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Vol. 8. Issue 1.
Pages 777-787 (January - March 2019)
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Vol. 8. Issue 1.
Pages 777-787 (January - March 2019)
Original Article
DOI: 10.1016/j.jmrt.2018.06.006
Open Access
Design of uniform nano α precipitates in a pre-deformed β-Ti alloy with high mechanical performance
Bingjie Zhanga, Tao Yangb, Mingda Huanga, Dong Wangb, Qiaoyan Suna,
, Yunzhi Wanga,b,c,**, Jun Suna
a State Key Laboratory for Mechanical Behavior of Materials, Xi’an Jiaotong University, Xi’an, Shaanxi 710049, China
b Center of Microstructure Science, Frontier Institute of Science and Technology, Xi’an Jiaotong University, Xi’an, Shaanxi 710049, China
c Department of Materials Science and Engineering, The Ohio State University, 1305 Kinnear Road, Columbus, OH 43212, USA
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Table 1. Element wet chemical analysis of the ingot.

A strategy for the microstructural design to achieve a critical upper limit of uniform nano α precipitates by controlling the amount of dislocations in the pre-deformed matrix after short time aging treatment was proposed in Ti–10Mo–8V–1Fe–3.5Al (all in wt.%, TB3 alloy) β alloy. The optimal processes focused on the interaction of defects, mainly point defects and dislocations that are generated during cold rolling (CR). Texture evolution of α and β phases was also specified. Amount of nano-scaled α precipitates was characterized by the number density of precipitates observed by SEM, which is saturated at ∼204.1μm−2. For the CR 20% sample, a good combination of tensile strength (∼1400MPa) and elongation (∼12%) was achieved after secondary aging at 550 for 1h, which provides an enlightenment for the re-engineering of traditional precipitate-hardening alloys.

β-Ti alloy
Short-time aging saturation
Uniform α precipitates
Mechanical properties
Full Text

Near-β Ti alloys, including Beta-C, Beta III, and Ti-8823, possess high specific strengths, high toughness, and good corrosion resistance and are excellent choices not only for structural components, but also as materials to replace Ni-based fasteners and steel fasteners in aerospace and automotive industries [1]. High strength of this type of age-hardenable alloys is obtained by aging treatment to precipitate fine, uniformly distributed incoherent α platelets of 12 orientation variants in metastable β matrix [2–6]. However, direct aging for such commercially available age-hardening alloys is supposed to induce non-uniform dispersion of coarse α precipitates and precipitate-free zones (PFZs) along grain boundaries (GBs). The occurrence of PFZs generally, explained in terms of either vacancy depletion or solute depletion, is considered by Thomas et al. [7] as deleterious layers due to preferential plastic deformation and fracture in the zones along GBs. Such unfavorable factors could induce limited strength and ductility of β-Ti alloys even after long aging duration.

Three main methodologies [2,3,8–10] for achieving finer and more uniform dispersion of α precipitates were explored by researchers in different β-Ti alloy systems: (1) ω-assisted α nucleation: the heterogeneous nucleation of the α phase from the precursory ω precipitates; (2) thermally induced compositional fluctuations within the β matrix via the pseudo-spinodal decomposition mechanism to assist α nucleation; (3) a high density of dislocations, grain, or sub-GBs by hot/cold deformation prior to aging to accelerate the kinetics of precipitation. Based on the third mechanism, Xu et al. [9,10] and other researchers [11–13] have studied different aging and mechanical responses to the corresponding pre-deformation degree as well as passes through cold/hot deformation. Furthermore, severe pre-deformation prior to a thermal treatment is often applied in the fastener process of β-Ti alloys to achieve the desired strength. The accumulation of dislocations during cold deformation leads to low ductility. The enhanced solutes transport is considered to induce undesirable consequences, e.g., creep in materials, which associates with edge-dislocation climb [14,15]. This raises a question: will the precipitate density continue to increase with increasing deformation degree? Alternatively, is there an optimum dislocation density, which could yield the saturation of precipitates and properties? The answer to these questions can help improving properties of the work-pieces with control of defect density by deformation [16].

