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In-situ investigation of the anisotropic mechanical properties of laser direct metal deposition Ti6Al4V alloy
Author’s Accepted Manuscript In-situ investigation of the anisotropic mechanical properties of laser direct metal deposition Ti6Al4V alloy Junxia Lu, Ling Chang, Jin Wang, Lijun Sang, Shikai Wu, Yuefei Zhang www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(17)31567-8 https://doi.org/10.1016/j.msea.2017.11.106 MSA35821
To appear in: Materials Science & Engineering A Received date: 3 August 2017 Revised date: 25 November 2017 Accepted date: 25 November 2017 Cite this article as: Junxia Lu, Ling Chang, Jin Wang, Lijun Sang, Shikai Wu and Yuefei Zhang, In-situ investigation of the anisotropic mechanical properties of laser direct metal deposition Ti6Al4V alloy, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2017.11.106 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
In-situ investigation of the anisotropic mechanical properties of laser direct metal deposition Ti6Al4V alloy Junxia Lua*, Ling Changa, Jin Wangb, Lijun Sangb, Shikai Wua, Yuefei Zhangb** a Institute of Laser Engineering, Beijing University of Technology, Chao Yang District, Ping Le Yuan 100#, Beijing, 100124, China b Institute of Microstructure and Property of Advanced Materials, Chao Yang District, Ping Le Yuan 100#, Beijing, 100124, China E-mail: [email protected] E-mail: [email protected] *Corresponding author: Junxia Lu, Tel.: +86 010-67396557 Fax: +86-10-67392635 **Corresponding author: Yuefei Zhang, Tel: +86-10-67392635 Fax: +86-10-67392635 Abstract This study compares the microstructure and tensile properties of Ti6Al4V components fabricated by laser direct metal deposition (LDMD) additive manufacturing (AM) in the transverse and longitudinal directions. The results show anisotropic tensile properties with the transverse direction having high tensile and fracture strengths and the longitudinal direction having a high elongation and reduction of cross section. The anisotropic mechanical properties are attributed to the anisotropic microstructural distribution. The transverse tensile specimen is composed of short columnar prior-β grains which grow perpendicular to the tensile direction, and have a lamellar structure. Along the β grain boundary, αGB and large α colonies were identified. However, the longitudinal specimen shows that the long β structure is parallel to the tensile axis and that the microstructure is composed of basket-woven α phases with shorter α plates and smaller colony sizes compared with those in the transverse specimen. The fracture mechanism induced by the anisotropic microstructure along the transverse and longitudinal directions was compared by examining the fracture process in real-time using uniaxial in-situ scanning electron microscopy (SEM) tensile testing. The results show that shear fracture, which is caused by the vertical β grain boundaries and large α colonies with long α plates, occurs in the transverse specimen. The shear mode is the main reason behind the enhanced tensile strength and fracture strength due to the high resistance to microcrack propagation. However, in the longitudinal specimens, symmetric necking behavior due to the fine α grains resulted in uniform deformation of the grains on both sides of the grain boundaries, inducing greater elongation. Key words: Additive manufacturing; Titanium alloy; Microstructure; Mechanical properties; In-situ tensile
1. Introduction The Ti6Al4V (wt.%) alloy represents about 50% of the titanium alloy market, with very large quantities being used by the aerospace industry because of its exceptionally high specific strength, good cryogenic properties, and excellent corrosion resistance [1-4]. Parts made of the alloy are
traditionally machined by die forging, hand forging or plate stock. However, these methods cause substantial wastage of material, especially when the geometries involved are complex, resulting in high buy-to-fly ratios [1]. For these types of components, laser near-net-shape additive manufacturing (AM) offers several potential benefits, including an increased raw material utilization ratio and a reduced number of machining operations with a consequential decrease in the cost and lead time [5]. Laser direct metal deposition (LDMD) is a near-net-shape AM technology in which successive layers of metal powder, wires or ribbon are melted onto a workpiece. In this process, a focused laser heat source is used to melt the metallic feedstock and produce a desired shape. Materials can be delivered to the workpiece either by first spreading the individual power layers and then, melting the individual layers selectively in a powder-bed fusion process, or by continuous direction through coaxial nozzles to the melt pool in a directed-energy deposition process. LDMD offers the ability to fabricate fully dense three-dimensional parts with highly complex geometric features and lightweight structures using low-density alloys, thereby resulting in a shorter manufacturing time, less raw material waste, and lower cost [6, 7]. Conventional manufacturing techniques, such as rolling and forging are associated with material deformation and rely on process control and material properties to produce metallic components. The mechanical properties of the product are determined by the process parameters. In