J. Cent. South Univ. (2021) 28: 1601-1614

DOI: https://doi.org/10.1007/s11771-021-4720-z

Selective laser melted near-beta titanium alloy Ti-5Al-5Mo-5V-1Cr-1Fe：Microstructure and mechanical properties

HUANG Hua-long(黄华龙)¹, LI Dan(李丹)¹, CHEN Chao(陈超)^{1, 2}, LI Rui-di(李瑞迪)¹,

ZHANG Xiao-yong(张晓泳)¹, LIU Shi-chao(刘世超)¹, ZHOU Ke-chao(周科朝)¹

1. State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China;

2. Shenzhen Research Institute, Central South University, Shenzhen 518057, China

Central South University Press and Springer-Verlag GmbH Germany, part of Springer Nature 2021

Abstract: In this work, a near-beta Ti-5Al-5Mo-5V-1Cr-1Fe titanium alloy was fabricated by selective laser melting (SLM), and the microstructure evolution together with the mechanical properties was studied. The as-fabricated alloy showed columnar β grains spreading over multiple layers and paralleling to the building direction. The distinct microstructure of as-fabricated alloy was composed of near-β (more than 98.1 %) with a submicron cellular structure. Different SLM processing parameters such as hatch spacing could affect the microstructure of as-fabricated alloy, which could thus further significantly affect the mechanical properties of as-fabricated alloy. In addition, the as-fabricated alloy with the distinct microstructure exhibits yield strength of 818 MPa combined with elongation of more than 19 %, which shows that SLM is a potential technology for manufacturing near-beta titanium components.

Key words: selective laser melting; Ti-5Al-5Mo-5V-1Cr-1Fe; near-β and β-titanium alloy; cellular structure; precipitation

Cite this article as: HUANG Hua-long, LI Dan, CHEN Chao, LI Rui-di, ZHANG Xiao-yong, LIU Shi-chao, ZHOU Ke-chao. Selective laser melted near-beta titanium alloy Ti-5Al-5Mo-5V-1Cr-1Fe：Microstructure and mechanical properties [J]. Journal of Central South University, 2021, 28(6): 1601-1614. DOI: https://doi.org/10.1007/s11771-021-4720-z.

1 Introduction

Selective laser melting (SLM) as an advanced additive manufacturing process, can build functional and complex parts by selectively melting consecutive powder layers with a laser beam [1-3]. Thus, SLM can be used to manufacture complex aerospace components that cannot easily be produced by other methods [4]. A large number of studies have focused on the SLM of titanium alloys due to its high level of specific strength and excellent biocompatibility. Until now, most studies of SLM Ti alloys have been carried out in near-α and (α+β)-titanium alloys, such as Ti-6Al-4V [5-15] and CP-Ti [16, 17] alloys. More and more studies have been performed to fabricate near-β and β-titanium alloy specimens by additive manufacturing [18-23], some of which have focused on near-β and β-titanium alloys by SLM. Some such studies have included selective laser melted Ti-5553 [24], Ti-10V-2Fe-3Al [25], Ti-24Nb-4Zr-8Sn [26], Ti6Al4V-ELI [27], Ti-37Nb-6Sn [28], Ti-25Nb-3Zr-3Mo-2Sn [29], Ti-13Nb-13Zr [30] and other alloys. Near-β and β-titanium alloys (VT22, Beta C, Ti-5553, Ti-10-2-3, etc.) are the most versatile titanium alloys and are also attractive materials for aerospace applications such as landing gear due to their high specific strength, good hardenability, excellent fatigue/crack propagation and corrosion resistance [24, 31, 32].

The other major advantage of β-titanium alloys is the tunable nature of the microstructure by varying the thermomechanical processing [33-35]. The ductility of near-β titanium alloy has been reported to be largely dependent on the prior β grain size; other mechanical properties, such as strength, were expected to be dependent on the distribution of α precipitates in the β matrix [33, 36]. As mentioned previously, in addition to a large number of studies focused on Ti-6Al-4V alloys, more and more works had been done to understand the evolution of the microstructure in near-β and β-titanium alloys obtained by additive manufacturing methods. LIU et al [19] studied laser melting deposited (LMD) Ti-5Al-5Mo-5V-1Cr-1Fe near-β titanium alloy, and found that its microstructures exhibited fine basketweave microstructure and continuous grain boundary α (αGB). LIU et al [26] studied a β-type Ti-24Nb-4Zr-8Sn manufactured by electron beam melting (EBM) and SLM, and found that the microstructure of EBM and SLM samples consisted of α+β phases and a single β phase, respectively. Particularly, the faster cooling rate during SLM promoted the formation of fine β dendrites. SCHWAB et al [24] studied the near-β titanium alloy Ti-5553 processed by SLM, and found that the microstructure of samples consisted of a pure β phase and the mechanical properties of samples exhibited a tensile strength of about 800 MPa and a strain up to 14%. Compared with LMD and EBM, SLM is more suitable for preparing parts with the complex and fine structures because of its higher precision characteristics. Therefore, it is of great significance to study the preparation of a near-β Ti-5Al-5Mo-5V-1Cr-1Fe titanium alloy by SLM. Of particular interest is the relationships between the SLM process, the microstructure and the mechanical properties of a near-β Ti-5Al-5Mo-5V-1Cr-1Fe titanium alloy obtained by SLM.

In this research, a typical near-β Ti-5Al-5Mo-5V-1Cr-1Fe titanium alloy with a nominal density of 4.62 g/cm³ was selected for SLM [37]. The aim of this research was to understand the relationship between the microstructure and mechanical properties of Ti-5Al-5Mo-5V-1Cr-1Fe titanium alloy. On this basis, the typical microstructure of the as-fabricated specimens was characterized, and its evolution during SLM was further investigated. In addition, the effects of the microstructure on the tensile properties were studied to evaluate the feasibility of the mechanical properties for aerospace applications [38].

