Microstructure homogenization and strength-ductility synergy improvement of the hybrid additive manufactured dual-phase titanium alloy by sub-critical annealing

ABSTRACT

This study employed the effect of sub-critical annealing heat treatment on the microstructure and anisotropy of the hybrid additive manufactured TC11 alloy. The results demonstrate that after sub-critical annealing, a gradient microstructure with the gradual changes in the morphology and content of primary α phases from the laser deposition zone to the substrate zone was formed. Moreover, the tensile anisotropy of the hybrid manufactured specimens was substantially reduced. In particular, when loading along to the deposition direction, the ultimate tensile strength of the heat-treated hybrid manufactured specimen was 1062 MPa, increased by 8.5%. And when loading perpendicular to the deposition direction, the elongation can reach 13.0%, increased by 86% compare to the as-deposited sample. These findings would also shed new insights for the modification of graded structural parts prepared by the hybrid manufacturing technique, laser additive repair and other methods.

KEYWORDS:

• Titanium alloys
• hybrid manufacturing
• heat treatment
• heterogeneous microstructure
• mechanical properties

1. Introduction

Due to the good toughness, superior specific strength and high corrosion resistance, titanium and its alloys have achieved broad application in load-bearing structures with complex geometry, such as aerospace, petrochemical and marine engineering [1–5]. Whereas, the structural complexity and processing volume make it challenging and limited to directly produce large integral structures by a single manufacturing process. Consequently, to overcome these difficulties, a new hybrid manufacturing technique integrating the conventional processing techniques and additive manufacturing technologies has been developed [6–10]. This approach can take full advantage of the high performance and flexibility of directed energy deposition (DED) technology as well as the efficiency [11–16] and low cost of traditional forging or rolling components.

Although the hybrid technique shows considerable promise for producing high performance complex structures, regulating the microstructure still presents certain challenges. On the one hand, in the laser deposition zone (LDZ), the as-deposited titanium alloys usually show the epitaxial growth of coarse β columnar crystal and intragranular lamellar (α+β) basketweave microstructures [17, 18]. The fine lamellar microstructures in the LDZ typically show high strength, poor plasticity and anisotropy, making it becomes the weak zone of the whole hybrid manufactured structure [19]. For example, Chi et al. [20] and Bambach et al. [21] both reported that the hybrid manufactured titanium alloys were fractured at the LDZ side during tensile process. Our previous research has also presented that during the plastic deformation of the bonding zone sample, microcracks preferentially forming at the grain boundaries in the LDZ and gradually expanding to the substrate result in final fracture [22]. On the other hand, the successive heating/cooling effects of the DED process on the substrates during the hybrid manufacturing technique could lead to the formation of the heterogeneous microstructures in the heat affected zone (HAZ), which would affect the performance of the final parts [23, 24]. For instance, Ling et al. [25] reported that the HAZ had the worst plasticity and toughness in the hybrid manufactured TC4 sample. Shen et al. [26] studied the tensile deformation behaviour of the TC17 titanium alloy formed by the hybrid manufactured technique and the results revealed that owing to the poor coordinated deformation among grains, the stress concentrated firstly in the HAZ. Consequently, the above questions of the hybrid manufactured titanium alloys in the as-deposited state limit its applications.

Heat treatments have been proved to be a feasible method to tune the final microstructure and adjust the mechanical properties [27–30]. Many studies have shown that the mechanical properties of the additive manufactured titanium alloys can be significantly improved after the sub-critical annealing at a temperature that slightly lower than the β_trans temperature [31–33]. During this heat treatment, the residual primary α (α_p) phases grow and coarsen, which are conductive to increasing the plasticity, and the ultrafine secondary α phases (α_s) would precipitate from the β phases leading to the improved strength. The combined effect of multiscale α phases can achieve a good match between the strength and plasticity of titanium alloys. For example, Zhao et al. [34] spheroidized the α phases in the laser DEDed TC4 titanium alloy using the sub-critical annealing treatment, resulting in a great increase in the elongation accompanied with a slight decrease in strength within in an acceptable range. Similarly, Zhang et al. [35] proposed that the strength and ductility of the DEDed TA15 titanium after sub-critical annealing heat treatment were simultaneously increased by about 13 MPa and 4.4%, respectively. However, few studies have systematically investigated the effects of heat treatment on the microstructures and mechanical performance of the hybrid manufactured titanium alloys with gradient structures. And it is not clear whether a single heat treatment schedule can meet the optimal strengthening conditions of the multiple components simultaneously. Therefore, there is necessary to design a heat treatment suitable for the hybrid additive manufactured titanium alloys, and study the effects of heat treatment on the microstructures and mechanical properties in detail.

In this study, the effect of heat treatment on the microstructure evolution as well as the tensile mechanical properties of the hybrid manufactured TC11 titanium part were investigated. The microstructural characteristics of the hybrid manufactured samples were analyzed in detail before and after sub-critical annealing treatment using different microscope techniques. The influence of heat treatment on the mechanical performance was evaluated by uniaxial tensile tests to uncover the relationship between microstructures and mechanical properties. Digital image correlation (DIC) and in-situ XRD tensile tests were also conducted to reveal the deformation behaviour of the heat-treated hybrid manufactured TC11 titanium alloy.