comparison, components produced using LDMD AM differ substantially from those produced using conventional manufacturing methods. The final mechanical properties are affected not only by the process parameters but also the solidification behavior of the material. The LDMD AM process involves rapid melting and solidification coupled with high supercooling at the solid/liquid interface. Repetitive thermal cycling by the reheating of the previously deposited layer induces the phase transformation of Ti6Al4V alloy from the body centered cubic (bcc) β-phase to face centered cubic (fcc) α-phase. The steep temperature gradients and repeated heating lead to macroscopically epitaxial grain growth [8-10]. A general observation in earlier studies was the presence of large columnar prior-β grain structures parallel to the build direction when selective laser melting (SLM) deposition is used [11-14]; these structures were also observed when the electron-beam (EB) AM [15-17] and wire + arc AM [18] techniques were used. The width of these prior-β grains depend on the cooling rate during solidification. For instance, the width of the columnar prior-β grain in Ti6Al4V alloy when EB melting was used is reported to be about 25-30 μm [19] while it is 220-2450 μm for laser beam deposition [20]. The coarse columnar β grains that grow from the substrate and extend across multiple layers often tend to favor intergranular failure [11] and the vertical columnar growth along the build direction causes anisotropic mechanical properties in the deposition direction and build direction [12，21]. Within the prior-β grains, a significant amount of acicular α' martensite or lamellar α structure can be observed; the quantities depend on the cooling rate [22]. Some authors also observed grain boundary α phases which distributed along the prior-β grain boundary [3, 18]. However, Bermingham et al [23] report that the grain boundary α was eliminated and the prior-β size decreased by the trace additions of boron to the Ti6Al4V feedstock material. Such nonhomogeneous microstructural distribution within a single component often leads to difficulties in characterizing and predicting the final structural performance. This is particularly relevant when designing parts for aircraft assembly application using AM technique [24]. The anisotropic tensile properties of AM Ti6Al4V were tested and the results showed greater
elongation and reduction of cross section in the perpendicular direction while higher values of yield strength and ultimate tensile strength were observed in the deposition direction [8, 11]. Some reports partially attribute the high elongation in the perpendicular orientation to the less porous components [25, 26]. In addition, the anisotropy is less pronounced for the powder-based processes compared with that for wire-based processes because of the lower heat input and the consequent higher cooling rate that occurs during powder-based, causing the formation of a finer microstructure. Furthermore, it has been reported that the mechanical properties vary along the height in the build [3]. Higher values of hardness have been measured at the top and bottom of the build, which is explained by the different temperature gradients and cooling rates at different positions in the build [27]. A post-build heat treatment was usually used to improve the comprehensive mechanical properties of AM Ti6Al4V alloy by transforming the acicular α′ martensite into equilibrium (α + β) microstructures or reducing the thermal stresses [1, 28-30]. However, the relationship between the microstructure and anisotropic mechanical properties in LDMD AM Ti6Al4V is not fully understood. In-situ SEM tensile testing plays an important role in studying the mechanical properties and deformation behavior by placing a tensile stage into SEM equipment. It can directly reveal the deformation process via real-time monitoring of the microcrack initiation, propagation, and final fracture. Understanding in-depth the real-time relationship between the microstructure and mechanical properties during tensile loading are very helpful for manipulating the mechanical properties and microstructure of the alloy by adjusting the processing parameters. Studies on forged titanium alloys have shown that during tensile loading, shear bands, slip bands, and dislocation movements are significant in the early stages of plastic deformation. By comparison, in the case of LDMD AM titanium alloy, which is composed of large columnar β grains interspersed with α phase, the fracture mechanism differ along the transverse and longitudinal directions when subjected to tensile loading. In this work, the evolution of the microstructure and the tensile fracture processes induced in LDMD Ti6Al4V was examined using in-situ tensile experiments by placing a tensile loading stage within a field-emission SEM (FE-SEM). This study aims to compare the anisotropic tensile properties by real-time observation of the crack initiation and propagation and the concomitant evolution of microstructure using in-situ FE-SEM tensile testing. On the basis of the in-situ FE-SEM tensile results, the relationship between the microstructure and mechanical properties as a function of the location and direction in an AM Ti6Al4V component was constructed. A mechanism that gives rise to the anisotropic mechanical properties is proposed.
