Effect of double annealing treatments on Ti-10V-2Fe-3Al alloy by directed energy deposition

Abstract

In order to improve the mechanical property of Ti-10V-2Fe-3Al (TB6) metastable β titanium alloy made by directed energy deposition (DED) technique, a group of double annealing treatments were designed. The microstructure and mechanical properties changing after every heat treatment step were systematically investigated. Compared with the traditional solution and aging heat treatment, a microstructure of the fork-like primary α phase (α_p) and plate-like secondary α phase (α_s) were observed after the new-designed double annealing treatment which the first annealing temperature was below β-transus with the ductility of the alloy increased significantly with slight strength reduction. The optimal tensile properties were obtained by the double annealing treatment with the first annealing temperature above β-transus (820 °C/2 h, AC + 520 °C/8 h, AC).

Keywords

Additive manufacturing directed energy deposition Ti-10V-2Fe-3Al microstructure mechanical property metastable β titanium alloy

Introduction

In recent years, especially in aerospace, the applications of metastable β titanium alloys are increasing because they offer the highest strength-to-weight ratios and have very attractive mechanical properties such as strength, toughness, hardenability and fatigue resistance [1,2]. Compare to the majority metastable β titanium alloys such as β–III and Ti-8V-8Mo-2Fe-3Al [3,4], Ti-10V-2Fe-3Al (TB6) alloy has lower densities, economical manufacturing costs and higher age hardening kinetics. It is widely used for aerospace structural components, such as the airframes and landing gear for Boeing 777, Airbus 340 and Airbus A-380, etc. [1,2,5,6].

Directed energy deposition (DED) is an advanced manufacturing technique that can fabricate fully dense near-net shape metal components layer-by-layer directly from CAD models [7 -11]. And, it is a potential way to fabricate high-performance titanium components with comparable large size and complicated shapes due to its lower cost, shorter processing times and greater designing flexibility [11,12]. Thus, extensive research has been carried out to explore directed energy deposition titanium alloys [13 -16]. However, the mechanical properties of as-deposited titanium alloys prepared by DED still need to improve further before they can be used in industry.

It is widely recognised that the mechanical properties of alloys are dependent, to a great extent, on their microstructure [17]. Therefore, many studies were carried out to improve the mechanical properties of titanium by optimising the heat treatment conditions. For example, Zhang et al. [18] analysed the microstructure changes of a wrought TB6 under different heat treatment conditions and discovered that when the solution temperature was 820°C, very fine β-grains could be obtained. G. Terlinde et al. [19] indicated that the recommended solution and aging treatment process for the wrought TB6 alloy was aging at 760 °C for 2 h followed by water quenching (WQ) and re-heating to 520 °C for 8 h with subsequent air cool (AC). However, most of the works were focused on the near α or α + β titanium alloy prepared by laser additive manufacturing such as Ti–4Al–1.5Mn [20], Ti–6Al–4V [21 -24], Ti–5Al–5Mo–5V–1Cr–1Fe [25] and Ti–6Al–2Zr–1Mo–1V [26], etc, only a few studies have been carried out on DEDed metastable β alloy.

As it was known for TB6 titanium alloy, the optimisation of solution and aging treatment conditions have been found to be an effective way to obtain comprehensive mechanical properties [27,28]. After solution treatment and aging process, TB6 alloy always exhibits a duplex microstructure with coarsened primary α phase (α_p) and fine secondary α phase (α_s). The traditional solution treatment and aging process contain two steps: (i) heating to 700–780 °C (T1), holding for 1–3 h, WQ; (ii) holding at 520–550 °C (T2) for 6–8 h,AC. Hence, the recommended solution treatment and aging process is 760 °C/2 h, WQ + 520 °C/8 h, AC. However, previous studies on different kinds of heat treatment routes of TB6 alloys mainly focused on the wrought parts [29,30]. Due to the obvious difference in microstructure between DEDed titanium alloys and wrought titanium alloys [31], and the dynamics of microstructure evolution for DEDed titanium alloys during heat treatment are also quite different from that of severely deformed microstructures [32,33]. Thus, the heat treatment schedules suitable for wrought alloys may not be optimal for DEDed alloys. And recent research has also shown that by annealing at a temperature higher than the β-transus, the morphology and size of the β grains for the near-β alloys will change, which could result in the improved anisotropy of the tensile properties [34]. Therefore, a new heat treatment suitable for DEDed TB6 titanium alloy is needed to be designed.