Thus, the major objective of this study is to improve the precipitate dispersion and refine precipitates to achieve better comprehensive mechanical properties via cold deformation and subsequent heat treatment. The selected processes can avoid the disadvantages of other severe plastic deformation methods such as equal channel angular pressing (ECAP), in which the ECAP-dies have limited bearing capacity to process high-strength alloys. By selecting parameters of pre-deformation and subsequent aging process, a series of combination of processing variables were explored in the presented β-Ti alloy.

2Experimental procedure

In this paper, a β-Ti alloy was selected. It is the Ti–10Mo–8V–1Fe–3.5Al (wt.%), known as TB3 alloy with [Mo]eq=15–17, often used as fastener material due to its high strength and good formability. This β-Ti alloy was prepared by arcmelting of a mixture of high-purity elements (99.99%) under a Ti-gettered argon atmosphere, with subsequent cogging, intermediate forging, and hot rolling into cylindrical rods of 15mm in diameter. Subsequently, solution treatment (ST) was carried out at 830°C for 0.5h followed by water quenching. This procedure resulted in β grains with an average size of 50μm. The alloy composition is shown in Table 1. Cold rolling was conducted on 6×6×10mm plates at room temperature using the reduction of 10, 20, 30, 50, 70, 80, and 90%. Plate samples were then sealed in evacuated quartz tubes and aged at 500°C for 1–12h. The sample (CR20%/500°C for 1h) was subsequently annealing-treated at 550°C for 1h. The cooling method of all the aging treatment is water cooling. And the general processes were summarized in Fig. 1.

Table 1.

Element wet chemical analysis of the ingot.

Sample  Al (wt.%)  Mo (wt.%)  V (wt.%)  Fe (wt.%)  O (wt.%)  N (wt.%)  H (wt.%) 
Top  3.39  10.2  8.06  1.08  0.1  0.01  0.001 
Bottom  3.4  10.2  8.01  1.02  0.1  0.01  0.001 
Figure 1.

Plot showing the experimental design adopted in this study.


X-ray diffraction method (XRD; Maxima_X XRD-7000) and transmission electron microscope (TEM; JEM-2100) were used to characterize the microstructure of the processed samples. The samples were subjected to orientation imaging microscopy using a field emission scanning electron microscope (JSM-7100F) with electron back-scattered diffraction (EBSD) facility. Field-emission scanning electron microscopy (FESEM; HysitronSU6600) was conducted in the morphology characterization of precipitates and energy spectrum analysis. The Vickers microhardness (HV) was measured by digital microhardness (HVS-50Z070702) under a load of 300gf for 10s. Tensile tests were performed using the Instron 1195 testing machine at a strain rate of 1mm/min.

3Results3.1Microstructure of single β phase by cold rolling3.1.1XRD and TEM characterization

The evolution of microstructure with CR reduction can be observed in TEM images and reflected by XRD measurements (Fig. 2). Fig. 2(a)–(d) of CR specimens show dislocation tangles and cells in different slip systems with increasing CR percentages. This can be associated with the elongated diffraction patterns of CR20% sample shown in the inset of Fig. 2(c). A distinct dislocation structure of subgrain interiors (lower dislocation density) and boundaries (dislocation tangling with higher density) was achieved at large strain (CR90%). XRD tests (Fig. 2(e)) of ST and CR samples at room temperature show that all the samples tested are composed of a single β phase. Moreover, {110} texture increases, which can be associated with the change of relative diffraction intensity I1 1 0/I2 0 0, calculated as 4.04 at ST and 39.4 at CR40%. Microstrain of the CR samples was determined by means of XRD measurements [17]. The XRD pattern shows that the measured Bragg reflection profile is a convolution of the functions representing both the instrumental and the physical broadening profile. The instrumental broadening profile was subtracted as a Gaussian type in the present work. According to XRD measurement, the mean microstrains in those samples of ST (CR0%), CR10%, CR20%, CR30%, and CR40% conditions are given as 0.04, 0.33, 0.42, 0.58 and 0.56%, respectively, as shown in Fig. 2(f).