2. Experimental details 2.1 Laser direct metal deposition The specimen was fabricated using the LDMD technique directly onto a Ti6Al4V base plate at room temperature of 23℃. An EOS M280 apparatus made by Electro Optical Systems, Germany, was used at the maximum power of 340 W. The apparatus was operated on the basis of the coaxial powder feeding principle. The pre-alloyed Ti6Al4V powder (in wt.%: 6.17 Al, 3.98 V, 0.173 O, 0.018 C, 0.0022 H, 0.009 N, balance Ti) with the particle size of approximately 40 μm was used during LDMD. The coupons were built using an alternating scanning strategy post the deposition
of each layer . In this way, the laser path was rotated by 180° for each consecutive pass within the same layer. The experiment was performed in a protective argon atmosphere to avoid oxidation. The experimental parameters are listed in Table 1. According to the parameters, a cube-shaped component with dimensions of 180 mm × 42 mm × 130 mm was fabricated, as shown in Fig. 1a. The staircase effect demonstrates the layer thickness. Table 1. LDMD deposition parameters used to fabricate Ti6Al4V AM specimens. Deposition parameters Laser power Scan speed Powder bed layer thickness Beam diameter Hatch space Layer thickness Protecting atmosphere flow
170 W 1250 mm/s 30 μm 0.2 mm 5 mm 700 μm 10 L/min
2.2 Characterization The microstructure of the as-deposited samples was examined using Olympus optical microscope (OM) and scanning electron microscopy (SEM, FEI QUANTA FEG650). The specimens used for metallographic analysis and in-situ study were prepared using standard polishing techniques such as grinding and polishing with SiC abrasive papers (from 400# to 2000#) until mirror smooth. The samples were then cleaned with ethanol and isopropyl alcohol and dried at room temperature for 24 h. A solution of Keller’s reagent (HNO3:HCl:HF = 2.5 mL:1.5 mL:1.0 mL) was used as the etchant for the OM and SEM metallography studies. 2.3 Tensile testing To compare the anisotropic mechanical properties in the as-deposited condition, five dog-bone-shaped specimens, each with the long axis along the transverse (X-direction) and longitudinal directions (Z-direction), were fabricated to avoid error. The specimens were machined according to GB/145-2001 by electrical discharge machining as the dimensions shown in Fig. 1b. In order to study the fracture mechanisms in the X- and Z-directions, in-situ tensile tests were performed in an FEI QUANTA FEG650 SEM at room temperature of 23℃. Three small specimens with dimensions of 32 mm × 5 mm × 0.42 mm (shown in the inset in Fig.5) were extracted along the X- and Z-directions respectively. Before testing, the specimens were polished to a mirror-smooth finish to be able to observe the fracture phenomenon clearly. The in-situ tensile stage shown in Figs. 2a and b was developed by the authors. The left- and right-side clamps on the tensile stage are controlled by a multistep gear-drive system, which ensures biaxial stretching along the central axis (Fig. 2c). The tensile loading rate is 1 μm/s. The time interval at which the images were captured is 4 frames/s. The SEM images were taken at pauses in the tensile process.
Fig. 1. (a) Schematic of the LDMD AM Ti6Al4V component and (b) dimensions of the tensile test specimens.
Fig. 2. In-situ SEM tensile system: (a) tensile testing stage mounted on the SEM sample stage, (b) side view of the tensile testing stage inside the microscope and (c) the in-situ setup of the tensile testing.
3. Results 3.1 As-built microstructure Figure 3 shows the OM micrographs of the tensile specimens (Fig. 3a for the X-direction and Fig. 3b for the Z-direction). Coarse columnar primary β grains having lengths much larger than the layer thickness were observed in both the specimens. The formation of large β grains is due to the good wettability between the neighboring layers or between the liquid and substrate because of the similarity in their compositions that, theoretically, reduces the wetting angle to 0º. This implies that there is no nucleation barrier during solidification. The growth direction is along the build direction itself as there is a steep thermal gradient in the liquid at the growth front and the predominant direction of the heat flow is perpendicular to the base plate [16]. Heat is mainly dissipated away along the Z-direction by conduction through any material that has been previously deposited as well as the build plate [24]. Primary β grains decorated with the fine α grains grow approximately parallel to the build direction and through the deposition layers. The distance between two adjacent interfaces is approximately 700 μm, which is consistent with the adjacent
layer thicknesses obtained from the experiments. The LDMD process characterized by the layer-by-layer additive results in different microstructures between neighboring layers since each layer is subjected to its own unique thermal history [11, 31].