In this study, four different double annealing heat treatments and three solution schedules were designed to find out the optimal heat treatment method which can attain the well strength-ductility balance of a TB6 titanium alloy. The changes of microstructure and mechanical properties after every heat treatment step were comprehensively explored. Meanwhile, we also discussed the relationships between microstructures and properties of TB6 alloys, and explanations for the excellent strength-toughness of TB6 alloys after the double annealing treatment with the first annealing temperature above β-transus.

Experimental procedures

The materials used in this study were the TB6 titanium alloy pre-alloy powders with a particle size of about 80–200 mesh (made by vacuum plasma rotating electrode atomisation technique). And the base-material was a pure titanium TA0 substrate with a thickness of about 3 mm. The chemical composition of the powders is shown in Table 1. Before directed energy deposition, the pre-alloyed powders were dried in a vacuum chamber at 180 ± 5 °C for 6 h. The directed energy deposition system consisted of a YSL-10000 fibre laser, a BSF-2 powder feeder, a co-axial powder delivery nozzle, and a Fagor-8055 computer numerical control (CNC) four-axis working table. In order to prevent the melt pool from oxidising, the experiments were conducted inside an argon-purged processing chamber with oxygen contents of less than 100 ppm. The DED processing parameters were as follows: laser power was 5000 W, beam diameter was 6 mm, laser scanning speed was 800–1500 mm/min, and powder feed rate was 15–25 g/min. The dimensions of the plate-like sample were approximately 300 mm × 300 mm × 35 mm.

Table 1

The chemical composition of powders and As-deposited substrate.

	Chemical composition (mass fraction, %wt)
	Al	V	Fe	C	N	H	O	Ti
Allowable Specification [35]	2.6∼3.4	9.0∼11.0	1.6∼2.2	≤0.05	≤0.05	≤0.015	≤0.13	balance
TB6 Powders	2.92	10.51	1.63	0.0069	0.0082	0.0018	0.12	balance

Table 2 shows the heat treatment processes. The effect of recommended heat treatment under condition of 760 °C/2 h, WQ + 520 °C/8 h, AC (HT8) on microstructure and tensile properties were investigated to discuss the suitability for DEDed TB6 alloy and to compare with other double annealing treatments. New double annealing treatment under condition of 790 °C/2 h, AC + 520 °C/8 h, AC (HT2) was designed to analyse the effect of AC cooling method on microstructure and tensile properties of TB6 alloy. Moreover, 820 °C/2 h, AC (HT3) and 820 °C/2 h, AC + 520 °C/8 h (HT4) were designed to study the difference of microstructure and tensile properties between below and above β-transus temperature, and the effect of second annealing temperature on microstructure and tensile properties of TB6 alloy was also investigated under 790 °C/2 h, AC + 500 °C/8 h (HT5) and 790 °C/2 h, AC + 510 °C/8 h (HT6). Finally, 790 °C/2 h, AC (HT1), 820 °C/2 h, AC (HT3) and 760 °C/2 h, WQ (HT7) double annealing treatments were selected to explore the effect of the second annealing treatment on microstructure and tensile properties of TB6 alloy.

Table 2

Heat treatment schedules for investigating the microstructure evolution and tensile properties.

No.	Heat treatment process
HT1	790 °C/2 h, AC
HT2	790 °C/2 h, AC + 520 °C/8 h, AC
HT3	820 °C/2 h, AC
HT4	820 °C/2 h, AC + 520 °C/8 h, AC
HT5	790 °C/2 h, AC + 500 °C/8 h, AC
HT6	790 °C/2 h, AC + 510 °C/8 h, AC
HT7	760 °C/2 h, WQ
HT8	760 °C/2 h, WQ + 520 °C/8 h, AC

Samples were prepared by conventional mechanical polishing and etched method. The etching was carried out using a mixture of 1 ml HF, 6 ml HNO₃ and 100 ml H₂O to reveal grains as well as grain boundaries. Then, microstructures of samples were characterised using an optical microscopy (OM, LECIA-DM4000M), a scanning electron microscopy (SEM, JEOL-6010) and a high-resolution transmission electron microscope (TEM, JEM-2100F). Different phases were quantitatively analysed using an Image-P software. The hardness measurements were carried out on a FUTURE-TECH DM800 hardness tester, with a constant load of 1000 g and dwelling time equal to 15 s. And the final Vickers hardness of each sample was obtained as an average of 10 measurements.