2 Experimental

2.1 Materials

The powders were produced by electrode induction gas atomization. The chemical composition of the gas atomized powders was given in Table 1. The chemical composition was in agreement with the standardized Ti-5Al-5Mo-5V-1Cr-1Fe alloy [37]. The Ti-5Al-5Mo-5V-1Cr-1Fe alloy powders with a bulk density of 2.05 g/cm³ exhibited a spherical morphology (Figure 1(a)). The powder particles were spherical with a diameter of 2.8-86.4 μm (the median size d50 was 31.5 μm), as shown in Figure 1(b).

Table 1 Chemical composition of as-received gas atomized Ti-5Al-5Mo-5V-1Cr-1Fe powder (mass fraction, %)

2.2 SLM method and preparation of specimens

The Ti-5Al-5Mo-5V-1Cr-1Fe alloys were produced on a Farsoon 271 M SLM system. As shown in Figure 2(a), the SLM system parameters were as follows: the laser power was 0-500 W, the layer thickness was 0.02-0.1 mm (the default value was 0.03 mm), the default laser hatch spacing was 0.12 mm, the default value of laser beam diameter was 90 μm, and the default value of building plate temperature was 100 °C. To avoid oxidation of the melt pool, the experiments were conducted inside an argon purged processing chamber with an oxygen content of less than 1×10^-3.

As shown in Figure 2(b), the specimens were placed perpendicular to the laser beam (Z axis). The specimens were fabricated by a focused and high-energy laser beam. A standard alternating X/Y raster scanning strategy was selected. There were bidirectional hatches of a layer “N” along the Y-axis, while the next layer “N+1” was rotated to 67°. After SLM, cuboid specimens (60 mm×10 mm×10 mm) were prepared from the titanium alloy substrate by wire electrical discharge machining. In addition, we calculated the energy density E (energy applied per unit volume of the specimen), which was used to represent the energy input by the following equation [39]:

　　　　　　　　　　　　　　　　(1)

where P is the laser power; v is the laser scanning velocity; h is the hatch spacing, and t is the layer thickness.

Figure 1 SEM image (a) and particle size distribution (b) of Ti-5Al-5Mo-5V-1Cr-1Fe alloy powders

Figure 2 Schematic of SLM:

2.3 Characterization

The density of the specimens was measured by the OHAUS AX224ZH (Archimedes principle). Each specimen was measured at least three times after the necessary drying treatment.

As shown in Figure 2(b), all the tensile testing specimens were machined by wire electrical discharge machining from as-fabricated cuboid specimens according to ASTM E8M standard. The mechanical properties were characterized by averaging three measured values for the tensile specimens at room temperature in air using an Instron 3369 mechanical testing machine. Uniaxial tensile tests were carried out at a fixed strain rate of 0.1 min^-1. The yield strength was measured using the 0.2 % offset method. The fracture surfaces and cross sections of the tensile testing specimens were observed by SEM.

Optical microscope (OM, DM2700P, Germany Leica) and scanning electron microscope (SEM) (Quanta 250 FEG, FEI, Czech) with a secondary electron signal for imaging were used to analyse the microstructures of the fabricated samples. Specimens for OM and SEM testing were prepared using a mixture of 1.5 mL HF, 3 mL HNO₃ and 100 mL H₂O after mechanical lapping and polishing. Electron back-scattered diffraction (EBSD) measurements were performed using a dual beam microscope system (FEI, Helios Nanolab G3 UC). The analytical methods combining the EDAX OIM Analysis^TM software and the EBSD results were used as in our previous work [40]. All specimens were investigated by X-ray diffraction (XRD) analysis (RigakuD/max 2500 V, Japan). To quantitatively analyse the effects of thermal cycles, the size of the melt pools was measured using Image-Pro Plus 6.0 software.

3 Results and discussion

3.1 Density

Table 2 shows the effect of the SLM parameters on the densities of the as-fabricated specimens at laser scanning velocities of 1000 mm/s, laser powers from 200 W to 300 W, hatch spacings of 0.05 mm and 0.12 mm. Since the energy density, E, with 0.05 mm hatch spacing was higher than that of the default hatch spacing, the average density of specimens with 0.05 mm hatch spacing was larger. However, all the relative densities were higher than 99.6% and similar when applying the standard alternating X/Y raster scanning strategy, which ranged from 4.60 to 4.62 g/cm³ with an average value of 4.61 g/cm³. These values are similar to what is known from literature for the relative densities of SLM-prepared specimens [24]. Unmelted powder particles were found at the fracture macromorphology of the as-fabricated specimen in Figure 13, when the energy density was 55.6 J/mm³. The selected SLM system parameters with high energy densities (≥55.6 J/mm³) could fabricate specimens with very high density.

