Promoting direct and ultrafast precipitation of the α phase in the metastable Ti-10V-6Cu alloy

Direct aging of water-quenched metastable β-titanium alloys usually follows the sequence: rapid growth of ωiso and then slow transformation into α. To inhibit ωiso growth and promote α nucleation, 6% pre-deformation was introduced in Ti-10V-6Cu alloy before 500°C/60s aging. An ultrafast α-formation and a few α-free kink-bands are observed, indicating the alteration of the precipitation sequence. For α-nucleation without ωiso as the sites, it shows that O' nanodomains induced by stress accumulation around slip bands mediate the ultrafast nucleation of hcp-α structure, while dislocation pipe diffusion accelerates the necessary phase separation. This provides new insights into regulating precipitation kinetics of age-hardenable β-titanium alloys. GRAPHICAL ABSTRACT


Introduction
β-Ti alloys are promising materials for aerospace applications due to their outstanding comprehensive mechanical properties [1,2].Upon aging, α phase (a hexagonal close-packed structure, hcp) precipitates within the β-Ti matrix (a body-centered cubic structure, bcc), tailoring strength, ductility and toughness of β-Ti alloys.Thus, it is of great importance to control the size, morphology and distribution of α-phase precipitates, which depends on the nucleation mechanism [3].Besides conventional homogeneous nucleation, the nucleation of the α phase can also occur heterogeneously at grain boundaries, dislocations or various intermediate domains such as ω, β', O' or O'' [3][4][5][6][7].Over the past several decades, great efforts have been attempted with the focus on optimizing the chemical composition and/or heat-treatment regime to produce a desirable degree of finely dispersed α microstructure [8][9][10][11].
ω-assisted nucleation of α phase has been well documented in many metastable β-Ti alloys.The phase transformation pathway during artificial aging (AA) follows a specific sequence of β+ω ath →β+ω iso →β+ω iso +α, but the detail of nucleation mechanisms varies over a wide range of compositions and processing windows [12].More specifically, the ω ath phase formed via atomic shuffling in water-quenched samples will transform to metastable ω iso and further grow up via a diffusionmediated process upon aging at temperatures below ω solvus.In this case, the activation energy barrier for subsequent α-nucleation can be reduced by (1) the stress field (structural) and ( 2) the concentration field (compositional) associated with the growth of ω iso particles [13].The nucleation of α-phase either at ω/β interface within the core or in the vicinity of ω iso particles has been experimentally evidenced [13][14][15].Therefore, it is essentially required to inhibit the growth of ω iso to provide more heterogeneous nucleation sites of α precipitates.In a two-step aging regime, lower-temperature aging at the first step induced the moderate growth of ω iso , which could favor the refinement of α-platelets and improve mechanical response during subsequent higher-temperature aging [12,16].However, considering the diffusion nature of various α and β stabilizers during α-nucleation, an ultralong period for compositional redistribution is still required, resulting in the slow age-hardening kinetics regarding metastable β-Ti alloys [17,18].Obviously, there is a trade-off between inhibiting ω iso -growth and adequate phase separation when applying conventional processing regimes.Although it was recently reported in the Ti-19 at.%V alloy that α can nucleate in the absence of influence from ω-phase when the applied aging temperature is greater than ω-solvus, the size of α precipitates is distinctly less refined compared with that from ω-assisted α nucleation mechanism due to the decreased nucleation sites [19].Therefore, alternative methods to promote α precipitation in thermodynamics and kinetics still need to be investigated.
It has been demonstrated that direct aging at 500°C for 60 s of water-quenched Ti-V-Cu alloy results in the rapid growth of the ω iso phase and no nucleation of the α phase [20].To inhibit the rapid growth of the ω iso phase and promote the nucleation of the α phase during aging, a ternary Ti-10V-6Cu (wt.%) alloy was selected in this work for solution-treatment and water quenching, followed by 6% pre-deformation to introduce deformation defects before aging at 500°C for 60 s.The design philosophy of this alloy is to reside in the 'slip' zone and be well removed from the twin-slip boundary in the metal d-orbital level and the bond order (M d -B o ) map [21], as seen in Figure S1 in the supplementary section.Dislocation slip, as the dominant mode during deformation, makes it possible to decouple the multiple deformation mechanisms and introduce kink bands (KBs) at coarse grains [22,23], to locate the preferred site of α-phase nucleation.Besides, as a deep-hardenable alloy, Ti-10V-6Cu is insensitive to cooling rate [17] and exhibits a phase separation process before the commencement of ω→α transition, thus providing favorable conditions for observing α-nucleation [20].The results show that a high density of fine α platelet is generated and a decreased fraction of ω phase under the same aging condition.The ultrafast α-precipitation behavior and an altered precipitation sequence achieved in this work are investigated to address the promoted nucleation mechanism of the α phase without the need for ω iso as the site.