The dimensions of the heat-treated samples for the tensile test were 70 mm × 35 mm × 15 mm. After heat treatments, three cylindrical specimens with a gage length of 70 mm and a diameter of 12 mm were cut by electric discharge wire from the centre of heat-treated samples. Then, the cylindrical specimens were processed into the tensile specimen with a diameter of 5 mm, as shown in Figure 1. After heat treatment, oxidised layer of about 0.2 mm was removed completely and which would not affect the experiment results. Room-temperature tensile properties were tested according to the test standard of ISO 6892-1: (GB/T 228.1—2021). Here, the loading direction of specimens was parallel to the deposition direction. The deformation strain was measured by an extensometer. Three tensile test samples were tested for each treatment to determine the mechanical properties. The fracture surfaces and cross sections of the tensile test specimens were examined by SEM and OM, respectively.

Figure 1

Geometry and dimension of the tensile test specimen.

Results and discussion

The effect of heat treatment on microstructure of TB6 alloys

Comparing solution treatment and aging treatment

The microstructure evolution of TB6 alloy after double annealing treatment is shown in Figure 2. After HT1, there were many fork-like primary α phase (α_p) and plate-like secondary α phase (α_s) appeared in β matrix with grain boundary α phase (α_GB) distributed inhomogeneously in the alloy, see Figure 2(a) and Figure 2(b). The distribution of α_GB was closely related to the orientation between two adjacent β grains [36]. And from the SEM photograph shown in Figure 2(b), the fork-like α_p was found to distribute along the preferred direction. After the second annealing treatment (HT2), α_s precipitated more completely and coarsened significantly with the slight coarsening of α_p and α_GB. SEM photograph shown in Figure 2(d) indicates that the α_p shows an obvious triangled precipitation position. Moreover, the forked shape of α_p became more pronounced after HT2, mainly due to the organisation of the alloy tending to be more stable after the second annealing treatment and α_s tending to precipitate along the preferred direction at the tip of α_p.

Figure 2

OM (a, c) and SEM (b, d) photographs showing the microstructure evolution of new solution treatment and aging: (a, b) HT1 (c, d) HT2.

Compared with the microstructure of samples after solution and aging treatment, as shown in Figure 3, the volume fraction and the shape of α_p as well as the shape of α_s all changed significantly. Comparing Figure 3(a,b), when cooling (first step) under WQ condition, the morphology of α_p was lath-like with a volume fraction of about 6.2%. No α_s phase was observed in the β matrix, which was due to the limited time for α_s to precipitate from the β matrix under such a fast cooling rate, as shown in Figure 3(a). Comparing Figure 3(a,c), it can be found that the α_s precipitated more completely after the second heat treatment, which was mainly because the α_s had sufficient time to precipitate out during the following aging process from β matrix. When the first cooling method changed to AC, the volume fraction of α_p increased to 17.3%, much larger than that under WQ condition, and some α_s could be found in the β matrix. However, as the AC process went on, the temperature decreased with the reduction of diffusivity of elements, α_s could hardly precipitate out completely, as shown in Figure 3(b).

Figure 3

SEM photographs showing the microstructures of different heat treatment: (a) HT7 (b) HT1 (c) HT8 (d) HT2.

Comparing Figure 3(c,d), we can find that the shape of α_s was significantly different at different cooling conditions. When the first step cooling method was WQ, α_s were fine and randomly distributed. However, when the first step cooling method was AC, α_s had a plate-like shape. It is because the cooling rate was relatively fast under WQ condition, with the incomplete phase transformation from β to α, most of the α stabilising elements were solid-solution in β matrix, causing the increase of driving force for phase transformation in the aging process. Therefore, the nucleation of α_s will be very fast, resulting in the formation of fine and dispersed α_s. If the first cooling method is AC, the driving force for phase transformation of the aging process will be relatively small, and the α_s has sufficient time to grow bigger along the preferred direction during AC process, forming a plate-like morphology. Hence in the following aging process, the new-formed α_s will precipitate out more completely and coarser.