3.2 Evolution of microstructure

3.2.1 Columnar microstructure

The typical microstructures of the specimens prepared with the default hatch spacing, powder layer thickness of 0.03 mm, laser scanning velocity of 1000 mm/s and laser power of 250 W were observed and analysed. As shown in Figure 3(a), typical columnar grains spreading over multiple layers and parallel to the built direction from the bottom to the top were observed at the macroscale. The columnar grains were up to 1 mm or more in the X-Z plane (along the specimen build direction) (Figure 3(b)). Fusion boundaries were visible, such as the proximate semi-elliptical lines shown in Figure 3(b) and the parallel tracks shown in Figure 3(c), indicating the melted tracks and cross sections of the laser beam scanning paths. The proximate semi-elliptical braided structures of the X-Z plane (along the specimen build direction) were approximately 143 μm in width and 27 μm in height. Rotated melted tracks were observed in Figure 3(c) and were similar to the hatch spacing (0.12 mm) and the rotation angle (67°) that were used in this research. As shown in Figure 4, all specimens were investigated by XRD analysis, but only the β phase was detected. To further confirm the phase composition, several sets of specimens were investigated by EBSD analysis, typical columnar β grains spreading over multiple layers and parallel to the built direction were clearly observed as a result of epitaxial growth, as shown in Figure 5(a); however, the approximate equiaxial microstructure can clearly be observed in Figure 5(b). But a small amount of α (less than 1.9%) was unevenly distributed inside the typical columnar β grains in Figures 5(c) and (d).

Table 2 Densities of as-fabricated specimens at different hatch spacings and power conditions

Figure 3 OM images of as-fabricated specimens:

The grain morphology in the X-Y plane (perpendicular to the specimen build direction) of the specimens prepared by 0.05 mm hatch spacing, 0.03 mm powder layer thickness, 1000 mm/s laser scanning velocity and 275 W laser power were observed and analysed in Figure 6(a). Columnar β grains were spherical with diameters mainly from 3.3 to 166.4 μm (the average dimension was 105 μm) in Figure 6(a). As shown in Figure 6(b), the grain morphology in the X-Y plane (perpendicular to the specimen build direction) of the specimens prepared by 0.12 mm hatch spacing, 0.03 mm powder layer thickness, 1000 mm/s laser scanning velocity and 275 W laser power were observed and analysed. Columnar β grains were spherical with diameters mainly from 2.5 to 109.2 μm (the average dimension was 50.7 μm) in Figure 6 (b). Since the energy density, E, with 0.05 mm hatch spacing was higher than that of the default hatch spacing, the specimens with 0.05 mm hatch spacing had coarser columnar grains in the X-Y plane (perpendicular to the specimen build direction).

Figure 4 XRD diffraction pattern of as-fabricated specimens with 250 W laser power

To better understand the evolution of the columnar grains, a schematic illustration is shown in Figure 7. During the SLM process, when the high energy density laser beam was focused on the current powder layer, the powders in the laser affected region could be completely melted to liquid and form a melt pool. The already solidified matrix was similar to a heat sink, and heat was rapidly transferred from the liquid phase to the already solidified matrix resulting in a large temperature gradient from the liquid phase to the substrate. At the same time, when a layer absorbed heat from the laser beam, the heat was quickly transferred to the next layer of the solidified substrate thereby promoting repeated heating and melting. Due to the extremely high cooling rate of SLM [13], the solidification of the β liquid phase had no nucleation barrier. Therefore, the large β grains rapidly and epitaxially grew from the bottom to the top and along the centre of the melt pools, which was parallel to the local heat transfer direction. Meanwhile, both the cooling rate and temperature gradient stabilized at the top of the melt pool during solidification and promoted the formation of large columnar β grains. In addition, the grain orientation was closely related to the scanning strategy. The type of control of texture via modification of the scanning strategy has previously been reported by VRANCKEN et al [27]. In this research, there was a bidirectional track in the X-Y plane (perpendicular to the specimen build direction) with a standard alternating X/Y raster (rotation angle 67°). The macrostructure of the as-fabricated Ti-5Al-5Mo-5V-1Cr-1Fe alloy specimens exhibited many irregular three-dimensional columnar β grains spreading over multiple layers and parallel to the built direction from the bottom to the top, which was different from the bamboo-like macrostructure of LMD [19].

Figure 5 EBSD inverse-pole figure (IPF) maps and phase maps of as-fabricated specimen:(The red region is β phase and the green region is α phase of (c) and (d))

Figure 6 EBSD maps showing grain morphology of as-fabricated specimens with hatch spacing of 0.05 mm (a) and 0.12 mm (b) in X-Y plane (perpendicular to the specimen build direction)

Figure 7 Schematic illustration of columnar grain evolution during SLM

3.2.2 Cellular microstructure

As shown in Figure 8, a lath-like microstructure was observed in the X-Z plane (along the specimen build direction), but a honeycomb microstructure was observed in the X-Y plane (perpendicular to the specimen build direction). Due to the extremely high cooling rate of SLM (10⁴–10⁶ K/s), the multilayer alloy rapidly solidified in the melt pool to form a nonequilibrium microstructure, which was accompanied by insufficient precipitation of the strengthening phase, a homogeneous precipitated phase and a regular morphology phase [41, 42]. Thus, the generation of the cellular structure was the consequence of a combination of structural supercooling theory and compositional fluctuations caused by extremely high temperature gradients and cooling rates. Because of the narrow solidification range of titanium alloy [27], cellular growth occurred mainly in the as-fabricated Ti-5Al-5Mo-5V-1Cr-1Fe alloy specimens, as shown in Figure 9(c). To further understand the cellular structures, the Gaussian laser energy distribution caused a corresponding heat input (Q_s) distribution in the melt pools, as shown in the following equation [42]:

                  (2)

where P_v is the absorbed laser power; f is the heat distribution factor that affects the power distribution; d is the laser beam radius; r is the radial distance from the centre of the laser beam; z is the depth in the current thickness direction, and h is the depth of the energy source.