Figure 2.

TEM of Ti–10Mo–8V–1Fe–3.5Al (wt.%) alloy showing deformed microstructure of matrix in TB3 alloy with the CR reduction of (a) 0%; (b) 10%; (c) 20%; (d) 90%; (e) XRD of ST and CR samples; (f) the variation trend of microstrain within β matrix with CR reduction.

3.1.2Texture development

Fig. 3(a)–(c) shows the texture data in the form of orientation distribution function (ODF) of ST, CR20%, CR40%, CR50% and CR 80% samples. The data were obtained by EBSD measurements scanning on the transverse direction (TD) plane parallel to both the RD and normal direction (ND). Fig. 3(a) indicates that there exists prominent α-fiber component of 0 0 11 1 0 and also strong 1 1 01 1 0 in the ST sample. The orientation density along α-fiber increases with CR reduction up to 20% (Fig. 3(b)) and then decreases at CR40% and CR50% as shown in Fig. 3(c) and (d). When the CR reduction increases to 80%, the main texture component is shifted to 1 1 01 1 0. Besides, the texture component 0 0 10 1 0 occurs in all the CR samples. The IQ maps of different CR reductions in Fig. 3(f) and (h) give insights into the process of dislocation glide, pile up, and finally leads to a three-dimensional network with increasing density.

Figure 3.

ODF at Euler angle φ2=45° for the β textures corresponding to ST and CR samples of (a) ST, (b) CR20%, (c) CR40%, (d) CR50%, and (e) CR80%; image quality map (IQ map) of (f) CR20%, (g) CR50%, and (h) CR80%; and the scale bar for all IQ maps.

3.2Microstructure of α+β phases by short-time aging

With the increase of the dislocation density, the morphology and density of precipitates would also have a corresponding change, which can be seen in the SEM image in Fig. 4. From the SEM micrographs, two significant microstructural variations can be observed. First, compared with direct aging after solution treatment (Fig. 4(a)), there is a substantial increase in precipitate density after pre-deformation by cold rolling (Fig. 4(b)–(e)). Ultrafine α precipitates (∼100nm, Fig. 4(c)) with a uniform dispersion can be observed at CR/aging samples. But the variation tendency of α precipitate density with CR reduction is non-monotonous. The number density of α precipitate reaches its zenith at ∼204.1μm−2 when CR reduction increases from 0 to CR20%. More dislocations in the deformed (>CR20%) samples would not induce more precipitates. This cannot further refine the precipitates that can be associated with a stable change of precipitate size from ∼574 to ∼111nm in the deformed samples shown in the statistics of Fig. 4(i). Second, the PFZs disappear along GBs in CR20% sample with aging process (CRA sample; Fig. 4(g)), compared with the ST sample with the same aging process (STA sample, Fig. 4(f)).

Figure 4.

FESEM of Ti–10Mo–8V–1Fe–3.5Al (wt.%) alloy aged at 500°C after different CR reduction. (a) Coarse precipitates of CR0 (ST)/1h sample with precipitate density of 21.3μm2; (b) CR10%/1h sample with a precipitate density of 122.2μm2; (c) CR20%/1h sample with a precipitate density of 204.1μm2; (d) CR50%/1h with a precipitate density of 204μm2; (e) CR70%/3h with a precipitate density of 200μm2; (f) GBs area of CR0 (ST)/1h sample; (g) GB area of CR20%/1h sample; (h) the relation between α precipitate density and CR reduction; (i) the relation between the size of α precipitate and CR reduction.