Fig. 3. Optical micrographs of the as-built LDMD Ti6Al4V alloy: (a) X-direction, and (b) Z-direction. Figure 4 shows SEM images of the as-received microstructures in the X- and Z-directions. The low-magnification image of the X-direction in Fig. 4a shows a lamellar structure interspersed with the transformation products composed of grain boundary α and colonies of primary α. At the β grain boundaries, the primary α plates have a low angle grain boundary to satisfy the Burgers orientation relationship (BOR) with the prior β phase. An approximately 45° tilting of the α plates toward the deposition direction X was also identified to be a result of epitaxial growth along the steepest temperature gradient. During solidification, the growth rates are also influenced by the crystallography through specific preferred growth directions. For the Ti6Al4V alloy, the favorable growth directions of β phase are the <100> directions since, kinetically, much faster growth occurs here during the solidification [32]. Thus, the nucleation along the <100> direction will outgrow the less favorably oriented sites along the temperature gradient. The α plates parallel to each other constitute an α colony, which is visible in the high-magnification image in Fig. 4b. The colony size is limited by the prior-β grain size because α phases are initially nucleated from the prior-β grain boundaries and subsequently, grow in colonies with the parallel α plates. With increasing cooling rates, the growth of the α colony from the prior-β grain boundaries is not fast enough and therefore, α plates start nucleating at the boundaries of other colonies giving rise to new α colonies. The growth direction of each α colony is in accordance with the BOR. The average thickness of the α plates is approximately 500 nm. The size of the α plates vary randomly within the specimen; thus, a correlation between the colony size and location cannot be established. This result indicates that the cooling rate in the part fabricated for this study was insufficient to form acicular α′ martensite. The absence of martensite may imply an increase in the ductility of the material [33]. The top regions in the specimen oriented in the Z-direction have a different microstructure from that prevalent in regions that are farther down. The five tensile specimens all exhibit a failure point that appears nearly at the middle of the specimen at room temperature. In order to observe the fracture process and deduce the fracture mechanism by in situ tensile testing, therefore, the microstructure was taken from the middle position in the Z-direction specimen. Fig. 4c exhibits a secondary electron morphology composed of the basket weave of α phases. The thickness of the single α plate is comparable to that shown in Fig. 4b at the same magnification because of the identical cooling rates. However, the size of the individual α plates and the colony size are both
much smaller than those of X-direction in Fig. 4b. The distribution with a smaller α plate size and smaller colony arrangement is the result of imperfect sectioning; i.e., the growth direction of the majority of the α plates is not coherent, but show an inclination towards the direction of deposition. As a result, the side view shows the formation of shorter α plates with a smaller colony distribution. Each α plate becomes a separate grain in the final solidified structure.