It's well known that large α plates have smaller total interfacial energy than that of small α plates. According to the common tangent rule, the equilibrium concentration of α stabilising elements around large α plates is lower than that around the small α plates [37]. And the high temperature at the beginning of AC process makes atoms diffuse quickly. Thus, the large α plate (α_p) prefers coarsening and small α plates (α_s) will dissolve to minimise the total free energy in system. It means that α_s prefers heterogeneous nucleate on α_p rather than homogeneous nucleation and grows along the preferred orientation. Meanwhile, the high concentration of α stabilising elements near the tip of α_p will also promote the precipitation of α_s, leading to the form of the unique fork-liked morphology with a high volume fraction of α_p [14,38]. Although large amounts of αs precipitate out inhomogeneously from the β matrix when the first cooling method is WQ, because of the high cooling rate, the precipitation method is still heterogeneous, which will not form a fork-like morphology. Thus, the effect of the secondary heat treatment on the morphology of TB6 alloys is mainly achieved by influencing the precipitation of α_s.

Heat treatment with the first annealing temperature above β-transus

To investigate the difference in microstructure when the first annealing temperature above β-transus, a double annealing treatment with the first annealing temperature above β-transus was designed, and the microstructure evolution is shown in Figure 4. Through the OM photographs, it can be seen that there was no α_p was observed in the whole heat treatment process. This is because when the first annealing temperature was above β-transus, α_p would completely dissolve in β matrix. After the first annealing heat treatment, some part of α_s phase precipitated from β matrix as the OM photograph shown in Figure 4(a). However, compared with the first annealing temperature below β-transus, the width of α_s plates was thinner and more dispersed, which was due to the larger thermal gradient and the homogeneous nucleation of alloys cooled from the β phase region under HT3 or HT4. Moreover, compared with Figure 4(b,d), the α_s plates have precipitated more completely after the second annealing treatment. It can be shown that even when the first annealing temperature was above β-transus, the driving force for phase transformation from this high-temperature gradient was still insufficient to obtain an equilibrium microstructure. Therefore, when the first annealing temperature exceeds the β-transformation temperature, a second heat treatment could be designed to adjust the microstructure of TB6 alloy to obtain excellent mechanical properties.

Figure 4

OM (a, c) and SEM (b, d) photographs showing the microstructure evolution of above β-transus heat treatment: (a, b) HT3 (c, d) HT4.

Heat treatment with different second annealing temperature

In order to investigate the effect of different second annealing temperatures on the microstructure of the TB6 alloy, OM photographs of the microstructure of the alloy at three different second annealing temperatures are shown in Figure 5. Compared with the Figure 5(b,d, and f), it can be found that there was no obvious change in the size of both α_p and α_GB. But α_s became much coarser as the second annealing temperature rises. That is to say; the second annealing treatment mainly affects the morphology of the α_s. This also verifies the conclusion in Section 3.1.1 that the effect of the second heat treatment on the morphology of α_p is indeed achieved by influencing the precipitation of α_s. When the first cooling method is AC, the driving force for phase transformation is small during the subsequent aging process, and α_s tends to precipitate along the preferred growth direction, forming a fork-like morphology without changing the thickness of α_p. However, as the temperature of the second heat treatment rose, the precipitation of α_s was more complete, resulting in the coarsening of α_s, as shown in Figure 5(b,d,f). This is because the increase in aging temperature leads to a decrease of the driving force for the phase transformation of the supersaturated β to α [39], with a decrease in the nucleation rate of α_s, resulting in a dispersed distribution of α_s. Moreover, the significantly increasing in the size of α_s can be attributed to the faster diffusion speed of atomic at higher temperatures.

Figure 5

OM (a, c, e) and SEM (b, d, f) photographs showing the microstructures of different aging temperature heat treatment: (a, b) HT5 (c, d) HT6 (e, f) HT2.

The effect of heat treatment on tensile properties of TB6 alloys

As it was known, the optimised mechanical properties of the alloy, closely dependent on microstructure, could be obtained by adjusting the morphology and volume fraction of different phases under different heat treatment schedules [40]. According to the literature [41,42], the mechanical properties of TB6 titanium alloy can be adjusted by controlling the volume fraction, size and morphology of α phase, the continuity of α_GB, and grain size. A good match between strength and plasticity can be achieved by reducing the continuity of α_GB and balancing the size and volume fraction of α_p and α_s [43 -46].