TEM (Fig. 5) indicates that the characteristic triangular arrangement of α plates in the sample of ST/500°C/1h changes its frequency of occurrence in the matrix with the increase in CR reduction (as shown in the areas marked in Fig. 5(b)–(e)). It has been determined that there exist 12 possible orientation relationships between α precipitates and β matrix [20]. After cold deformation with subsequent aging process, the number of variants corresponding to 〈111〉β zone axis decreases from three sets usually observed in the ST/aging samples of β titanium alloy [25], to two sets. Moreover, the orientation relationships between α variants and β matrix were determined as (0001)a1//(110)β and (0001)a2//(011)β, as shown in Fig. 5(f). From the XRD results of STA and CRA samples (Fig. 6), the texture intensities of (110)α and (200)β increase with the increment of CR reduction while the peak of (110)β decreases.

Figure 5.

TEM of Ti–10Mo–8V–1Fe–3.5Al (wt.%) alloy aged at 500°C after different CR reduction at (a) CR0 (ST)/1h; (b) CR20%/1h; (c) CR50%/1h; (d) CR70%/1h; (e) CR90%/1h; (f) the selected area electron diffraction (SAED) pattern of [111] β zone axis in CR70%/1h sample and schematic illustration of the SAED pattern (seen in the inset, varied points for different variants).

Fig. 6.

XRD of CR samples with subsequent aging process.


Fig. 7 shows that the orientation of β grain as represented by the (110) pole figure has the six 〈110〉 poles. Pole figures for material in the undeformed condition (Fig. 7(a)) indicate that not all α precipitates and β phases obeyed the Burgers OR when aged at 500°C for 1h within fully recrystallized matrix. The spread of the (0001)α and (110)β poles increases with strain (Fig. 7(b)–(c)), thereby suggesting a deviation from the ideal Burgers OR. However, even after CR20%, there exist four clusters in the (0001)α pole figure (Fig. 7(b)), which still have counterparts in the (110)β (Fig. 7(b)). Only after CR80% (Fig. 7(c)), most of the correspondences disappear indicating a loss of the Burgers OR.

Figure 7.

{110} β pole figure and {0001} α pole figure for CR samples with subsequent aging process: (a) ST/500°C/1h, (b) CR20%/500°C/1h, and (c) CR80%/500°C/1h.

3.3Mechanical behaviors of the samples with ultrafine α+β microstructure

The strengthening by CR can be associated with microhardness (HV) variation of deformed β matrix. As shown in the red curve of Fig. 8(a), the microhardness increases until CR70% and then shows no more increment even under further deformation. However, after a short-time aging treatment as the blue curve indicates in Fig. 8(a), the HV (∼470) of CR70%/500°C/0.5h and CR80%/500°C/0.5h doubles that of the CR70% and CR80% samples prior to aging. This value is higher than the peak hardness of ST/aging samples. When the aging time was prolonged to 1h, the maximum HV value occurs at CR20%/500°C/1h, which is almost equal to that of CR70%/500°C/0.5h and CR80%/500°C/0.5h samples.

Figure 8.

Microhardness (HV) variation. (a) HV changing with CR reduction under fixed aging time; (b) HV variation with time under fixed deformation.


Fig. 8(b) shows the response of microhardness to longer holding time during aging process. When the holding time was prolonged to 3h, the HV of the selected deformed samples (CR20% and CR70%) show a decrease. The decreasing tendency indicates that overaging happens 6h ahead of that of the ST/aging samples seen from the black curve in Fig. 8(b), in which the peak aging hardness occurs after 9h aging.