Fig. 4. SEM microstructures of as-deposited LDMD Ti6Al4V components: (a) low-magnification image of the front view showing the grain boundary α nucleating along the β grain boundary, (b) local magnification showing the large α colony and (c) side view showing the basket-weave structure. 3.2 As-built tensile properties Five specimens each, oriented along the X- and Z-directions respectively, were subjected to tensile testing. The measured mechanical properties at 23℃ are tabulated in Table 2. Differences in the data are observed to be a function of specimen location and orientation. The average tensile and yield strengths of the five specimens are approximately 993 and 923 MPa in the X-direction and approximately 957 and 868 MPa in the Z-direction, respectively. The average elongation and reduction of cross section along the Z-direction are measured as approximately 14 and 34%, respectively. Along the X-direction, the respective lower values of 8 and 18% indicate the anisotropy in mechanical properties. Leuders et al conjectured that the low ductility in the X-direction can be attributed to the presence of lack-of-fusion porosity [34]. In the present study, porosity was observed neither in the X-direction specimens nor in Z-direction specimens. It may be concluded that the underlying mechanism for the anisotropy in mechanical properties between X- and Z-directions in this study is directly related to the anisotropic microstructure of the LDMD AM build. Table 2. Mechanical properties of the as-built LDMD Ti6Al4V measured along the X- and Z-directions. Specimens
Tensile strength, MPa
Yield strength, MPa
Elongation, %
Reduction of cross section, %
X1 X2 X3 X4 X5 Average Z1 Z2
993 993 996 986 996 992.8 958 930
925 924 922 918 924 922.6 871 847
8.5 6.0 10.0 7.5 8.0 8 12.0 15.0
18 13 20 24 16 18.2 27 38
Z3 Z4 Z5 Average
962 966 968 956.8
867 872 883 868
13.0 16.0 13.0 13.8
38 40 28 34.2
3.3 In-situ tensile testing To elucidate the tensile fracture behavior of the as-built LDMD Ti6Al4V alloy along the X- and Z-directions, we investigated the crack initiation and growth process using in-situ tensile testing. Figure 5 shows the force-displacement curves along the X- and Z-directions, and the inset shows the in-situ tensile specimen size. The stepwise appearance of the curves is mainly due to the relaxation of the specimen and due to the intermittent starting and stopping of the experiment to enable acquisition of SEM images at various stages. From the tensile curves, the X-direction specimen exhibits a maximum force of 716 N and a maximum displacement of 1.368 mm. The maximum force and displacement in the Z-direction specimen are 707 N and 1.440 mm, respectively. Notably, the tendency of high displacement in the Z-direction and the high force in the X-direction agree with that of the high elongation in the Z-direction and the high strength in the X-direction. The difference in the values is caused by the different specimen size.
Fig. 5. Force-displacement curves from the in-situ tensile experiments (the inset shows the in-situ tensile specimen size). 3.4 In-situ tensile deformation process Figure 6 shows the in-situ tensile deformation process for an X-direction specimen with tensile axis parallel to the scanning direction. When the displacement reached 1.011 mm (point a on the red line in Fig. 5), some phenomena are observed along the grain boundary α, as shown in Fig. 6(a). However, the deformation in the specimen was relatively uniform. As the displacement increased to 1.165 mm (point b on the red line in Fig. 5), necking occurred, accompanied by shear deformation, as shown in Fig. 6(b), due to the plastic deformation resulting from exceeding the yield strength. At the free surface of the top part, microcracks appeared along the α plates due to stress concentrations. Notably, the relative shear motion along the cracks on the top was also
observed at an angle of approximately 45°. The plastic flow and shear motion induce chaotic grain-boundary tortuosity within the plastic deformation zone. As the displacement further increased to 1.309 mm (point c on the red line in Fig. 5), the severe shear motion led to fracture along the shear direction, as shown in Fig. 6(c). The fracture mode is trans-crystalline with an inclination of approximately 45° between the fractured surface and tensile axis, which is consistent with the direction of the shear motion. Ductile fracture characteristics with tortuous grain boundary deformation (Fig. 6c) and deep equi-axed dimples (Fig. 6d) were also identified. In addition, the fracture surface seems to follow the prior β grain boundary, as visible in Fig. 6c. The elongation is believed to stem from the traces of prior β grain boundaries, which would constitute the crack propagation path as the material fractures [35]. From the fracture surface profile in Fig. 6c, elongated features resembling the shape of the columnar prior β grains can be observed. The fractured structure that was composed of the fully dense material did not reveal the presence of pores or unmolten areas, suggesting suitable overlap between the passes and sufficient filling of the melted material. Therefore, it can be concluded that porosity did not play a substantial role in determining tensile fracture properties in the transverse specimen. The shear motion of the fractured surface and the elongated features in the profile suggest lower values of elongation.