According to Eq. (2), the heat input (Q_s) decreased exponentially with increasing r in the same horizontal plane in Figure 9(b). So the temperature at the centre of the laser track was much higher than the temperature at the edge region. Then,thermal flux is generated from the centre to the edge due to thermal dissipation [43]. The high dissipation and cooling rate resulted in a liquid metal temperature (T_L) in the central region that was lower than the melting point (T_M) and a degree of subcooling (△T=T_M-T_L) sufficient to form new grain nucleation. Therefore, the liquid metal would undergo simultaneous nucleation and random orientation in the same horizontal plane. In addition, the growth rate of the crystal nucleus was higher at the melt pool interface in the same vertical plane. Equiaxed crystals were easily formed in the X-Y plane (perpendicular to the specimen build direction). As shown in Figure 9(a), the equiaxed crystal had a hexagonal cellular structure in the X-Y plane (perpendicular to the specimen build direction) (approximately 500 nm), which could be explained theoretically by the low energy boundary theory. According to the minimum Gibbs free energy theory, the grain boundaries tended to have the same angle, and the three crystals preferentially formed an intersecting interface. The interface energy (γ) of the ternary grain boundary should satisfy the following equation:

γ₁=γ₂=γ₃                                                   (3)

According to the low energy boundary theory, the interface energy (γ) and the boundary angle (α) should satisfy the following equation:

                       (4)

According to Eqs. (3) and (4), the values of the boundary angles should be equal, i.e., α₁=α₂=α₃. On the other hand, some intersections of four crystals were observed in Figure 9(a), which could not be stably maintained and would be broken down into the intersection of three crystals. Therefore, the coexistence of intersections of three crystals and intersections of four crystals was evidence that the as-fabricated specimens would form a non-equilibrium microstructure due to the extremely high cooling rate of SLM. For elongated cellular structures, the input thermal flux and latent heat of crystallization were released from the centre and reduced the degree of subcooling thereby inhibiting the nucleation of new grains. Therefore, the elongated cellular structures were ubiquitous due to the high crystal growth rate along the heat flux direction.

Figure 8 SEM observation of as-fabricated specimens:

Figure 9 Microstructural evolution analyses of cellular structures:

As shown in Figure 9(c), the X-Z plane (along the specimen build direction) of the melt pool had an elongated cellular microstructure, and transition from planar to cellular solidification at the bottom of the melt pool (0-3 μm) was observed. Planar growth is rare for alloys. The planar front was unstable in the as-fabricated specimens and hard to observe. The mechanism of the transition is shown in Figure 9(d). The evolution of the structure was mainly determined by the ratio of the temperature gradient to the solidification rate (G/R) [41]. Generally, no component subcooled to cause planar growth, narrow components subcooled to cause cellular growth, and broad components subcooled to cause dendrite growth. When the alloy composition was consistent and as G/R decreased, the crystal morphology changed from planar crystals to cellular crystals, cellular dendrites, columnar branches and equiaxed dendrites. At the bottom of the melt pool, the G value was the greatest due to the reduction in the heat input (Q_s). R was close to 0; thus, the G/R ratio was large and caused planar growth. Similar results were also obtained in other alloys, such as Ti6Al4V+10 Mo, obtained by SLM [9]. R gradually increased upward from the bottom of the melt pool, and the G/R value gradually decreased. Therefore, cellular structures grew from the boundary towards the centre of the melt pool, and this direction was also the direction of the heat flow. In addition, dendrite crystals or others formed at the top of the melt pool primarily because these regions were affected by the subsequent laser.

As shown in Figure 10(a), the grain orientations were observed in the X-Y plane (perpendicular to the specimen build direction) and were mainly parallel to the built direction. This direction is an easy growth direction for a body-centred cubic crystal because the growth is the fastest when the <100> crystal direction coincides with the maximum temperature gradient [27]. According to the heat transfer distribution of the melt pool [43], the maximum temperature gradient was perpendicular to the melt pool boundary, and the melt pool boundary was concave. The grains grew from the boundary to the centre of the melt pool, and the increasing angle between the grain orientation and built direction was higher near the top of the melt pool, as shown in Figure 10(b). However, the top of the deposited layer underwent remelting during the laser beam scanning of the secondary layer, which remelted all grains near the top of the melt pool on the first layer. What remained was the bottom part of the melt pool on the first layer where the particles were mainly parallel to the built direction. As shown in Figure 10(c), residual horizontally grown grains in the X-Z plane (along the specimen build direction) after partial remelting of the next layer were observed, and these grains are marked by a red elliptical line and correspond to the position in Figure 10(b). As shown in Figure 10(d), the grain orientations of Grain A and Grain B were different on both sides of the melted track; thus, the partial remelting of the next layer was inferred to affect the grain orientations in the X-Y plane (perpendicular to the specimen build direction). In particular, the smearing effect provided some possibilities for adjusting the scanning strategy to control the grain orientations by SLM.

Figure 10 Grain orientations of as-fabricated specimens:

3.2.3 Phase transitions

As shown in Figure 5, the near-β phase (more than 98.1 %) was detected in the as-fabricated Ti-5Al-5Mo-5V-1Cr-1Fe alloy specimens. According to the research of VRANCKEN et al [27], the amount of beta stabilizer required to maintain a pure β phase at ambient temperature depends on the molybdenum equivalent; approximately 10 wt% of molybdenum stabilized completely and inhibited the α′ martensite transition, and the β phase was completely retained by SLM. The molybdenum equivalents are given by the following:

[Mo]_eq=1.0w(Mo)+0.67w(V)+0.44w(W)+

0.28w(Nb)+0.22w(Ta)+1.6w(Cr)+

1.25w(Ni)+1.7w(Co)+2.9w(Fe)-1.0w(Al) (5)

The effect of Mo on the α′ martensitic transformation was in the following two aspects: reducing the critical cooling rate for retaining β and reducing the initial temperature of α′ martensite by the addition of β stabilizer. Since the Mo_eq of the Ti-5Al-5Mo-5V-1Cr-1Fe alloy was 7.85, abundant β phase can be maintained at room temperature during rapid cooling of SLM.