Table 3 lists the tensile properties of laser additive manufactured TB6 alloy in this study, including the yield strength (YS), ultimate tensile strength (UTS), elongation (EL) and reduction of area (RA). Compared to the sample with WQ as the first cooling method (HT8), samples first cooled by AC showed a significant increase in ductility with slight strength reduction. It is because the sample after HT8 had a large amount of fine and dispersed α_s in the β matrix, as shown in Figure 3(c), making dislocation movement difficult. The microstructure of the secondary annealing treatment with AC as the first cooling step consisted of fork-like α_p and plate-like α_s, as shown in Figure 3(d) and Figure 5(b,d,f), which showed better ductility. It indicates that the double annealing treatment with AC as the first cooling step could improve the plasticity of the samples after the solution and aging heat treatment.

Table 3

Tensile properties of laser additive manufactured TB6 alloy.

Heat treatment process	UTS, MPa	YS, MPa	EL, %	RA, %
760 °C/2 h, WQ + 520 °C/8 h, AC (HT8)	1142 ± 11	1092 ± 12	4.2 ± 0.8	8.3 ± 2.3
790 °C/2 h,AC (HT1)	1019 ± 20	916 ± 15	10.0 ± 1.8	20.0 ± 1.7
790 °C/2 h, AC + 500 °C/8 h, AC (HT5)	1066 ± 17	1021 ± 18	10.0 ± 1.9	29.0 ± 8.1
790 °C/2 h, AC + 510 °C/8 h, AC (HT6)	1056 ± 19	1007 ± 13	9.3 ± 3.3	27.0 ± 9.1
790 °C/2 h, AC + 520 °C/8 h, AC (HT2)	1033 ± 15	979 ± 18	10.7 ± 2.6	35.7 ± 7.9
820 °C/2 h,AC (HT3)	1060 ± 40	970 ± 36	10.7 ± 3.1	39.0 ± 11.8
820 °C/2 h, AC + 520 °C/8 h, AC (HT4)	1107 ± 17	1057 ± 15	9.5 ± 0.8	20 ± 4.4
Allowable specification [35]	≥1105	≥1035	≥8	≥15

Compared with the tensile properties of samples just after the first annealing treatment, both UTS and YS rose significantly and the ductility almost did not change after the double annealing heat treatments. That is to say, when AC is used as a coolling condition at the first step, the second annealing treatment mainly increased the strength of the sample. The improvement of strength is due to the α_s precipitating more completely in the second annealing treatment process, which hindered the dislocation movement. And as the second annealing temperature rose, the strength of the alloys reduced and the ductility also almost did not change. Combining with the microstructure, as shown in Figure 5, the coarsening of α_s lead to the decline of strength. When the first annealing temperature was above β-transus, the strength of the TB6 alloy was higher, with nearly no change of alloy ductility. This is because α would be wholly transformed into β when the first annealing temperature was above β-transus. Therefore, no α_P could be found in the β matrix after the first annealing treatment, as shown in Figure 4, which will lead to a saturation of alloying elements in β. In the subsequent AC process, the driving force for phase transition is large, which will lead to a large amount of finer and more dispersed α_s precipitating in β matrix during the first cooling process, making the dislocation movement more difficult. After the second annealing treatment, α_s precipitates more completely, so the strength of TB6 alloy is further improved. Moreover, after HT4 heat treatment, the specimen has fine plate-like α_s, which are easier to deform than fine and dispersed α_s, resulting in improved ductility. Furthermore, it is noteworthy that both the strength and ductility can meet the requirements of the allowable specification of TB6 alloy when the first annealing temperature is above β-transus temperature.

In order to explain the high ductility of the sample under the condition of 820 °C/2 h, AC + 520 °C/8 h (HT4), the fracture surfaces and cross sections of tensile samples are examined. The fracture surface after recommended heat treatment (HT8) is shown in Figure 6(a). And the sample after HT4 is further given in Figure 6(b). It can be seen that the fracture surface is mainly characterised by a ductile transgranular fracture feature. Comparing with the HT8 sample which failed by intergranular fracture with a few large dimples observed, see Figure 6(a). The proportion of intergranular fracture features of the HT4 sample significantly reduces (see Figure 6(b)) and the area of large dimpled increases (see Figure 7(a)). The fracture subsurface image (see Figure 7(b)) also indicates that the cracks propagate along a transgranular path, it means that the tensile fracture mechanism of HT4 sample changes from intergranular to transgranular. It can be assumed that the fracture mechanism of double annealing heat-treated samples with the first annealing temperature above β-transus is dominated by transgranular fracture, resulting in an excellent match between the strength and plasticity of the HT4 sample.