Fig. 6. In-situ tensile fracture process for an X-direction specimen: (a) a displacement of 1.011 mm, (b) a displacement of 1.165 mm, (c) a displacement of 1.309 mm and (d) fracture morphology. To explain why the Z-direction specimen has a higher elongation and lower tensile strength than the X-direction specimen, in-situ tensile testing of the Z-direction specimen with the tensile axis aligned along the build direction was conducted. The results are shown in Fig. 7. At a displacement of 1.067 mm (point a on the Z-direction force-displacement curve in Fig. 5), initial necking was observed, with the necking zone showing homogeneous deformation (Fig. 7a). As the displacement increased to 1.279 mm (point b on the Z-direction curve in Fig. 5), homogenous deformation along with necking and microstructural changes were observed as shown in Fig. 7b. The crack was also first initiated along the top edge. The ambient α plates caused the extension of homogeneously tortuous grain boundary and plastic flow, showing ductile characteristics. As the
displacement further increased to 1.342 mm (point c on the Z-direction curve in Fig. 5), a trans-crystalline fracture occurred with smooth cup-cone shaped fracture morphology (Fig. 7c) and deep dimples (Fig. 7d). In addition, the cross section of the fracture surface was very flat and smooth (Fig. 7c). This is because the orientation of the prior β grains is along the horizontal direction; therefore, the fracture surface did not show any cracks or microstructural elongation. The observations of the fracture in the Z-direction, with deep dimples and cup-cone fracture surfaces, support the higher values of elongation.
Fig. 7. In-situ tensile fracture process for a Z-direction specimen: (a) a displacement of 1.067 mm, (b) a displacement of 1.279 mm, (c) a displacement of 1.342 mm and (d) fracture morphology.
4. Discussion Observations made in this study show that the location and direction of specimens extracted from the AM build strongly affect the microstructure and tensile properties of laser-processed Ti6Al4V. The effect of pores, as mentioned in many other studies, often results in limited ductility; these were not observed in the present study and hence, are not discussed. 4.1 Effect of columnar β-grains’ direction on mechanical properties Ti6Al4V is an (α+β) alloy and is especially sensitive to its thermal history. Layer-by-layer manufacturing introduces the columnar prior-β grains with a length of several millimeters growing in the build direction. The transverse-extracted specimens in this study were tested with the columnar β grains approximatively perpendicular to the tensile axis; the results showed high strength and low elongation. When the longitudinal specimens were tested, the stress axis was approximately parallel to the direction of growth of the columnar β grains, and the result was high elongation and low strength. This result is consistent with many other studies, that recorded high elongation and reduction of cross section in the perpendicular specimens, and a high yield strength and ultimate tensile strength in parallel specimens [12, 33, 36]. The difference in tensile properties is mainly because the growth direction of large columnar β grains is approximately perpendicular to the tensile axis in the transverse specimen, which is a result of the short axes of the prior-β
grains boundaries and the grain boundary α phases being subjected to the tensile load. Grain-boundary α is a soft, thin continuous layer that grows along the prior-β grain boundary. As schematically explained by Carroll et al. [12], it will be more detrimental to the mechanical properties of X-direction specimens. This is because, in this direction, the tensile stress acts as the separating force on the grain boundary α. More prior-β grains boundaries in the fracture surface (see in Fig. 6c) could also explain the resistance to fracture. However, in the longitudinal specimen where the tensile direction is along the growth direction for the large columnar β grains, the long axes of the prior-β grain boundaries and grain-boundary α phases are subjected to the tensile load. The number of grain boundaries in a cross-section of the specimens will differ from that in the X-direction specimens. Therefore, the long axes of the prior-β grain boundaries and grain boundary α serve as a path where damage can accumulate preferentially, leading to failure at a low stress and high elongation. 4.2 Effect of α colonies and their crystal orientation on the mechanical properties According to Lütjering [37], in the case of titanium alloys with a lamellar microstructure, the α colony size is the most important factor that influences the mechanical properties because it determines the effective slip length [1, 23, 37, 38]. However, the α colony size is similar for the two specimen orientations in the same LDMD Ti6Al4V build and therefore, the anisotropy in mechanical properties cannot be attributed to the α colony size. Due to the effect of thermal gradient, the growing direction of prior β grains and thereby, the grain boundary α, will deflect from the vertical direction [12]. The large columnar prior β grains cause a significant difference in the number of grain boundaries in a cross-section of the specimen in X- and Z-direction because of the orientation effect. As per the result, the size of α colonies and single α plate present the difference along the X- and Z-direction. Therefore, the α colonies and their crystal orientation can be assumed to be the main reason for the anisotropy in mechanical properties [39]. In the X-direction specimen, big α-colony size with long α plates (see Fig. 4b) and more prior β grains (see the fracture profile in Fig. 6c) cause the initial microcrack formation along the α plate boundaries at a displacement of 1.011 mm as shown in Fig. 8a. Microcracks usually nucleate and grow relatively easily along the primary α phase boundary because the homogenous microstructure there presents very few obstacles. The observation of crack initiation and propagation along the primary α phase boundary also corroborates the detrimental influence of grain boundary α on tensile elongation. However, Lütjering [1, 37] speculated that the colony boundaries are strong barriers that inhibit cracks from passing through and propagating. Therefore, a large α colony is beneficial for improving the strength by resisting the propagation of microcracks due to increased crack roughness and crack closure [37]. At a displacement of 1.309 mm (Fig. 8b), the severe shear motion causes the boundaries of the α plates to offer little resistance to deformation. A large number of microcracks appear across one or several α plates. Some of the fractured α plates exhibit stepped staircase-like features. This behavior is analogous to the cleavage fracture usually observed for materials with a higher grain-boundary fraction. The fracture surface in Fig. 8c also shows trans-crystalline fracture characteristics.