Materials and methods
Ti-10V-6Cu (wt.%) alloy was prepared by arc-melt casting.Cylinders with a size of 6×9 mm were cut from the as-cast ingot and were encapsulated in an argonfilled quartz tube.Solution treatment was carried out at 1100°C for 48 h and then water was quenched.After that, one set of cylinder samples (diameter 6 mm, height 10 mm) were subjected to 6% compression at a fixed rate of 0.005 s −1 using an Instron universe machine.Subsequently, short-term isothermal aging at 500°C for 60s was conducted by soaking the pre-deformed sample in a hot oil bath furnace (without encapsulation).For reference, other solutionized samples were directly aged at the same condition.Differential scanning calorimetry (DSC) was conducted on Netzsch 404F3 at 5°C/min.Foils for transmission-electron-microscopy (TEM) observations were thinned by twin-jet electro-polishing.Talos F200X G2 was operated at 200 kV for the TEM mode.High-angle-annular-dark-field scanning TEM (HAADF-STEM) mode was performed on a Titan Themis-Z aberration-corrected microscope equipped with energydispersive spectroscopy (EDS) at 300 kV.Fast Fourier transformation (FFT), inverse fast Fourier transformation (IFFT) and geometric phase analysis (GPA) were obtained by GMS 3. A JEOL 2100F microscope, equipped with a NanoMEGAS DigiSTAR scan generator, was used for conducting scanning-precession-electron-diffraction (SPED) experiments to construct the phase and orientation maps [24].The precession angle employed was 0.5 o and the step size was set as 1 nm.