Figure 6

Fracture surface of room temperature tensile specimens (a)showing the intergranular fracture morphologies after HT8 (b) showing the predominantly transcrystalline fracture morphologies after HT4.

Figure 7

Fracture surface of room temperature tensile specimens after HT4 (a) High magnification SEM fracture photograph (b) OM photograph of fracture subsurface.

The low ductility of DEDed samples after traditional solution treatment and aging is because of the fine and dispersed α_s, which hinder the movement of the dislocations. During deformation, due to the poor ability of grain boundaries to resist deformation, the dislocation will concentrate in the grain boundary with the formation of stress. So the cracks mainly nucleate along grain boundary during loading. Finally, the specimen is characterised by low plasticity and intergranular fracture, as shown in Figure 6(a). Here, for double annealing heat-treated samples, better strength-plasticity synergy is obtained because of the change of microstructure leading to the ability of intracrystalline deformation increasing. When the first step cooling method is AC, the α_s plates are thin and elongated. The reinforcement effect of the thin and elongated α_s plates is not as strong as the fine and dispersed α_s. And micro-hardness testing also shows that the micro-hardness of samples after double annealing treatment is about 371.6 HV which is lower than 402.6 HV of the TB6 specimen after recommended solution treatment and aging. Double annealing treatment is suitable for the laser additive manufactured titanium alloys, but it may not be appropriate for the wrought material. For wrought TB6 alloy, the deformation of the wrought-based processes that break up the grain boundaries. For example, the bi-modal microstructure in wrought material usually exhibits excellent mechanical properties.

Conclusion

In the present study, the TB6 titanium alloy was prepared by DED and a group of double-annealing heat treatments were designed to find the optimal heat treatment schedule. The microstructure and mechanical properties changes after every heat treatment step were systematically investigated. The main findings can be succinctly summarised as follows:

When the first cooling method was AC, the microstructure of DEDed TB6 alloys after double annealing treatment consisted of fork-like α_p and plate-like α_s with α_GB distributed inhomogeneously in the alloy. And when the first cooling method changed to WQ, the morphology of α_p was lath-like with no α_s phase observed in β matrix. After the subsequent aging process, fine and randomly distributed α_s hindered the dislocation movement, resulting in the highest strength of TB6 alloys after HT8.

The optimal tensile properties can be obtained by double annealing treatment with the first annealing temperature above β transformation temperature (HT4). Compared with the recommended heat treatment (HT8), with the strength merely decreased by about 35 MPa (about 3.06%), its elongation and reduction of the area increased by about 5.3% (about 126.25%) and 11.7% (about 140.96%), respectively.

Biography

Hongwen Deng was involved in investigation, methodology, resources and writing – original draft. Zongge Jiao carried out investigation and resources. Xu Cheng was responsible for funding acquisition, resources, supervision and writing – review & editing. Shijie Qi was responsible for funding acquisition and writing – review & editing. responseXiangjun Tian, Yudai Wang and Shuquan Zhang took part in resources, supervision and writing – review & editing.

Footnotes

Acknowledgements

This work was supported by the National Natural Science Foundation of China [52090044], and the Science and Technology Innovation 2025 Major Project of Ningbo [2022Z014].

Disclosure statement

No potential conflict of interest was reported by the author(s).

References

Briggs

RRBARD.

The use of β titanium alloys in the aerospace industry. J Mater Eng Perform 2005;14(6):681–685.

Colombo-Pulgarin

Biffi

Vedani

Beta titanium alloys processed by laser powder Bed fusion: a review. J Mater Eng Perform. 2021;30(9):6365–6388.

G. T. Terlinde

TWD

Williams

JC.

Microstructure, tensile deformation, and fracture in aged Ti-10V-2Fe-3Al. Metall Trans A. 1983;14A:2101–2115.

T. W. Duerig

GTT

Williams

JC.