2. Experimental methods

2.1. Hybrid manufacturing process and heat treatment

The hybrid manufactured specimens were achieved by depositing the TC11 titanium alloy on the rolled plates. The spherical pre-alloyed TC11 powders for the laser DED process had an average size 177 ± 40 μm, as shown in Figure 1(a). The deposition experiments were carried out using a self-developed DED system, which was equipped with an 8 KW fiber-coupled diode, a beam diameter of 6 mm, a computer numerical control multi-axis motion system, a coaxial powder feeding nozzle and an inert atmosphere protection chamber. During the DED process, the laser power was set at 3500 W, a powder feed rate of 16 g/min, a scanning speed of 1000 mm/min and the oxygen content was less than 50 ppm. The scanning strategy for forming TC11 bulk samples is shown in Figure 1(a). Because the rolled substrates had three different directions including rolling direction (RD), normal direction (ND) and transverse direction (TD), for better analysis, the hybrid manufactured TC11 bulk specimens were deposited at the ND-RD and TD-RD planes, respectively, as depicted in Figure 1(b). Table 1 shows the chemical composition of the laser deposition zone and the substrate zone. Then the specimens were first annealed at 530 °C for 2 h to release the residual stress and it was marked as-deposited alloys. Before formulating the heat treatment schedule, the β_trans temperature of laser DEDed TC11 alloy was determined by the metallographic method and it was 1015 ± 5 °C. Thus, 990 °C was selected as the sub-critical annealing temperature. Specifically, a subset of the hybrid manufactured TC11 samples was sub-critical annealed at 990 °C for 1 h followed by 530 °C for 6 h. And the two cooling methods were air cooling (AC), as illustrated in Figure 1(f). This heat treatment was indicated by the 990HT in this study.

Figure 1. (a) Powder characteristics and schematic illustration of the laser DED process. (b) Different TC11 bulk samples fabricated by hybrid manufacturing technique. (c) The relative positions of the metallographic samples and different kinds of tensile samples. (d) Schematic illustration of in-situ XRD tensile experimental setup. (e) Dimensions of rod-shaped specimens, DIC and in-situ XRD coupons, unit in mm. (f) The heat treatment path for the hybrid manufactured TC11 samples.

Table 1. Chemical composition of the LDZ and SZ (wt.%).

Elements	Ti	Al	Mo	Zr	Si	C	H	O	N
LDZ	Bal.	6.66	3.31	1.80	0.28	0.01	0.002	0.13	0.01
SZ	Bal.	6.60	3.40	1.80	0.30	0.01	0.003	0.12	0.005

2.2. Metallographic characterisation

Metallographic samples of the hybrid manufactured TC11 alloy at different stages were cut through wire electrical discharge machining. These specimens were cut from the cross-section of the hybrid manufactured TC11 plate, as shown in Figure 1(c). They were then ground, polished and etched using Kroll’s agent. The preliminary microstructures were observed using Digital Camera (Nikon D90), Optical Microscope (OM, Leica DM4000) and Scanning Electron Microscope (SEM, Apreo S Lovac). The phase content and size were calculated by Image J software. The texture evolution was analyzed by electron back-scattered diffraction (EBSD) with a field-emission source of 20 kV and a 1.5 μm step size through JSM-7001F (Pegasus XM2). Then the recorded data were processed using TSL OIM software and MATLAB toolbox Mtex-5.7.

2.3. Mechanical properties

According to whether it is along the deposition direction and the sampling location, five different kinds of rod-shaped tensile specimens including the deposition zone (D-L, D-T), the bonding zone (D + S-L, D + S-T) and the rolled substrate (S) were taken out from hybrid manufactured TC11 samples of the as-deposited and heat-treated states. It should be noted that the L and T represent the tensile direction parallel and perpendicular to the deposition direction, respectively. In order to reduce the measurement error, three rod-shaped tensile specimens of five types were performed in accordance with the GB/T228.1-2016 test standard. Besides, the tensile strain in the bonding zone of the heat-treated hybrid manufactured TC11 alloy was analyzed by means of the digital image correlation (DIC) tests. During the DIC experiment, two cameras took pictures every 400 ms to capture the gauge length. In-situ X-ray diffraction (XRD) measurement during the tensile deformation was also conducted to analyze the difference of deformation behaviour between the laser deposition zone and substrate of the D + S-L heat-treated sample. The experiment was carried out on an additional tensile apparatus equipped with a wide-angle XRD using the NANOPIX-WE system (Rigaku Co, Tokyo, Japan). The system was equipped with a rotating Mo anode target as an X-ray source (corresponding to a wavelength of λ, 0.7093 Å). Figure 1(c) also shows the hybrid manufactured specimens and the positions where the rod-shaped samples, DIC and in-situ XRD coupons were extracted. Figure 1(d) illustrates the schematic of in-situ XRD tensile experimental setup. The dimensions of each specimen are illustrated in Figure 1(e). The gauge dimensions of the DIC and in-situ XRD coupons are 30 × 6 × 2 mm³ and 15 × 4 × 0.8 mm³, respectively.