Results and discussion
Figure 1 shows the DSC curves of the solutionized and pre-strained sample during heating to explore the microstructural evolution.For the solutionized sample, a wide endothermal change is found at a low temperature range (350-480°C), indicating the formation and growth of isothermal ω iso .The corresponding peak is inconspicuous due to the pre-formed ω ath phase.Above ∼ 480°C, a large exothermic peak appears with a distinct slope, which is attributed to nucleation of α phase by transforming from ω domains.With further increasing temperature, the α phase continuously grows.In comparison, the pre-strained sample presents a different transformation sequence between 350 and -600°C.The purple curve in Figure 1 keeps decreasing until ∼ 410°C without obvious fluctuation.In this case, the potential ω ath →ω iso transition is not detected.In the range of 410-560°C, it is interesting to find a wide endothermal peak with the absence of any abrupt exothermal peak, suggesting the promotion of α-nucleation.Speculation is that the deformed structure could facilitate the ω→α transition at a wider and lower temperature range (410-560°C), followed by the growth of α during further heating.
TEM characterizations are performed on the Ti-10V-6Cu alloy at various states to explore the transformation sequences, as displayed in Figure 2. In the solutiontreated sample, ω ath domains with very fine-scaled size ( ∼ 2-4 nm) distribute homogeneously in the β-matrix as the dark-field (DF) image shows in Figure 2(a).The reflections at the 1/3 and 2/3112 positions in the selected area diffraction pattern (SADP) along [110] β arise from the ω phase.After 6% pre-compression, besides the stress-induced ω phase (Figure S2), the most important feature is that several band-like structures are generated in Figure 2(b), with the thickness varying from ∼ 20 nm to ∼ 200 nm.Most of the bands are parallel to each other and align along < 112 > β .EDS mapping (Figure S3) shows that the bands own the same composition as the matrix.In addition, there is a distinct discrepancy between the bright band and the dark matrix, indicating the deflection of the bands away from the Bragg diffraction condition.From SADP in Figure 2(c), the motifs of the two ω-variants are still visible.The corresponding DF images in Figure 2(d and e) clearly show that these bands are ω-devoid, as similarly observed by Lai.et al. in Ti-25Nb-0.7Ta-2Zralloy [25], where the ω-free deformation bands were delineated as channels for dislocation moving.Herein, these bands could be intrinsically defined as KBs which locally rotated in a more favorable Schmid-factor direction to produce local crystallographic re-orientation [23,26].As indicated in Figure 2(c), an extra set of bcc spots marked by orange rectangles can be identified, which is slightly off the zone axis and has a low misorientation angle (rotation axis is [-110]).This provides strong evidence to indicate the BCC structure of KBs.In this case, most of the KBs belong to the < 110 > -type, as reported in [27,28], and the corresponding slip system is identified as (112) .It should be mentioned that the coarse grains induced by homogenization (1100°C/48 h) in the present work facilitate the formation of KBs, which is similar to the findings in [22,29].
After short-term aging at 500°C for 60s, precipitation behaviors are observed for comparison in Ti-10V-6Cu samples with and without pre-strain.Figure 2(f) is a DF image from the directly aged sample acquired using ω reflections in SADP.A dense ω iso phase can be found, which is expected to be transformed from the waterquenched ω ath domains.This is consistent with the findings by Ng. et al. that aging at 500°C for 60 s results in no nucleation of α phases but facilitates the rapid growth of ω phase in Ti-10V-6Cu alloy [17,20].Surprisingly, in this work, the appearance of α phase reflections is found in the SADP of the pre-strained and aged sample, as marked by the yellow arrows in Figure 2(g).In addition, the ω spots become weak, suggesting a decreased fraction of the ω-phase after aging.Further microstructure observations in Figure 2(h and i) confirm the presence of ultrafine acicular α-precipitates (length: 10-20 nm), which coexist with minor ω-phase presented in the β-matrix.This result implies that the precipitation sequence is altered.With respect to the water-quenched sample, direct aging leads to the growth of ω phase, while small deformation applied before direct aging promotes the nucleation of α phase, which agrees well with the (β+ω)→(β+ω+α) transformation, as predicted in Figure 1.It should be noticed that the distribution of α is not homogeneous in the matrix.A series of precipitate-free zones (PFZs) are observed where the KBs reside, as marked in Figure 2(i).High-resolution TEM (HRTEM) analysis (Figure 3(a-c)) is performed to study the absence of αprecipitation within the KBs.In Figure 3(a), the