Phase transformations and tensile properties of Ti-10V-2Fe-3AI. Metall Trans A. 1980;11A:1987–1998.

Zygula

Wojtaszek

Lypchanskyi

The investigation on flow behavior of powder metallurgy Ti-10V-2Fe-3Al alloy using the prasad stability criterion. Metallurg Mater Trans A-Phys Metallur Mater Sci. 2019;50A(11):5314–5323.

Banerjee

Williams

JC.

Perspectives on titanium science and technology. Acta Mater. 2013;61(3):844–879.

Lewandowski

Seifi

Metal additive manufacturing: a review of mechanical properties. Annu Rev Mater Res. 2015: 1–4.

Abbott

Arcella

FG.

Laser forming titanium components. Adv Mater Process. 1998;153(5):29–30.

Arcella F G

FFH.

Producing titanium aerospace components from powder using laser forming. JOM. 2000;52(5):28–30.

10.

Herderick

Additive manufacturing of metals: a review. Mater Sci Technol. 2011;2011:1413–1425.

11.

Wang

HM.

Materials’ fundamental issues of laser additive manufacturing for high-performance large metallic components. Acta Aeronaut Astronaut Sin. 2014;35(10):2690–2698.

12.

Wang Huaming

Xiangming

Progress and challenges of laser direct manufacturing of large titanium structural components (invited paper). Chin J Lasers. 2009;36:3204–3209.

13.

Allison M. Beese

BEC.

Review of mechanical properties of Ti-6Al-4 V made by laser-based additive manufacturing using powder feedstock. J Miner, Met Mater Soc. 2015;68(3):724–734.

14.

Liu

Wang

Tian

Microstructure and tensile properties of laser melting deposited Ti–5Al–5Mo–5V–1Cr–1Fe near β titanium alloy. Mater Sci Eng A. 2013;586:323–329.

15.

Yanyan Zhu

Wang

Microstructure evolution and layer bands of laser melting deposition Ti-6.5Al-3.5Mo-1.5Zr-0.3Si titanium alloy. J Alloys Compd 2014;616:468–474.

16.

Zhu

Tian

Microstructure and mechanical properties of hybrid fabricated Ti–6.5Al–3.5Mo–1.5Zr–0.3Si titanium alloy by laser additive manufacturing. Mater Sci Eng A. 2014;607(0):427–434.

17.

Chen

Study on the relationship between microstructure and mechanical property in a metastable beta titanium alloy. J Alloys Compd. 2015;627:222–230.

18.

Zhang

Hao

Xiong

Effects of solution and aging treatment parameters on the microstructure evolution of Ti-10V-2Fe-3Al alloy. High Temp Mater Processes. 2020;39:501–509.

19.

Terlinde

Fischer

Beta titanium alloys. J Aircr. 1996;26(3):335–346.

20.

Tian

Zhang

Effect of annealing temperature on the notch impact toughness of a laser melting deposited titanium alloy Ti-4Al-1.5Mn. Mater Sci Eng, A. 2010;527(7–8):1821–1827.

21.

Kelly

Kampe

SL.

Microstructural evolution in laser-deposited multilayer Ti-6Al-4 V builds: Part II. Thermal modeling. Metallurg Mater Trans A. 2004;35(6):1869–1879.

22.

Kelly

Kampe

SL.

Microstructural evolution in laser-deposited multilayer Ti-6Al-4 V builds: Part I. Microstructural characterization. Metallurg Mater Trans A. 2004;35(6): 1861–1867.

23.

Zhai

Galarraga

Lados

DA.

Microstructure, static properties, and fatigue crack growth mechanisms in Ti-6Al-4 V fabricated by additive manufacturing: LENS and EBM. Eng Fail Anal. 2016;69:3–14.

24.

Han

Tang

Kou

Experiments and crystal plasticity finite element simulations of nanoindentation on Ti–6Al–4 V alloy. Mater Sci Eng A. 2015;625(0):28–35.

25.

Liu

Wang

Tian

Subtransus triplex heat treatment of laser melting deposited Ti–5Al–5Mo–5V–1Cr–1Fe near β titanium alloy. Mater Sci Eng A. 2014;590:30–36.

26.

Liang

Y-J

Liu

Wang

H-M.