3. Results

3.1. Microstructure characterisation

3.1.1. As-deposited specimen

 Figure 2 shows the microstructures of the as-deposited hybrid manufactured TC11 alloy. As displayed in Figure 2(a), the as-deposited hybrid manufactured TC11 sample can be categorised into three zones: the laser deposition zone (LDZ) in the upper part, the heat affected zone (HAZ) in the middle and the substrate zone (SZ) in the lower part. No defects are found in these three zones of the sample. As illustrated in Figure 2(a–c), the LDZ shows a columnar grain structure that penetrates multiple layers. And the average width of these columnar β grains is about 0.8 mm. Additionally, the microstructure within prior β grains of the LDZ shows a fine basket-weave structure includes continuous grain boundary α (α_GB), parallel-arranged α colony and lamellar α. The total content and thickness of α phases are about 76.3% and 0.7–1.2 μm, respectively. For the HAZ, as Figure 2(d) shows, in the region near the fusion line, recrystallization occurs and coarse equiaxed grains are formed due to the large heat input. Additionally, there are a lot of α colonies in these equiaxed β grains (Figure 2(e)). More detailed microstructural characteristics in other regions of the HAZ and the microstructure formation mechanism can be found in our previous study [22]. The microstructure of the SZ is indicated in Figure 2(f–g), showing an elongated and streamlined grain morphology along the rolling direction, and the volume fraction of α phases is about 82%, which is slightly higher than that in the LDZ.

Figure 2. Microstructures of the as-deposited hybrid manufactured sample: (a) overview of the specimen composed by three different zones [22]; OM and magnified SEM images for the (b)-(c) LDZ, (d)-(e) HAZ and (f)-(g) SZ, respectively.

3.1.2. Microstructures after sub-critical annealing treatment

Figure 3 shows the microstructures of the heat-treated hybrid manufactured TC11 alloy. For the LDZ, after the 990HT heat treatment, the β grain morphology remain unchanged compare with that of the as-deposited sample, as shown in Figure 3(a). It is consistent with the previous research results that when the peak temperature is lower than the β_trans temperature, columnar grains are observed to persist [36]. However, the basket-weave structure within prior β grain transforms into a special bimodal structure that is composed of coarse α_p laths with a width of about 1.6 μm along with ultrafine lamellar α_s in the β_trans matrix, as revealed in Figure 3(b). The α_GB becomes discontinuous due to the high solution temperature near β_trans temperature, which promotes the dissolution of continuous α_GB. In the SZ, the microstructure changed from a typical deformed elongated structure into the conventional bimodal microstructure in which the α_p phases are nearly equiaxed due to the subsequent recrystallization (Figure 3(c)). Figure 3(d) shows the overall microstructure for the HAZ. And according to the distance from the fusion line, the HAZ can be subdivided into three different sub-regions: the top-HAZ, the middle-HAZ as well as the bottom-HAZ. SEM images of Figure 3(e–g) present the detailed microstructural characteristics in these three subzones. On the whole, the α_p laths, equiaxed α_p and ultrafine β_trans exist simultaneously in the three regions of the HAZ, but the volume fraction and the average aspect ratio of the coarse α_p phases are substantially change along the deposition direction. According to Figure 4(a), the volume percentage of the α_p phases gradually decreases from the LDZ (∼ 44.9%) to the SZ (∼ 29.1%). And in the HAZ, the volume fraction of the α_p phases (including the α_p laths and equiaxed α_p phases) follow the sequence of top-HAZ > middle-HAZ > bottom-HAZ. It is possible to pinpoint that, as shown in Figure 4(b), the average aspect ratio of the α_p laths gradually decreases from the LDZ to HAZ, whereas the average diameter of the equiaxed α_p phases maintain constant (∼ 4.6μm) in the HAZ and SZ.

Figure 3. Microstructures of the heat-treated hybrid manufactured sample: (a) overview of the specimen; (b)-(c) OM and magnified SEM images for the LDZ and SZ, respectively; (d) overall microstructural characteristics for the HAZ; (e)-(g) microstructures of the dotted box at the corresponding locations in (d).

Figure 4. (a) The volume percentage of the α_p phases and (b) the average aspect ratio of the α_p laths and the average diameter of the equiaxed α_p phases at different locations in the heat-treated hybrid manufactured TC11 alloy.