The mechanical properties of the specimens characterized in terms of ultimate tensile strength and elongation subjected to three CR reduction and aging time conditions are shown in Fig. 9(a). The tensile strength of the aging hardening samples after different CR reduction shows no obvious increment when aging precipitation reaches the maximum volume fraction, as seen in Fig. 9(a). The elongation of the samples changes from 30% (ST/500°C/1h) to ∼2.7% (CR20%/500°C/1h) and 2.5% (CR80%/500°C/1h). However, the elongation increases to 11.7%, with a tensile strength of 1345MPa, when the CR20%/500°C/1h sample was held at 550°C for another 1h. Compared with other available titanium alloys shown in Fig. 9(b), the combination between tensile strength and elongation in such state is quite excellent after the secondary heat treatment at 550°C for 1h. Fig. 9(c) and (d) describes the corresponding microstructure of CR20%/500°C/1h (2#) and CR20%/500°C/1h/550°C/1h (4#) samples, from which α precipitate of 4# sample is coarser (∼250nm) than that of single heat treatment (2# sample, ∼100nm).

Figure 9.

Mechanical properties of CR/aging samples. (a) Tensile strength and elongation (% strain) of CR/500°C/1h samples; (b) Comparison of tensile strength and strain to failure (elongation) of TB3 alloy treated by different processes with other commercially available titanium alloys [29]; (c) the microstructure corresponding to CR20%/500°C/1h (2#) and (d) that of CR20%/500°C/1h/550°C/1h (4#) samples.


Compared with the ω-assisted α nucleation and the pseudo-spinodal decomposition mechanisms, the introduction of defect-assisted nucleation is very effective to boost precipitation in the matrix, when considering that the former two mechanisms are limited by the composition of alloy and stability of the selected materials. It has been identified that precipitation is greatly accelerated by heavy cold deformation mainly through

(1) reducing the misfit strain energy contribution to the activation energy barrier [16];

(2) nucleation on dislocations may be assisted by solute segregation [18];

(3) the dislocation can assist in growth of an embryo beyond the critical size [19].

Besides the effect of precipitates on the mechanical properties, it is confirmed that the effect of crystallographic texture is also far more relevant than just playing a very limited use [20] in titanium alloys. Based on the new experimental results, a detailed description of the texture evolution, precipitation behavior, and mechanical response are proposed in following sections.

4.1Evolution of texture and precipitation behavior during CR and subsequent aging process4.1.1Texture of β matrix

It is to be mentioned here that BCC rolling texture can be largely depicted through φ2=45° sections. Texture evolution of ST and as-rolled samples by the section of φ2=45° is shown in Fig. 3(a)–(e). Generally, the major texture components in the deformation of BCC metals and alloys are the α-fiber {001} <110>, {114} <110>, {112} <110> to {111} <110> and γ-fiber {111} <110> to {111} <112> (visible in φ2=45° section of the ODF) [21]. In the present investigation, the CR sample deformed to CR20% reduction in thickness shows the presence of discontinuous γ-fiber but relatively weak α-fiber shown in Fig. 3(c)–(e). These are somewhat distinct to the usual texture of BCC metals [21]. Furthermore, the γ-fiber components are not strengthened with increasing thickness reduction from CR40% to CR80%. These differences, on the one hand, may be attributed to the starting texture of the ST sample. On the other hand, extensive deformation band formation within the matrix of CR20%–50% samples shown in IQ maps (Fig. 3(g) and (h)) would gradually reorient the grains and weaken the α-fiber and finally lead to the formation of other texture components such as the strong orientation of 1 1 01 1 0 in the CR80% sample. It appears that the formation of deformation heterogeneities plays an important role in weakening the deformation texture, especially after heavy deformation [35].

4.1.2Texture of α precipitates

The α phase formed in the β matrix is well known to be oriented according to the Burgers orientation relationship (OR) [22], viz. {110}β//{0001}α and 〈1 1 1〉β//〈1 1 2¯ 0〉α. Because of the symmetry of the two crystals, theoretically, 12 distinct α orientations can form in a single parent β grain. When all α variants in a β grain occur with equal statistical probability, the transformation is said to evolve without variant selection. However, variant selection takes place during phase transformation because of certain physical reasons. Fig. 7(b) shows that there are four α clusters in the (0001) α pole figure that have the counterparts in the (110)β pole figure at CR20%/500°C/1h. However, the CR80%/500°C/1h sample (Fig. 7(c)) only shows one overlapping pole in a single β grain, which indicates the stronger variant selection along with the increase of CR degree. Besides, the spread of the (0001) α and the (110)β pole figure increases with CR degree (Fig. 7(b) and (c)), thereby suggesting a deviation from the ideal Burgers OR. The strong variant selection is often considered to be deleterious to the mechanical properties of materials [23].