3.3 Tensile properties

As shown in Figure 11(a), the typical as-fabricated specimens with 0.12 mm hatch spacing, 250 W laser power and 1000 mm/s laser scanning velocity had a good combination of strength and ductility with an ultimate tensile strength (UTS) of (938±4) MPa, a yield strength (σ_0.2) of (801±16) MPa, an elongation to failure of (18.5±1.0)% and a reduction of area to failure of (28.3±3.3)%. The as-fabricated specimens exhibited good extensibility by SLM in contrast to other processing technologies [18, 19, 31]. Due to the low ultimate tensile strength, the as-fabricated specimens could not meet the permissible specifications [11]. In general, residual stresses are considerably large in the as-deposited additive manufacturing parts [1], which are not the final state of titanium alloys in use. An additional post processing heat treatment which can reduce residual stresses and regulate the β→α phase transformation [19], should deliver significantly better strength properties. The anisotropy in strength and ductility was attributed to the {100} fiber texture and columnar grain morphology of the SLM-prepared specimens [40]. Longitudinal SLM-prepared specimens should show lower tensile strength but higher elongation than transverse SLM-prepared specimens, owing to the columnar grains and accelerated damage in transverse specimens under Mode I opening tension when applying tensions [40]. This will be topic of future investigations.

Figure 11 (a) Typical tensile stress-strain curves of as-fabricated specimens; Tensile strength (b) and elongation (c) of as-fabricated specimens with 1000 mm/s laser scanning velocity at different hatch spacings and power conditions

In order to understand the effect of different hatch spacings on tensile properties, several sets of specimens with the same laser scanning velocity (1000 mm/s) were selected for tensile testing. As shown in Figures 11(b) and (c), the specimens with default scan spacing had higher average tensile strength and elongation at a power of 275 W. QIU et al [25] studied a selectively laser melted beta titanium alloy, and found that a low laser power and a short exposure time (i.e., low energy density) led to development of fine columnar β grains and widespread cell structures whereas increased laser power and exposure time (i.e., high energy density) resulted in pronounced grain growth, increased texture and significantly decreased cell structures [25]. Similarly, the higher energy density of specimens with small hatch spacing (0.05 mm) resulted in grain growth and led to development of coarse columnar β grains in Figures 5(a) and (c), which would significantly reduce the tensile strength and elongation.

For the specimens with default scan spacing, the best comprehensive mechanical properties were obtained with 250 W laser power and 1000 mm/s laser scanning velocity. It is generally accepted that the mechanical properties are mainly determined by the as-deposited microstructure [19]. To investigate the cause of the low ultimate tensile strength, the fracture morphology of the tensile specimens was studied, as shown in Figure 12. All specimens had ductile fractures, and the fractures exhibited typical tensile equiaxed dimples. As shown in Figure 12(a), two fracture modes were observed, namely, intergranular fracture and dimple fracture. As shown in Figure 12(b), a very pronounced dimple morphology was observed in the high magnification image. The secondary cracks were surrounded by shallow dimples that were irregular in shape and uneven in size distribution. The blind spots without cellular structure preferentially produced intergranular fractures and acted as sources of secondary cracks. Both the dimples and secondary cracks prolonged the distance and time of crack propagation, indicating that the as-deposited specimens had good ductility. Deformed dimples and secondary cracks indicated mixed fracture mechanisms, including predominant microvoid coalescence and the intergranular cracking mechanism. The low strength of the as-fabricated specimens was related to the microstructure with a nearly pure β phase. As a result, a small amount of segregated α did not play a role in dispersion strengthening. The as-fabricated specimens were composed of submicron cellular structures, which could maintain good tensile strength. However, the segregated α was mainly precipitated inside the columnar β matrix where almost no grain boundary α was precipitated. Therefore, a relatively single β phase with cellular structure was beneficial to increase ductility.

Figure 12 SEM images showing fracture morphology in the X-Z plane (along specimen build direction) of as-fabricated specimens (a) at a low multiple and (b) at a high multiple

In particular, unmelted powder particles were found only at the fracture macromorphology of the as-fabricated specimen with the lowest density (4.60 g/cm³) in Figure 13, which was most likely due to the lowest energy density (55.56 J/mm³).

Figure 13 SEM image showing fracture morphology of as-fabricated specimen with 0.12 mm hatch spacing, 1000 mm/s laser scanning velocity and 200 W laser power

4 Conclusions

A Ti-5Al-5Mo-5V-1Cr-1Fe near-β titanium alloy was fabricated by SLM. The microstructure and mechanical properties of the as-fabricated Ti-5Al-5Mo-5V-1Cr-1Fe were investigated. The following conclusions can be drawn.

1) The selected SLM system parameters with high energy densities (≥55.6 J/mm³) could fabricate Ti-5Al-5Mo-5V-1Cr-1Fe alloy with very high density. As the hatch spacing increased from 0.05 mm to 0.12 mm, the as-fabricated Ti-5Al-5Mo-5V-1Cr-1Fe alloy had finer columnar β grains in the X-Y plane (perpendicular to the specimen build direction). The as-fabricated Ti-5Al-5Mo-5V-1Cr-1Fe alloy consisted of vast columnar β (more than 98.1 %) grains, which were up to 1 mm or more and parallel to the specimen build direction. Additionally, hexagonal cellular structures (approximately 500 nm) were observed in the as-fabricated Ti-5Al-5Mo-5V-1Cr-1Fe alloy.