3.1.3. Crystallographic texture

Figure 5 shows the inverse pole figure orientation maps and pole figures of the as-deposited hybrid manufactured TC11 sample. The observation of microstructure is on the YOZ and TD-RD cross sections for the LDZ and SZ, respectively. As presented in Figure 5(a and b), the α phases show high texture strength in some certain directions of the LDZ. It can be ascribed to the strong β<001>// building direction texture formed during the laser scanning process and this fiber texture would transform into weak textured α phase strictly followed the Burgers orientation relationship [37]. As shown in Figure 5(c and d), in the SZ, some α grains present a texture state of (0002) basal plane is around 30° away from the perfect <0001 > texture [38]. For comparison with the crystallographic orientation evolution of the hybrid manufactured samples before and after heat treatment. EBSD test is carried out on the heat-treated parts. As depicted in Figure 6(a and b), the β phases in the LDZ of the 990HT sample also display the strong β<001>// building direction texture. It is consisted with the results of previous studies [39, 40]. From the above α_p phases measured results, after the 990HT heat treatment, there are still about 44.9% of the α_p phases in the LDZ, and these uniformly distributed α_p phases are theoretically the same as the texture of the as-deposited state. In addition, the precipitation of the α_s phases also followed the Burgers orientation relationship. Therefore, the crystallographic textures of α orientations in the LDZ (Figure 6(c, d)) are similar to the as-deposited sample. The inverse pole figure orientation maps and pole figures of the α phases in the SZ (Figure 6(e, f)) indicated that the appeared textures of the α phases in rolled microstructure are maintained during the 990HT heat treatment process, but the maximum intensity of texture changes.

Figure 5. EBSD results of the as-deposited hybrid manufactured TC11 sample. (a) Reconstructed β-orientation map from the SZ to the LDZ and the corresponding (b) pole figures in the LDZ. Inverse pole figure orientation maps and pole figures of the α phases in the (c)-(d) LDZ and (e)-(f) SZ.

Figure 6. EBSD results of the heat-treated hybrid manufactured TC11 sample. (a) Reconstructed β-orientation map from the SZ to the LDZ and the corresponding (b) pole figures in the LDZ. Inverse pole figure orientation maps and pole figures of the α phases in the (c)-(d) LDZ and (e)-(f) SZ.

3.2. Mechanical properties

3.2.1. Tensile properties of the L-direction samples

The tensile performances of the L-direction as-deposited and heat-treated hybrid manufactured TC11 alloy are listed in Table 2. The ultimate tensile strength (UTS) of the 990HT D-L, D + S-L and S samples are all increased compared with the as-deposited samples, and the UTS of the D + S-L sample can reach 1062 MPa, which is increased by 8.5%. In terms of plasticity, both the elongation (EL) and reduce in area (RA) of the D-L, D + S-L samples in the heat-treated state are similar to that of the as-deposited state. In addition, the EL and RA of the S sample are slightly decreased with acceptable after the 990HT heat treatment.

Table 2. Room temperature properties of the as-deposited and heat-treated hybrid manufactured TC11 alloy with the L-direction.

State	Samples	UTS (MPa)	YS (MPa)	EL (%)	RA (%)
As-deposited [22]	D-L	979 ± 19	898 ± 16	13.5 ± 1.4	27.0 ± 2.8
	D + S-L	976 ± 4	904 ± 4	10.5 ± 0.7	35.5 ± 2.1
	S	1093 ± 6	1043 ± 6	17.0 ± 1.1	32.5 ± 0.7
990HT	D-L	1053 ± 13	879 ± 17	16.3 ± 1.1	31.0 ± 1.4
	D + S-L	1062 ± 8	921 ± 14	10.6 ± 0.7	30.0 ± 0.7
	S	1178 ± 7	1045 ± 11	13.5 ± 0.8	30.0 ± 0.9

In this study, the heterogeneous deformation behaviours of the heat-treated hybrid manufactured TC11 alloy are mainly studied. Thus, Figure 7(a) also presented the detail engineering strain–stress curves of the 990HT D-L, D + S-L and S samples. It can be confirmed that the UTS (∼1178 MPa) and YS (∼1045 MPa) of the 990HT S sample are higher than those of the 990HT D-L and D + S-L samples. The UTS of the 990HT D + S-L sample is comparable to that of the 990HT D-L sample, while the 990HT D + S-L sample show a higher YS (∼921 MPa). In contrast, compared with the 990HT D + S-L and S samples, the 990HT D-L sample exhibit the highest EL (∼16.3%) and RA (∼31.0%). And the EL is about 10.6% of the 990HT D + S-L sample, which is a little lower than that of the 990HT S sample.

Figure 7. Tensile results of the L-direction heat-treated hybrid manufactured TC11 alloy. (a) Strain-stress curves of the heat-treated D-L, D + S-L and S samples. (b) Macroscopic image near the fracture and the crack propagate path within the β grains of the 990HT D + S-L sample. (c) DIC observations on deformation behaviours of the 990HT D + S-L sample over time.

Subsequently, as shown in Figure 7(b), the fracture position of the 990HT D + S-L specimen is located in the LDZ. The OM micrographs of the fracture profile, indicating that the crack propagation path is tortuous and behaves transgranular ductile fracture. The enlarged SEM image presents that the α_p laths near the fracture has a large-scale deformation. Figure 7(c) displays the strain evolution contour maps from elastic deformation stage to necking characterised by DIC analysis throughout the tensile process of the 990HT D + S-L sample. At the beginning of the tensile test (22 s), the cloud image shows that the whole specimen has uniform elastic deformation, which means that the LDZ, HAZ and SZ bear similar strain. With the progress of the tensile test subsequently, after 40 s, the inhomogeneous local strain concentration in the LDZ is more obvious than that in the HAZ and SZ. Then the strain is further concentrated between 40 and 81 s, and the whole sample undergoes significant deformation. After 81 s, the strain concentration is aggravated, and the obvious necking appears in the LDZ, which eventually leads to the fracture in this zone (137 s). Therefore, when the loading direction parallel to the deposition direction, the strength and plasticity of the LDZ are the kay factors affecting the performance of the 990HT D + S-L sample.