band and the β-matrix are divided by a yellow dashed line.The former one is relatively 'clean' with only a set of bcc reflections, as shown in FFT-1, while the latter one owns several 'wrinkled' areas and consists of both α and ω phases, as indicated in FFT-2.A small rotation angle of ∼ 1.9 o along the [-110] β axis is identified between the two regions.The IFFT image obtained through all masked α-reflections shown in Figure 3(b) further proves that α only precipitates in the β-matrix instead of the KBs, and more specifically, located at the 'wrinkled' areas where dislocations accumulate [30].This can be evidenced by the Geometric Phase Analysis (GPA) result in Figure 3(c), where the KBs exhibits much relaxed stress distribution in contrast to the β-matrix.Different from the ω iso →α transformation mechanism during the direct aging process; however, the nucleation of α-precipitation in this work can only be initiated at regions with high residual stress.
The SPED technique was also performed to provide statistical results of α-precipitates and misorientations between the rotated KBs and the matrix.As seen in the phase map in Figure 3(d), a representative area containing a KB in the center was indexed using a step size of 1 nm.Again, there is no preferential nucleation of the α phase in the KB.In such a case, the area fraction of the αphase is 21%.From the corresponding inverse pole figure map in Figure 3(e), it can be found that the orientation within the KB is uneven.Reversely, it produces local variation in a wide range of ∼ 20 nm, as indicated by the 1D misorientation profile in Figure S4.The misorientation angle of KB is in the range of 1-5°(Figure S4) for the β matrix.The schematic figure of this configuration is given in Figure 3(f).
The key question is to elucidate the underlying mechanism of altering precipitation sequence and ultrafast formation of the α phase achieved in this work.To bridge the relation between the defects and the promoted nucleation of the α phase without the need for the ω iso phase as the site, Figure 4(a) presents an atomic-scale HAADF-STEM image of a relatively large α-precipitate.It can be seen that α owns the hcp structure and shows a darker contrast tn the β-matrix.The classic Burgers orientation relationship between α and β, (0001) α //(110) β and [-12-10] α // [1][2][3][4][5][6][7][8][9][10][11] β , can be identified as extensively reported in [3,9,31,32].In addition, some spherical particles with much brighter contrast can be observed with the same bcc structure either in the matrix or at the α/β interfaces, which are identified as Cu-rich clusters as shown in the corresponding EDS maps in Figure 4(c-e).In addition, the α phase is enriched in Ti and depleted in V and Cu.It proves that the modulation of composition (phase separation) [20] and structure transition synergistically occur, resulting in α-precipitation (see Figure S5).A specific check of the α structure in Figure 4(b), using IFFT of all α-reflections, demonstrates the presence of slip bands at the α interface and local ordering change in the center region.An enlargement of the white rectangle region is thus displayed in Figure 4(f).There is one dislocation found in the center, as shown in Figure 4(g), and strong lattice distortion around the dislocation can be reflected by the slight deviation of the nearby (1-100) α planes from the perfect lattice positions.High residual stress locally around dislocations and interfaces facilitates α-phase precipitation [33].Besides, in the vicinity of the dislocation within the α-phase, the solute redistribution is more significant compared with other regions.This is probably ascribed to the faster pipe diffusion with the assistance of dislocation since it can serve as an efficient diffusion path for solute atoms [34].As a result, the phase separation process was accelerated.To verify the nucleation pathway arising from slip bands/dislocations, the structure of a very fine α-embryo with a length of ∼ 10 nm was captured in Figure 5   which are known to form via 100 < 1-10 > shuffle in bcc structure [35], can be observed on each side of the acicular α-embryo.Further atomic intensity profiles of the marked layers in Figure 5(c) confirm the atomic displacement of ∼ 0.3Å from the parent β positions in < 110 > β directions on every alternate 110 β plane.In this case, the ω phase is absent, different from the ω-assisted αnucleation theory [36].It has been demonstrated that ω dissolved during aging while the O phase first developed in β and then transformed to fine α [8].The deformationinduced defects assist subtle shuffle of atoms and quickly cause the formation of O' domains, which transform into the hcp-α phase during aging.For the α-embryo at an early stage, the associated atomic shear and/or element diffusion accompanied by the transition is expected by using a combination of STEM and atom probe tomography techniques.