2) The as-fabricated Ti-5Al-5Mo-5V-1Cr-1Fe alloy with the distinct microstructure exhibits yield strength of 818 MPa combined with elongation of more than 19%. Compared to the allowable specifications, the as-fabricated Ti-5Al-5Mo-5V-1Cr-1Fe alloy by SLM had higher elongation and reduction of area but lower ultimate tensile strength.

Contributors

CHEN Chao provided the concept and edited the draft of manuscript. HUANG Hua-long conducted the literature review and wrote the first draft of the manuscript. LI Dan and LIU Shi-chao analyzed the measured data. LI Rui-di, ZHANG Xiao-yong, and ZHOU Ke-chao edited the draft of manuscript. All authors replied to reviewers’ comments and revised the final version.

Conflict of interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

References

[1] GU D D, MEINERS W, WISSENBACH K, POPRAWE R. Laser additive manufacturing of metallic components: Materials, processes and mechanisms [J]. International Materials Reviews, 2012, 57(3): 133-164. DOI: 10.1179/ 1743280411Y.0000000014.

[2] DUTTA B, FROES F H. Additive manufacturing of titanium alloys [M]. Oxford: Butterworth-Heinemann Elsevier, 2016. DOI: 10.1016/B978-0-12-804782-8.00005-7.

[3] LI Neng, HUANG Shuai, ZHANG Guo-dong, QIN Ren-yao, LIU Wei, XIONG Hua-ping, SHI Gong-qi, BLACKBURN J. Progress in additive manufacturing on new materials: A review [J]. Journal of Materials Science & Technology, 2019, 35(2): 242-269. DOI: 10.1016/j.jmst.2018.09.002.

[4] STOLT R, ELGH F. Introducing design for selective laser melting in aerospace industry [J]. Journal of Computational Design and Engineering, 2020, 7(4): 489-497. DOI: 10.1093/jcde/qwaa042.

[5] LIU Min, LIU Shi-chao, CHEN Wei, CHEN Chao, LV Ya-ping, ZHANG Xiao-yong, LEI Peng-fei, LIN Yong-cheng, ZHOU Ke-chao. Effect of trace lanthanum hexaboride on the phase, grain structure, and texture of electron beam melted Ti-6Al-4V [J]. Additive Manufacturing, 2019, 30: 100873. DOI: 10.1016/j.addma.2019.100873.

[6] ZHANG Qi, LIANG Zheng-long, CAO Miao, LIU Zi-fan, ZHANG An-feng, LU Bing-heng. Microstructure and mechanical properties of Ti6Al4V alloy prepared by selective laser melting combined with precision forging [J]. Transactions of Nonferrous Metals Society of China, 2017, 27(5): 1036-1042. DOI: 10.1016/S1003-6326(17)60121-3.

[7] GUO Kuai-kuai, LIU Chang-sheng, CHEN Sui-yuan, DONG Huan-huan, WANG Si-yu. High pressure EIGA preparation and 3D printing capability of Ti-6Al-4V powder [J]. Transactions of Nonferrous Metals Society of China, 2020, 30(1): 147-159. DOI: 10.1016/S1003-6326(19)65187-3.

[8] THIJS L, VERHAEGHE F, CRAEGHS T, HUMBEECK J V, KRUTH J P. A study of the microstructural evolution during selective laser melting of Ti-6Al-4V [J]. Acta Materialia, 2010, 58(9): 3303-3312. DOI: 10.1016/j.actamat.2010.02.004.

[9] ZHAO Xiao-li, LI Shu-jun, ZHANG Man, LIU Yan-dong, SERCOMBE T B, WANG Shao-gang, HAO Yu-lin, YANG Rui, MURR L E. Comparison of the microstructures and mechanical properties of Ti-6Al-4V fabricated by selective laser melting and electron beam melting [J]. Materials & Design, 2016, 95: 21-31. DOI: 10.1016/j.matdes.2015.12.135.

[10] NEIKTER M, WORACEK R, MAIMAITIYILI T, SCHEFFZUK C, STROBL M, ANTTI M L, AKERFELDT P, PEDERSON R, BJERKEN C. Alpha texture variations in additive manufactured Ti-6Al-4V investigated with neutron diffraction [J]. Additive Manufacturing, 2018, 23: 225-234. DOI: 10.1016/j.addma.2018.08.018.

[11] YU Han-chen, YANG Jing-jing, YIN Jie, WANG Ze-min, ZENG Xiao-yan. Comparison on mechanical anisotropies of selective laser melted Ti-6Al-4V alloy and 304 stainless steel [J]. Materials Science and Engineering A, 2017, 695: 92-100. DOI: 10.1016/j.msea.2017.04.031.

[12] RAFI H K, KARTHIK N V, GONG Hai-jun, STARR T L, STUCKER B E. Microstructures and mechanical properties of Ti6Al4V parts fabricated by selective laser melting and electron beam melting [J]. Journal of Materials Engineering and Performance, 2013, 22(12): 3872-3883. DOI: 10.1007/s11665-013-0658-0.

[13] YANG Jing-jing, YU Han-chen, YIN Jie, GAO Ming, WANG Ze-min, ZENG Xiao-yan. Formation and control of martensite in Ti-6Al-4V alloy produced by selective laser melting [J]. Materials & Design, 2016, 108: 308-318. DOI: 10.1016/ j.matdes.2016.06.117.

[14] MISHRA A K, KUMAR A. Numerical and experimental analysis of the effect of volumetric energy absorption in powder layer on thermal-fluidic transport in selective laser melting of Ti6Al4V [J]. Optics & Laser Technology, 2019, 111: 227-239. DOI: 10.1016/j.optlastec.2018.09.054.