The in-situ XRD tensile results can quantitatively determine the lattice strains of different α phases crystal planes between the LDZ and SZ of the 990HT D + S-L sample during deformation. By calculating the evolution of lattice stress with applied stress, the load partitioning of α phases in different zones can be calculated. One-dimensional diffraction profiles are obtained along the tensile axis, that is., the integral range from −5 to +5° along the longitudinal direction, using the FIT2D software. Note that due to the texture of the LDZ and SZ, it is not achievable to acquire reliable d-spacing information for all α phases. Thus, in this analysis, {10$\overline{1}$0}_α, {10$\overline{1}$1}_α and {10$\overline{1}$2}_α planes with peak intensities are sufficient to fit are chosen for calculation.

The 2D diffraction patterns in the LDZ and SZ of the 990HT D + S-L sample before deformation are shown in Figure 8(a, b), and the diffraction peaks positions of these two zones are similar. More specifically, the diffraction peak of {0002}_α disappears with peak intensity of {20$\overline{2}$0}_α increases in the LDZ compare with the peaks in the SZ. Figure 8(c–f) display the evolution of the X-ray diffraction peaks positions in the LDZ and SZ of the 990HT D + S-L sample under different tensile strains. The diffraction peaks in the LDZ and SZ shift to the lower 2θ angle, indicating that along the tensile direction the d_hkl is increased. Peak broadening is also observed of the two zones, especially in the late stage of deformation. Shifting and broadening of the diffraction peaks indicate the development of lattice strain and the accumulation of dislocations [41]. On the whole, the shift of the {0002}_α, {10$\overline{1}$1}_α and {10$\overline{1}$1}_α diffraction peaks in the SZ lags behind the shift of those peaks in the LDZ after 2% strain, and the variation of the diffraction peaks of {10$\overline{1}$0}_α is larger than that of the {10$\overline{1}$1}_α and{10$\overline{1}$2}_α in the LDZ and SZ (Figure 8(e, f)).

Figure 8. The in-situ XRD tensile test results of the 990HT D + S-L sample. XRD profiles of the un-strained sample in the (a) LDZ and (b) SZ. Evolution of the X-ray diffraction peaks positions in the (c), (e) LDZ and (d), (f) SZ in different tensile strains along the loading direction.

The shifting of the diffraction peaks during the tensile process is generally caused by the lattice strain, and it is can be calculated using the equation below [42], (1) ${ϵ}_{\mathrm{hkl}}^i={\displaystyle \frac{{\mathrm{d}}_{\mathrm{hkl}}^i\left(\sigma \right)-{\mathrm{d}}_{\mathrm{hkl}}^i\left(0\right)}{{\mathrm{d}}_{\mathrm{hkl}}^i\left(0\right)}}$(1) where ${\mathrm{d}}_{\mathrm{hkl}}^i\left(\sigma \right)$ and $d_{\mathrm{hkl}}^i\left(0\right)$ are the d-spacing of the {hkl} plane in phase i under applied stress σ and prior to loading, respectively. The lattice strain along the tensile direction can be calculated by the Eq. (1) and combined with the position of each {hkl} peak recorded in the real time during deformation. Figure 9 shows the evolution of the lattice strains along the tensile direction for {10$\overline{1}$0}_α, {10$\overline{1}$1}_α, {10$\overline{1}$2}_α in the LDZ and SZ. In the initial elastic deformation stage (before about 2% strain), the lattice strains of all peaks in the LDZ and SZ of the 990HT D + S-L sample increase linearly with the macroscopic strain. With the increase macroscopic of strain (2-4% strain), the lattice strain of the LDZ and SZ all increase slowly and after 4% macroscopic strain, the lattice strains tend to be stable in the two zones. According to the previous results [43, 44] that after the end of elastic deformation, lattice strains increase slowly or remain stable of all {hkl} planes owing to the activation of dislocation slipping. Comparing the lattice strain values in Figure 9, the {10$\overline{1}$0}_α, {10$\overline{1}$1}_α, {10$\overline{1}$2}_α phases in the SZ are yielded later than those in the LDZ, indicating that the yield and plastic deformation are occurred in the LDZ of the 990HT D + S-L sample. This phenomenon corresponds with the DIC results.

Figure 9. Evolution of the lattice strains in the (a) LDZ and (b) SZ of the 990HT D + S-L sample during the tensile deformation.

