D = 5.0 mm (the position including fracture surface). XRD peaks of the  phase 10 can be seen in the gauge section between D = 3.5 and 5.0 mm. 11  12 Figure 4. Microstructure observations of a fractured -Ti tensile specimen. The -Ti 13 material was fabricated from CP2. a) Photograph of a fractured specimen. The dashed line 14 indicates the position of the cross section for the observations. The cross section is 15 perpendicular to the flat surface of the tensile specimen. b) Overview back-scattered electron 16 image (BEI) of the investigated area. Detailed observations were conducted at positions 17 where  →  transformations occur: D = 3.5, 4.0, 4.5 and 5.0 mm. BEIs (c), band contrast 18 (BC) images (d), and phase maps (e) are shown at these positions. In c), the red arrows 19 37  indicate the locations of pores. At D = 4.0mm, the lamellar  phase was observed within the 1  phase matrix. 2  3 Figure 5. Crystal orientation maps by EBSD at a position of D = 3.75 mm (a-d) and grain 4 boundary maps at different positions (e-h) within the fractured -Ti tensile specimen. a) A 5 phase map illustrating the distribution of the  and  phases at a position of D = 3.75 mm. 6 The presence of thin lamellar -phase is clearly observable within  grains. b-d) Inverse pole 7 figure maps of , , and both phases at a position of D =3.75 mm, respectively. e-h) Grain 8 boundary maps at D = 3.75, 4.0, 4.5, and 4.65 mm, respectively. The grain boundaries 9 between the  and  phases are represented by colored lines. The red and blue lines represent 10 the grain boundaries that satisfy the orientation relationships, OR I and OR II, respectively. i) 11 The ratios of / boundary line length (BLL) that satisfy the orientation relationship of OR I 12 at different D positions are represented by red symbols. The ratios of / BLL that satisfy 13 OR II at different D positions are represented by blue symbols. The former exhibits a 14 decrease, while the latter displays a slight increase with D. 15  16 Figure 6. Results of time-resolved synchrotron XRD measurements of an -Ti material 17 during a tensile test. The -Ti material was fabricated from CP2. a) All the XRD profiles 18 obtained by the time-resolved data collection (every 2 s). The color contrast represents peak 19 38  intensity. b) The profile recorded immediately after the crosshead was moved. c) The profile 1 recorded instantly prior to the failure of the tensile specimen. d) Nominal stress versus 2 elapsed time plot. This corresponds to the nominal stress-stain curve because the crosshead 3 speed is constant (0.5 m/s). e) Calculated volume fraction of the  phase as a function of 4 elapsed time. The stress-induced  →  transformation proceeds in the plastic regime. f) A 5 schematic illustration of the experimental setup. 6 39   1 40   1   2 41  Table 1. Summary of mechanical properties of pure titanium materials determined by miniature tensile tests. 1 Phase   +    Material 5N  CP2  CP4  CP4-cw  Ti64   5N  CP2  CP4  CP4-cw Grain size (m) -   -   -   -    -    65(16) 12(4) 5.2(17) 3.4(11) 0.2% (MPa) 118(18) 257(15) 640(50) 973(5)  990(40)  264(9) 634(13) 990(20) 1130(30) TS (MPa) 284(11) 574(14) 930(20) 1121(7)  1190(70)  527(6) 833(11) 1090(30) 1220(30) f (%) 64(8) 40(2) 28(2) 22(1)  25(3)  38(2) 34(1) 30(2) 16(2) The averages of the three tensile test runs for each material are shown with one standard deviation in parentheses. 2 Phase: a phase or phases that are present in the tensile specimens before starting the tensile tests. 3 5N: 99.999% Ti; CP2: CP-Ti Grade2; CP4: CP-Ti Grade4; CP4-cw; cold-worked CP-Ti Grade4; Ti64: Ti-6Al-4V. 4 0.2%: 0.2% offset yield strength; TS: tensile strength; f: elongation to failure. 5   6 42   1   2 43   1   2 44   1   2 45   1