4.1.3Precipitation behavior within pre-deformed β matrix

The processes of nucleation, growth, and coarsening of precipitates are significantly accelerated due to the heterogeneous nucleation of precipitates on dislocations and pipe diffusion through dislocations [8–11]. For the materials subjected to severe plastic deformation, dislocations density ρ can be related to microstrain ɛ approximately by [24]:

where D is the grain size and b is the magnitude of Burgers vector for dislocations. Based on Eq. (1), the dislocation densities corresponding to microstrain can be evaluated as 1.29×1014m−2 (ST), 1.01×1014m−2 (CR10%), 1.29×1014m−2 (CR20%), 1.82×1014m−2 (CR30%), and 1.73×1014m−2 (CR40%). Fig. 2 shows that the mean microstrain (dislocation density) does not monotonically increase but tends to saturate at CR30%. A similar tendency of microstrain and dislocation density was confirmed by other studies on large strain deformation [25,26]. These dislocations induced by CR boost α precipitation. The phenomenon that the number density of α precipitates abundantly increases within the deformed β matrix is in good agreement in the present study with previous results [12,13]. However, the present statistical data (Fig. 4(h) and (i)) indicate that the number and size of α precipitates do not increase monotonically but saturate at about CR20% with the increment of dislocation density within the β matrix. This indicates that the saturation of α precipitates does not rely on the maximum dislocation density during CR. For this viewpoint, Perrard et al. [19] have suggested that the initial dislocation density has a pronounced influence on the maximum precipitate density and proposed a criterion for the end of the nucleation stage on the dislocations. In the criterion, Perrard et al. [19] supposed that nucleation will stop if the time of a solute atom, arriving on a dislocation participating in the nucleation of a new precipitate, is longer than that of diffusing along the dislocation to feed the growth of an existing precipitate. Moreover, their criterion can help explain the aging saturation of the pre-deformed samples in the presented work.

Besides, Fig. 4(g) shows the disappearance of PFZs at CR20% prior to aging compared with that of direct aging treatment (Fig. 4(f)). In previous researches, many trials have been made to decrease the width of the PFZ, where the β phase is stabilized as stabilizing elements since the α stabilizing elements (Al and O) diffuse into α layers along the β GBs [27,29]. The modifications of processes such as aging temperature, quenching procedures, and the addition of alloying elements have been explored by Smith et al. [28]. From our experimental results, pre-deformation plays an effective role on the elimination of PFZs in β Ti-alloys. Meanwhile, from the comparisons between ST/aging in Fig. 5(a) and other CR/aging samples (Fig. 5(b)–(e)), it can be seen that dislocations have a significant selection effect on the 12 α variants during the β→α phase transition, resulting in some individual variants occurring more frequently than others [30,31]. Thus, it can be seen that the delta distributions of α precipitates (Fig. 5(a)) are replaced by the selected variants, e.g. the variant of (0001)α//(011)β in the matrix of the CR70%/500°C/1h sample. Generally, however, the delta distributions are experimentally observed in the ST/aging process most frequently because of the lowest energy compared with other pair-wise combinations of the 12 variants [32].