3.2.2. Tensile properties of the T-direction samples

The tensile performance of the different samples (D-T, D + S-T and S) of the as-deposited and heat-treated hybrid manufactured TC11 alloy are listed in Table 3. After sub-critical annealing heat treatment, in comparison with the as-deposited samples, the UTS and YS of the 990HT D-T and D + S-T samples are a little decrease, but they are all increased in the 990HT S samples. Moreover, the EL and RA are greatly improved in the 990HT D-T and D + S-T samples. In addition, the EL and RA of the 990HT D + S-T specimen can reach 13.0% and 29.0%, increased by 86% and 76%, respectively, which is comparable to that of the 990HT S sample.

Table 3. Room temperature properties of the as-deposited and heat-treated hybrid manufactured TC11 alloy with the T-direction.

State	Samples	UTS (MPa)	YS (MPa)	EL (%)	RA (%)
As-deposited [22]	D-T	1168 ± 6	1091 ± 13	6.5 ± 1.4	14.0 ± 0.0
	D + S-T	1205 ± 14	1120 ± 17	7.0 ± 0.7	16.5 ± 0.7
	S	1084 ± 4	1035 ± 7	15.5 ± 1.3	33.5 ± 0.7
990HT	D-T	1104 ± 8	956 ± 2	10.5 ± 0.6	23.0 ± 2.8
	D + S-T	1134 ± 11	990 ± 7	13.0 ± 0.8	29.0 ± 0.6
	S	1182 ± 10	1056 ± 7	13.1 ± 0.2	31.0 ± 0.8

The detail engineering strain–stress curves of the 990HT D-T, D + S-T and S samples are also displayed in Figure 10(a). It should be noted that the UTS (∼1134 MPa) and YS (∼990 MPa) of the 990HT D + S-T sample are between those of the 990HT D-T and S samples. The result is consistent with the strength criterion of the typical metal laminated composites [45, 46]. Furthermore, as shown in the true strain–stress curves of tension (Figure 10(b), ignoring the necking errors), these curves at the elastic deformation stage present three stages. More specifically, stage I changes linearly, while stage II is nearly linear but with a reduced slope and stage Ⅲ varies nonlinearly. Figure 10(c–h) show the microstructures of fracture profile in the different 990HT samples. The fracture mechanism of the 990HT D-T sample is transgranular, i.e. the crack passes through columnar grain boundaries resulting in an increase in plasticity. At the same time, the α_p laths within the β grains near the fracture have undergone severe deformation. As for the 990HT S sample, the crack propagation path is more tortuous and the original equiaxed α_p are elongated and deformed to a certain extent. While for the 990HT D + S-T sample, because the three zones including the LDZ, HAZ and SZ are deformed at the same time, the crack propagation path (Figure 10(e)) is different from the former two types of samples, and it is observed that the deforming degree of the α_p laths near the fracture in the upper LDZ (Figure 10(h)) is more severe than that of the 990HT D-T sample. Figure 11 shows the fractures of the three different tensile samples. The macroscopic images (Figure 11(a, c and e)) show that the shear lip region of the 990HT D-T sample is smaller than that of the 990HT S and D + S-T samples, which corresponds to the low plasticity of the 990HT D-T sample. In addition, the fracture surfaces of the 990HT D-T sample (Figure 11(a, b)) show several cleavage steps accompanied with some dimples and tearing ridges, indicating a quasi-cleavage fracture mechanism. Furthermore, the large and deep dimple characteristics in Figure 11(d and f) support the high plasticity of the 990HT S and D + S-T samples.

Figure 10. Tensile results of the T-direction heat-treated hybrid manufactured TC11 alloy. (a) Strain-stress curves of the 990HT D-T, D + S-T and S samples. (b) The elastic deformation stages of stress–strain curves of different samples. (c)-(e) OM and (f)-(h) SEM micrographs of cracks in the 990HT D-T, D + S-T and S samples, respectively.

Figure 11. SEM images of the fracture for the different 990HT hybrid manufactured samples: (a)-(b) the D-T sample; (c)-(d) the S sample; (e)-(f) the D + S-T sample.

The tensile strain evolution results (Figure 12(a–c)) measured by DIC of the 990HT D + S-T sample are used to further analyze the deformation behaviour in different zones. At the beginning of the tensile process (23 s), similar to the 990HT D + S-L sample, deformation occurs uniformly throughout the whole specimen. As the tensile test progresses (after 58 s), the strain in different zones of the sample began to differ. Specifically, the strain level at the LDZ (left) side is greater than that at the SZ (right) side. Moreover, as indicated in Figure 12(b, c), the strain distribution from top to bottom at the LDZ (left) side is more uneven than that at the SZ (right) side. Different from the 990HT D + S-L sample that the strain is concentrated in the LDZ, the fracture position is random in the 990HT D + S-T sample. It can be seen from the DIC contour map that from 90 s to 157 s, the lower part of the 990HT D + S-T sample shows obvious stress concentration and necking phenomenon.

Figure 12. (a) DIC observations on deformation behaviours of the 990HT D + S-T sample over time. The results showing the strain evolution along a path aligned with the tensile direction over time in the (b) LDZ and (c) SZ.