Microstructure and mechanical behavior of commercial purity Ti/Ti–6Al–2Zr–1Mo–1 V structurally graded material fabricated by laser additive manufacturing. Scr Mater. 2014;74:80–83.

27.

Sugiura

Nakamura

Hamai

Mechanical properties of beta solution treated and aged Ti-10V-2Fe-3Al alloy. Tetsu Hagane-J Iron Steel Inst Japan. 2000;86(3):181–188.

28.

Zhang

Qiao

Microstructural evolution in the surface of Ti-10V-2Fe-3Al alloy by solution treatments. Prog. Nat Sci-Mater Int. 2020;30(1):106–109.

29.

Bogucki

Mosor

Nykiel

Effect of heat treatment conditions on the morphology of alpha phase and mechanical properties in Ti-10V-2Fe-3Al titanium alloy. Arch Metall Mater. 2014;59(4):1269–1273.

30.

Al-Salihi

Bettles

Muddle

The ageing behavior of titanium alloy Ti-10V-2Fe-3Al. 7th pacific Rim International Conference on Advanced Materials and Processing, Cairns, AUSTRALIA, Aug 02-06; Cairns, Australia. 2010; pp 843–846.

31.

Zhu

Chen

Tang

Influence of heat treatments on microstructure and mechanical properties of laser additive manufacturing Ti-5Al-2Sn-2Zr-4Mo-4Cr titanium alloy. Trans Nonferr Met Soc China. 2018;28(1):36–46.

32.

Wang

Lei

Effect of heat treatment on microstructures and tensile properties of TA19 alloy fabricated by laser metal deposition. Mater Sci Eng A-Struct Mater Prop Microstruct Process. 2020;782. 139284.

33.

Thijs

Verhaeghe

Craeghs

A study of the micro structural evolution during selective laser melting of Ti-6Al-4 V. Acta Mater. 2010;58(9):3303–3312.

34.

Liu

Lyu

A review on additive manufacturing of titanium alloys for aerospace applications: directed energy deposition and beyond Ti-6Al-4 V. JOM. 2021;73(6):1804–1818.

35.

Editorial, C. A. M. H.

Titanium and copper alloys. Beijing: Standards Press of China; 2001.

36.

Jiao

The spatial distribution of a phase in laser melting deposition additive manufactured Ti-10V-2Fe-3Al alloy. Mater Des. 2018;154:108–116.

37.

Callister

WD.

Materials science and engineering: an introduction (2nd edition). Mater Des. 1991;12(1):59.

38.

Liu

Structure and formation mechanism of α/α interface in laser melting deposited α + β titanium alloy. J Alloys Compd 2016;657:278–285.

39.

Zhao

Shi

Cao

Effects of solution cooling modes and aging temperature on microstructure and tensile properties of TB8 titanium alloy. Hot Working Technol. 2022;51(14):126–130.

40.

Terlinde

Duerig

Williams

JC.

The effect of heat teatment on microstructure and tenstile properties of Ti -10V-2Fe-3Al. Proc 4th Int Conf Titanium. 1980;2:1571–1581.

41.

Wang

Mechanical behavior and microstructural evolution during cyclic tensile loading-unloading deformation in metastable Ti-10V-2Fe-3Al alloy. Mater Sci Eng A-Struct Mater Prop Microstruct Process. 2022;835.142663.

42.

Wang

Microstructure and microhardness evolution of Ti-10V-2Fe-3Al alloy under tensile/torsional deformation modes. J Alloys Compd. 2021;881.160484.

43.

Duerig

Albrecht

Richter

Formation and reversion of stress induced martensite in Ti-10V-2Fe-3Al. Acta Metall. 1982;30(12):2161–2172.

44.

Jackson

Dashwood

Christodoulou

The microstructural evolution of near beta alloy Ti-10V-2Fe-3Al during subtransus forging. Metall Mater Trans A. 2005;36(10):2871–2871.

45.

Chamanfar

Huang

Pasang

Microstructure and mechanical properties of laser welded Ti-10V-2Fe-3Al (Ti1023) titanium alloy. J Mater Res Technol. 2020;9(4):7721–7731.

46.

Bao

Chen

Zhang

Achieving high strength-ductility synergy in a hierarchical structured metastable beta-titanium alloy using through-transus forging. J Mater Res Technol. 2021;11:1622–1636.