Conclusion
The Ti-10V-6Cu alloy suffers from very slow precipitation kinetics with a specific sequence from ω ath to ω iso and finally to α.After direct aging at 500°C for 60 s, α-nucleation does not occur but a rapid growth of ω iso can be found.To inhibit the rapid growth of the ω iso phase and promote α-precipitation, we introduced 6% pre-deformation before aging for water-quenched Ti-10V-6Cu alloy with the focus on the very early-stage of α-nucleation and the corresponding mechanism.Surprisingly, a decreased fraction of the ω-phase with the formation of ultrafine and dense acicular α-precipitates is found also after aging at 500°C for 60 s.Apart from precipitate-free KBs, this ultrafast and direct α-precipitation happens in the βmatrix, which implies the altered precipitation sequence from (β+ω ath )→(β+ω iso ) to (β+ω ath )→(β+ω iso +α).Atomic-scale characterizations reveal that O' domains assisted by local stress accumulation around deformationinduced slip bands mediate ultrafast nucleation of the hcp-α structure, while dislocation pipe diffusion accelerates the necessary phase separation.This work sheds light on tailoring the precipitation kinetics of age-hardenable titanium alloys for high performance.

Figure 1 .
Figure 1.DSC exothermic curves of the solutionized and prestrained Ti-10V-6Cu samples during the continuous heating process.

Figure 2 .
Figure 2. TEM micrographs of the solutionized and pre-strained samples before and after 500 o C/60s aging.All images are taken along the < 110 > β zone axis except for (f).(a) DF image of the solutionized sample recorded using ω reflections from the SADP (see inset), showing uniformly distributed nanometer-scale ω domains.(b) BF images of the pre-strained sample with several bands observed paralleling to the [-11-2] β direction, while the SADP in (c) illustrates another set of extra reflections.DF images of two kinds of ω variants are shown in (d) and (e), respectively.(f) DF image of the coarsened ω phase together with SADP of the solution-treated sample after aging.(g) SADP of the pre-strained and aged sample, showing the presence of α-reflection spots as marked by the yellow arrows.(h, j) DF images of the same sample acquired from the white circle in (g) reveal the inhomogeneous distribution of nano α-precipitates, while obvious precipitation-free zones can be found distributing along < 112 > β directions.

Figure 3 .
Figure 3. Microstructural characterization of the PFZ (or KB) observed in the pre-strained and aged Ti-10V-6Cu sample.(a) HRTEM micrograph containing both the matrix and the kink band with the FFT patterns for each on the right side.(b) IFFT graph of (a) using all α reflections, showing α precipitates are absent in the kind band.(c) GPA analysis of (a) showing the planar stress distribution ( xy ).(d) The phase map and (e) inverse pole figure (IPF) map were acquired using the SPED technique, while the kink band is marked by yellow dashed lines.Note the 1D misorientation profile along the black solid line in (e) is provided in Figure S3.(f) Schematic figure showing the crystallographic relationship between the kink band and the β-matrix.
(a).The α-embryo nucleated around slip bands.An enlarged image shown in Figure 5(b) reveals that orthorhombic O' nano-domains,

Figure 4 .
Figure 4. Atomic-scale HAADF-STEM images viewing along the [110] β direction revealing the α-nucleation mechanism induced by prestrain.(a) Micrograph of a relatively large α-precipitate within the β-matrix with the interfaces marked by dashed yellow lines.The α phase exhibits dark contrast while some bright spherical domains are identified as Cu-rich clusters.Note the area in the green box was Fourier-filtered to provide a better image quality.(b) IFFT image obtained using all α-reflections showing the presence of slip bands at α interface.(c-e) EDS mapping results of (a) showing the distribution of elements.(f) Enlargement of white box in (b), and (g) the corresponding fringe image obtained through the IFFT of masked (1-100) α reflections.A dislocation within the α phase can be identified, while an extra (1-100) α plane is marked by the yellow solid line in the insert.

Figure 5 .
Figure 5. (a) Atomic-scale HAADF-STEM image showing a representative α-embryo observed heterogeneously nucleating on slip bands.(b) Fourier-filtered image with higher magnification showing the atomic structure of the α-embryo, while O' nano-domains can be observed between the α-embryo and β-matrix.Corresponding atomic intensity profiles of the marked layers from (b) are provided in (c), proving the presence of O' nano-domain.