4. Discussion

4.1. Effect of heat treatment on the mechanical properties of the hybrid manufactured samples

Previous research has indicated that the mechanical properties of titanium alloys are closely related to the microstructures [47]. Generally, the original β grain, the volume fraction, the orientation, the size and shape of α phase are the key factors affecting the mechanical properties of α+β type titanium alloys [48, 49]. Consequently, by adjusting the achieved microstructure, heat treatment can effectively alter the final mechanical properties. In order to intuitively show the changes of tensile properties in different zones perpendicular and parallel to the deposition direction before and after heat treatment of the hybrid manufactured samples, the UTS and EL are summarised in Figure 13. Obviously, after the 990HT heat treatment, the tensile anisotropy in the three zones of the sample has been greatly reduced. It is confirmed that the 990HT heat treatment could improve the comprehensive properties of the LDZ and SZ with different microstructural characteristics.

Figure 13. Longitudinal and transverse tensile properties of different samples in the as-deposited and 990HT states.

For the LDZ of the as-deposited specimen, it exhibits the same high-strength and low-plasticity properties as other titanium alloys made by additive manufacturing [50–52]. Due to the fine and hard α phases and the continuous α_GB phases in the β matrix, they can hinder the movement of the dislocations, making plastic deformation become difficult. Additionally, the high-volume fraction of α phase with a large aspect ratio can lead to an increased stress concentration at the α/β interface during the tensile process, which results in plastic instability. Therefore, for loading along the deposition direction (as-deposited D-L sample), the cracks need to be transferred across the columnar grains boundaries to expand, thereby improving the ductility. However, when loading perpendicular to the deposition direction (as-deposited D-T sample), the continuous α_GB phases are perpendicular to the maximum tensile stress, which has a larger hindering effect on slipping and it is easy to induce rapid crack propagation along the continuous α_GB phases, resulting in low plasticity. After 900HT heat treatment, the orientation and size of the original β grains are almost unchanged, while a special bimodal microstructure is obtained, which is composed of coarse α_p laths and superfine β_trans matrix. Additionally, the α_GB phases gradually appear spheroidization and discontinuity. According to the Hall–Petch effect [53], the coarsening of the α_p would lead to a decline in strength, whereas the increase in the α_p lath volume fraction and the discontinuous α_GB are associated with the increase in plasticity of the 900HT D sample, especially the 900HT D-T sample. As for the SZ, after 900HT, the appeared texture of α phase during hot rolling is maintained while the typical deformed streamlined structure into the conventional bimodal microstructure with equiaxed α_p and fine lamella α. The equiaxed α_p grains play a coordinating role in the deformation and the lamella α phases are beneficial to increase the number of phase interfaces and thus improving the strength [54, 55]. The combined effect of the two morphological structures makes the 900HT S sample have an excellent strength and plasticity matching.

4.2. Deformation and fracture mechanism of the 990HT D + S sample

From the above results, there is a gradient microstructure from the top LDZ to the bottom SZ of the 990HT D + S sample (Figure 3). It is widely accepted that the microstructure determines the mechanical properties, thus the heterogeneous 990HT D + S sample can behave a special deformation behaviour which is preliminarily observed in the DIC test and it is different from the homogeneous structure in the tensile deformation process. Figure 14(a) shows that there is a significant strength mismatch between the LDZ and SZ of the 990HT D + S specimen. This indicates that when the tensile stress is increased to the yield strength of the LDZ, the LDZ begins to yield and undergo plastic deformation. Consequently, the strain observed in the 990HT D + S sample during the tensile test is mainly attributed to this zone because the tensile stress has not reached yield strength of the SZ. Further support for this conclusion can be obtained from the test results of DIC. Figure 14(b) presents the displacement-time curves of the selected points in the tensile gauge length of the 990HT D + S sample. One point is chosen as a datum point on the bottom of the gauge length and the displacement of different points can be measured by l_it = |y_it - y₀|, where l_it is the displacement of point i at moment t. Clearly, the plastic deformation is mainly concentrated in the LDZ, yet the plastic deformation in the SZ can be neglected. Finally, the tortuous transgranular ductile fracture occurs in the LDZ as shown in the Figure 15(a).

Figure 14. (a) The schematic illustration of the mismatch strength in the 990HT D + S-L sample. (b) The displacement-time curves of the selected points during DIC test in the tensile gauge length of the 990HT D + S-L sample.

Figure 15. The crack propagation diagrams of the 990HT (a) D + S-L and (b) D + S-T samples.