[15] CHEN Cai-yan, GU Dong-dong, DAI Dong-hua, DU Lei, WANG Rui, MA Cheng-long, XIA Mu-jian. Laser additive manufacturing of layered TiB₂/Ti6Al4V multi-material parts: Understanding thermal behavior evolution [J]. Optics & Laser Technology, 2019, 119: 105666. DOI: 10.1016/ j.optlastec.2019.105666.

[16] ATTAR H, LOBER L, FUNK A, CALIN M, ZHANG L C, PRASHANTH K G, SCUDINO S, ZHANG Y S, ECKERT J. Mechanical behavior of porous commercially pure Ti and Ti-TiB composite materials manufactured by selective laser melting [J]. Materials Science and Engineering A, 2015, 625: 350-356. DOI: 10.1016/j.msea.2014.12.036.

[17] ATTAR H, BONISCH M, CALIN M, ZHANG L C, ZHURAVLEVA K, FUNK A, SCUDINO S, YANG C, ECKERT J. Comparative study of microstructures and mechanical properties of in situ Ti-TiB composites produced by selective laser melting, powder metallurgy, and casting technologies [J]. Journal of Materials Research, 2014, 29(17): 1941-1950. DOI: 10.1557/jmr.2014.122.

[18] LI Zhen, TIAN Xiang-jun, TANG Hai-bo, WANG Hua-ming. Low cycle fatigue behavior of laser melting deposited TC18 titanium alloy [J]. Transactions of Nonferrous Metals Society of China, 2013, 23(9): 2591-2597. DOI: 10.1016/S1003-6326(13)62772-7.

[19] LIU C M, WANG H M, TIAN X J, TANG H B, LIU D. Microstructure and tensile properties of laser melting deposited Ti-5Al-5Mo-5V-1Cr-1Fe near β titanium alloy [J]. Materials Science and Engineering A, 2013, 586: 323-329. DOI: 10.1016/j.msea.2013.08.032.

[20] LIU Y J, WANG H L, LI S J, WANG S G, WANG W J, HOU W T, HAO Y L, YANG R, ZHANG L C. Compressive and fatigue behavior of beta-type titanium porous structures fabricated by electron beam melting [J]. Acta Materialia, 2017, 126: 58-66. DOI: 10.1016/j.actamat.2016.12.052.

[21] QIU Chun-lei, RAVI G A, ATTALLAH M M. Microstructural control during direct laser deposition of a β-titanium alloy [J]. Materials & Design, 2015, 81: 21-30. DOI: 10.1016/j.matdes. 2015.05.031.

[22] XIE Le-chun, GUO Hao-jie, SONG Yan-li, LIU Chang, WANG Zhong-qi, HUA Lin, WANG Li-qiang, ZHANG Lai-chang. Effects of electroshock treatment on microstructure evolution and texture distribution of near-β titanium alloy manufactured by directed energy deposition [J]. Materials Characterization, 2020, 161: 110137. DOI: 10.1016/ j.matchar.2020.110137.

[23] MANTRI S A, BANERJEE R. Microstructure and micro-texture evolution of additively manufactured β-Ti alloys [J]. Additive Manufacturing, 2018, 23: 86-98. DOI: 10.1016/ j.addma.2018.07.013.

[24] SCHWAB H, PALM F, KUHN U, ECKERT J. Microstructure and mechanical properties of the near-beta titanium alloy Ti-5553 processed by selective laser melting [J]. Materials & Design, 2016, 105: 75-80. DOI: 10.1016/j.matdes.2016. 04.103.

[25] QIU Chun-lei, LIU Qi. Multi-scale microstructural development and mechanical properties of a selectively laser melted beta titanium alloy [J]. Additive Manufacturing, 2019, 30: 100893. DOI: 10.1016/j.addma.2019.100893.

[26] LIU Y J, LI S J, WANG H L, HOU W T, HAO Y L, YANG R, SERCOMBE T B, ZHANG L C. Microstructure, defects and mechanical behavior of beta-type titanium porous structures manufactured by electron beam melting and selective laser melting [J]. Acta Materialia, 2016, 113: 56-67. DOI: 10.1016/ j.actamat.2016.04.029.

[27] VRANCKEN B, THIJS L, KRUTH J P, VAN HUMBEECK J. Microstructure and mechanical properties of a novel β titanium metallic composite by selective laser melting [J]. Acta Materialia, 2014, 68: 150-158. DOI: 10.1016/j.actamat. 2014.01.018.

[28] CHEN Wei, CHEN Chao, ZI Xu-hui, CHENG Xiao-fan, ZHANG Xiao-yong, LIN Yong-cheng, ZHOU Ke-chao. Controlling the microstructure and mechanical properties of a metastable β titanium alloy by selective laser melting [J]. Materials Science and Engineering A, 2018, 726: 240-250. DOI: 10.1016/j.msea.2018.04.087.

[29] LIU Y J, ZHANG Y S, ZHANG L C. Transformation-induced plasticity and high strength in beta titanium alloy manufactured by selective laser melting [J]. Materialia, 2019, 6: 100299. DOI: 10.1016/j.mtla.2019.100299.

[30] ZHOU Li-bo, YUAN Tie-chui, TANG Jian-zhong, HE Jian-jun, LI Rui-di. Mechanical and corrosion behavior of titanium alloys additively manufactured by selective laser melting-A comparison between nearly β titanium, α titanium and α+β titanium [J]. Optics & Laser Technology, 2019, 119: 105625. DOI: 10.1016/j.optlastec.2019.105625.