4.2The effect of precipitate saturation on the mechanical properties in pre-deformed β Ti-alloys4.2.1The influence of direct short-time aging on the precipitate strengthening effect

The microstructure and mechanical properties have been investigated in cold-rolled Ti-1300 alloy [13] and Ti–15Mo–3Al–2.7Nb–0.2Si alloy [33] with differing degrees of reduction. But the present results try to clarify the evolution trend of aging behavior and mechanical behaviors along with the microstructural change. Moreover, aging parameters were also examined here to achieve better strengthening effect with energy-saving process. As shown in Fig. 9(a), the tensile strength and elongation both reach a limit at CR20% and stay stable along with the increment of further deformation prior to aging treatment, which could provide enlightenment in the practice of cold forming with aging process to produce work-pieces of β-Ti alloys.

In discussing the relationship between CR and aging time, the industrially recommended aging time interval is in 6–14h (curve marked by black oval line in Fig. 10) and the range of tensile strength distributes from 1200 to 1400MPa after aging treatment in defect-free alloys. Many trials and studies have been carried out based on these rules. Namely, aging time was often prolonged to achieve higher strengthening effect even in the heavily pre-deformed alloys that can be seen from the statistical distribution of data points in Fig. 10 green area. However, the shorter aging duration at 500°C for 1h of the CR20% sample can strengthen the TB3 alloy to ∼1600MPa which is higher than most statistic data of Ti-alloys, as the marked blue point shown in Fig. 10. Therefore, there may exist possibilities to treat the pre-deformed samples with less time at proper deformation reduction to attain even better strengthening effect. Meanwhile over-aging can be avoided for some β-Ti alloys having fast aging response. It also indicates that selection of aging time corresponding to differing defect populations may need more detailed consideration to tailor the properties of β-Ti alloy.

Figure 10.

Comparison of aging time and tensile strength of pre-deformed TB3 alloy with other commercially available titanium alloys, Titanium Alloys Handbook, Report No. MCIC-HB-02 [34].

4.2.2Optimization of the combination between strength and ductility by the secondary CR/aging process

β-Ti alloys aged after heavily cold deformation always show an inverse proportion of strength to the ductility [34]. To improve the strength–ductility balance, recovery treatment has been conducted to eliminate the high density of dislocations by the subsequent higher temperature recovering. By recovering the samples near the turning point (CR20%/500°C/1h) at 550 °C for 1h (Fig. 9(a)), an improved ductility (∼11.5%) and also a relatively high strength (∼1345MPa) were achieved. Although the prolonged aging duration was believed to induce the coarsening effect of precipitates within the β grain matrix (Fig. 9(c)), the best combination of strength and ductility is achieved after heat treatment at higher temperature. However, further explorations on the effects of recovery on the evolution of α precipitates still need to be done in the future work.


There exists a upper limit of nucleation sites related to dislocation density to obtain the maximum number of α precipitates corresponding to dislocation density of 1.29×1014m−2 (mean microstrain 0.33%, CR20%) in Ti–10Mo–8V–1Fe–3.5Al (wt.%) alloy. The number density of α precipitates increases from ∼21.3μm−2 (ST/aging) to ∼200μm−2 with the CR reduction up to CR20%. However, the dislocation density saturates at about CR30%.

The uniform dispersion of ultrafine α precipitates (∼100–300nm) at CR20% was achieved and characterized by the disappearance of PFZs along GBs. The variant selection of α precipitates is stronger in CR/aging samples than that of the ST/sample. Tensile strength increases from 793 to 1600MPa after aged at 500°C for 1h. The aging time selected was shorter than most industrially applied aging duration.

Recovery treatment at higher temperature (550°C) than direct aging (500°C) near the turning point of hardness (CR20%/500°C/1h) enhanced the ductility of the strengthened samples.

Conflicts of interest

The authors declare no conflicts of interest.


The study was supported by the 973 Program of China (2014CB644003 and 2010CB631003) and the National Natural Science Foundation of China (51271136, 51471136 and 51321003). The authors greatly appreciate the help of Northwest Institute for Non-ferrous Metal Research and Dr. Rajeshwar Reddy Eleti on the analysis of texture.

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