It’s reported that the plastic deformation behaviour of α phases is a key factor affecting the yield strength of titanium alloys with a high content of α phases [56]. Moreover, the plastic deformation ability can be characterised by Schmid factor (m), which is the ratio of the critical resolved shear stress (CRSS) to the maximum contact pressure [57]. It serves as an important parameter to evaluate the activity of the slip system. And the larger the Schmid factor, the larger the resolved shear stresses will be and the more easily the slip system can be activated. The main slip systems of α phase with HCP structure are basal slip system ({0001} < 11$\overline{2}$0>), prismatic slip system ({1$\overline{1}$00} < 11$\overline{2}$0>) and pyramidal slip system ({10$\overline{1}$1} < 11$\overline{2}$0 > or {10$\overline{1}$1} < 11$\overline{2}$3>) [58]. Among the three slip systems, the CRSS value of basal slip system and prismatic slip system are almost equal, but all much lower than that of pyramidal slip system. When the loading direction is parallel to the deposition direction, due to the α phase texture variance between the LDZ and SZ, the Schmid factor values for basal slip system and prismatic slip system in the two zones of the 990HT D + S sample are different. Calculated according to the method in the literature [36], the average Schmid factor for basal slip system and pyramidal slip system in the LDZ and SZ are 0.46 and 0.39, respectively. It consists with the in-situ XRD tensile results that the α phases in the LDZ of the 990HT samples are easier to slip than those in the SZ.

As for the 990HT D + S-T sample, when the loading perpendicular to the deposition direction, the LDZ and SZ are all suffer the elastic and plastic deformation. In the elastic deformation stage, similar to the homogeneous material, the strain is uniform and no abnormal phenomenon occurs. However, in the subsequent plastic deformation process, as shown in Figure 10(b), it can be subdivided into three stages. Based on the loading condition and the rule-of-mixture [59, 60], in the stage I, Δσ = [(E_LDZ + E_SZ)/2]·Δϵ. Additionally, as for stage II, Δσ = [((dσ_LDZ/dϵ_LDZ) + E_SZ)/2]·Δϵ can be obtained. Therefore, according to the two equations, the strain–stress curve showing strictly linear transition at stage I, and the decreasing slope of the 990HT D + S-T sample curve at stage II can be attributed to the as decreasing dσ_LDZ/dϵ_LDZ with increasing strain. This mechanical mismatch between the LDZ and SZ will generate a plastic strain gradient. It has been shown that the strain gradient will be accommodated by geometrically necessary dislocations (GNDs). As presented in Figure 16, the GNDs densities in the LDZ and SZ of the 990HT D + S-T sample are all higher than those in the 990HT D-T and S samples, respectively. The pile up of GNDs can produce long-range back stress and create the hetero-deformation induced (HDI) hardening strengthening to enhance the ductility [61–64]. This may be the main reason why the 990HT D + S-T sample has the same high ductility as the 990HT S sample. Moreover, combining the fracture morphologies (Figure 10(g, h)) and DIC results (Figure 12), the micro-cracks preferentially formed in the LDZ and gradually spread to the SZ, resulting in the final fracture of the 990HT D + S-T sample, as depicted in Figure 15(b).

Figure 16. The GNDs density mappings at the same position of different tensile fracture samples: (a) the 990HT D-T sample; (b) the 990HT S sample; (c)-(d) the LDZ and SZ of the 990HT D + S-T sample, respectively.

5. Conclusions

In this work, the hybrid manufactured TC11 alloy was prepared and subjected to the sub-critical annealing treatment. The microstructure evolution and corresponding tensile performance of the hybrid manufactured TC11 alloy before and after heat treatment were evaluated. Furthermore, the deformation and fracture mechanism of the heat-treated hybrid manufactured TC11 alloy were studied. The main conclusions are:

1. The hybrid manufactured TC11 samples including three zones: the LDZ, HAZ and SZ. After sub-critical annealing treatment, the fine basket-weave structures in the LDZ transformed into a special bimodal microstructure contained coarse α_p laths and superfine β_trans matrix. The α_GB also became discontinuous due to the high solution temperature near β_trans temperature. In the SZ, the microstructure changed from a typical deformed streamlined structure into the conventional bimodal microstructure. As for the HAZ, a gradient microstructure that was accompanied with the gradual changes in the morphology and content of α_p phases from top to bottom was formed.

2. The tensile anisotropy of the hybrid manufactured specimens was substantially reduced. When loading along to the deposition direction, the UTS of the hybrid manufactured sample, the LDZ sample and the SZ sample were all increased. When loading perpendicular to the deposition direction, the ductility of the hybrid manufactured sample and the LDZ sample were also improved. In particular, the EL and RA of the heat-treated hybrid manufactured TC11 specimens can reach 13.0% and 29.0%, increased by 86% and 76%, respectively, which was comparable to that of the SZ sample.

3. The occurrence of mechanical mismatch phenomena in the heat-treated hybrid manufactured TC11 samples resulted in non-uniform strain distribution during the tensile deformation process. When loading along to the deposition direction, the average Schmid factor for basal slip system and pyramidal slip system in the LDZ and SZ were 0.46 and 0.39, respectively. This consisted with the in-situ XRD tensile and DIC results that the α phases in the LDZ were easier to slip. When loading perpendicular to the deposition direction, mechanical mismatch between the LDZ and SZ may produce long-range back stress and create the HDI hardening strengthening to enhance the ductility.

Disclosure statement

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

Data availability statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.

Additional information

Funding

This work was supported by the National Natural Science Foundation of China under Grant number 52071005 and the Fundamental Research Funds for the Central Universities under Grant number YWF-23-L-1136.