[31] IVASISHIN O M, MARKOVSKY P E, MATVIYCHUK Y V, SEMIATIN S L, WARD C H, FOX S. A comparative study of the mechanical properties of high-strength β-titanium alloys [J]. Journal of Alloys and Compounds, 2008, 457(1, 2): 296-309. DOI: 10.1016/j.jallcom.2007.03.070.

[32] SINGH P, PUNGOTRA H, KALSI N S. On the characteristics of titanium alloys for the aircraft applications [J]. Materials Today: Proceedings, 2017, 4(8): 8971-8982. DOI: 10.1016/ j.matpr.2017.07.249.

[33] BOYER R R. An overview on the use of titanium in the aerospace industry [J]. Materials Science and Engineering A, 1996, 213(1, 2): 103-114. DOI: 10.1016/0921-5093(96) 10233-1.

[34] AHMED T, RACK H J. Phase transformations during cooling in α+β titanium alloys [J]. Materials Science and Engineering A, 1998, 243(1, 2): 206-211. DOI: 10.1016/S0921-5093(97)00802-2.

[35] BANERJEE D, WILLIAMS J C. Perspectives on titanium science and technology [J]. Acta Materialia, 2013, 61(3): 844-879. DOI: 10.1016/j.actamat.2012.10.043.

[36] WEISS I, SEMIATIN S L. Thermomechanical processing of beta titanium alloys—An overview [J]. Materials Science and Engineering A, 1998, 243(1, 2): 46-65. DOI: 10.1016/S0921-5093(97)00783-1.

[37] MOISEYEV V N. Titanium alloys: Russian aircraft and aerospace application [M]. Boca Raton: CRC Press, 2005: 75-77. DOI: 10.1201/9781420037678.

[38] XU W, BRANDT M, SUN S, ELAMBASSERIL J, LIU Q, LATHAM K, XIA K, QIAN M. Additive manufacturing of strong and ductile Ti-6Al-4V by selective laser melting via in situ martensite decomposition [J]. Acta Materialia, 2015, 85: 74-84. DOI: 10.1016/j.actamat.2014.11.028.

[39] NI Mang, CHEN Chao, WANG Xiao-jun, WANG Peng-wei, LI Rui-di, ZHANG Xiao-yong, ZHOU Ke-chao. Anisotropic tensile behavior of in situ precipitation strengthened Inconel 718 fabricated by additive manufacturing [J]. Materials Science and Engineering A, 2017, 701: 344-351. DOI: 10.1016/j.msea.2017.06.098.

[40] CARROLL B E, PALMER T A, BEESE A M. Anisotropic tensile behavior of Ti-6Al-4V components fabricated with directed energy deposition additive manufacturing [J]. Acta Materialia, 2015, 87: 309-320. DOI: 10.1016/j.actamat. 2014.12.054.

[41] ZHANG Ming-kang, YANG Yong-qiang, WANG Di, XIAO Ze-feng, SONG Chang-hui, WENG Chang-wei. Effect of heat treatment on the microstructure and mechanical properties of Ti₆Al₄V gradient structures manufactured by selective laser melting [J]. Materials Science and Engineering A, 2018, 736: 288-297. DOI: 10.1016/j.msea.2018.08.084.

[42] HOZOORBAKHSH A, ISMAIL M I S, AZIZ N B A. A computational analysis of heat transfer and fluid flow in high-speed scanning of laser micro-welding [J]. International Communications in Heat and Mass Transfer, 2015, 68: 178-187. DOI: 10.1016/j.icheatmasstransfer.2015.08.013.

[43] LE T N, LO Y L. Effects of sulfur concentration and Marangoni convection on melt-pool formation in transition mode of selective laser melting process [J]. Materials & Design, 2019, 179: 107866. DOI: 10.1016/j.matdes.2019. 107866.

[44] LUO Lei, MAO Xiao-nan, YANG Guan-jun, NIU Rong-rong. Brief Introduction for BT22 titanium alloy [J]. Material & Heat Treatment, 2009, 38(14): 14-16. DOI: 10.14158/j.cnki. 1001-3814.2009.14.038. (in Chinese)

(Edited by HE Yun-bin)

中文导读

选择性激光熔融法制备Ti-5Al-5Mo-5V-1Cr-1Fe近β钛合金的微观组织和力学性能

摘要：本文采用选择性激光熔融法制备了Ti-5Al-5Mo-5V-1Cr-1Fe近β钛合金，并研究了Ti-5Al-5Mo-5V-1Cr-1Fe合金在打印过程的组织演变和力学性能。Ti-5Al-5Mo-5V-1Cr-1Fe合金组织为近β钛合金(相含量大于98.1%)且在β晶粒内部存在大量亚微米级的晶胞组织，β柱状晶平行于构建方向。扫描间距、扫描功率等选择性激光熔融工艺参数显著影响Ti-5Al-5Mo-5V-1Cr-1Fe合金的微观组织和力学性能。制备的Ti-5Al-5Mo-5V-1Cr-1Fe合金屈服强度为818 MPa，伸长率超过19%，表明选择性激光熔融法是一种有潜力的近β钛零件制造技术。

关键词：选择性激光熔融；Ti-5Al-5Mo-5V-1Cr-1Fe；近β和β钛合金；晶胞组织；沉淀

Foundation item: Project(2019B010943001) supported by Key-area Research and Development Program of Guangdong Province, China; Project(2020) supported by the Fund of State Key Laboratory of Powder Metallurgy, Central South University, China

Received date: 2020-12-14; Accepted date: 2021-05-08

Corresponding author: CHEN Chao, PhD, Associated Professor; E-mail: pkhqchenchao@126.com; ORCID: https://orcid.org/0000-0001-6897-2998