Fig. 8 Microstructures of the X-direction specimen during in-situ tensile loading: (a) microcrack initiation at a displacement of 1.011 mm, (b) α-plate shear fracturing at a displacement of 1.309 mm, and (c) transcrytalline fracture. In the Z-direction specimen, because of α colonies and their crystal orientation, the fine basket-woven structure (see Fig. 4c) with smaller α colonies decreases the slip length. Therefore, more uniform deformation is observed at a displacement of 1.067 mm (Fig. 9a). The deformation during tensile testing leads to substantial changes in the microstructure. These α plates are elongated and bent, indicating high plasticity. As the displacement increases to 1.279 mm, a crack is observed at the α-plate boundaries because of uniaxial tension (Fig. 9b). The sharp angles of the microcrack cause a local stress concentration and play substantial roles in the initial fracture by promoting crack initiation under dynamic tensile loading, particularly for Z-direction tensile loading [29]. The fracture in Fig. 9c also shows the shearing of unfavorably oriented α plates.
Fig. 9 Microstructures of the Z-direction specimen during in-situ tensile loading: (a) uniform deformation at a displacement of 1.067 mm, (b) microcracking along the α plates at a displacement of 1.279 mm, and (c) transcrystalline fracture.
5. Conclusions The anisotropic mechanical properties of LDMD AM Ti6Al4V components were compared for transverse and longitudinal orientations. The fracture mechanisms were clarified on the basis of microstructural distributions and fracture processes, as tested by in-situ SEM tensile loading. The conclusions are summarized as follows: 1) LDMD AM produced compact and defect-free Ti6Al4V components. The anisotropic mechanical properties were characterized with high strength and low elongation in the transverse direction and high elongation and reduction of area in the perpendicular direction. The anisotropic mechanical properties are caused by the anisotropic microstructural distribution. 2) The microstructure in the transverse tensile specimen was composed of α phases having soft grain boundary, and short and thin columnar prior-β grains that grew approximately perpendicular to the tensile axis, rendering the tensile strength and fracture strength significantly higher than in the longitudinal direction. However, the longitudinally oriented tensile specimen consists of long columnar prior-β grains parallel to the tensile axis and small α colonies with short α plates, which
result in the preferential accumulation of damage. The results of the study indicate that the longitudinal specimen shows a low stress and high elongation. 3) The anisotropic tensile properties were also examined by in-situ tensile testing. The transverse specimen fractured along the prior-β grain boundary with plenty of shear bands. The vertical β-grain boundaries and large α colonies contribute to the high strength by resisting microcrack propagation. The longitudinal specimen fractured with symmetric necking due to the fine α grains, which results in uniform deformation of the grains on both sides of the grain boundaries.
Acknowledgements The authors wish to thank the National Natural Science Foundation of China (Grant number: 51505010) and Beijing Natural Science Foundation (Grant number: 2152007)
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PII: DOI: Reference:
S0921-5093(17)31567-8 https://doi.org/10.1016/j.msea.2017.11.106 MSA35821
To appear in: Materials Science & Engineering A Received date: 3 August 2017 Revised date: 25 November 2017 Accepted date: 25 November 2017 Cite this article as: Junxia Lu, Ling Chang, Jin Wang, Lijun Sang, Shikai Wu and Yuefei Zhang, In-situ investigation of the anisotropic mechanical properties of laser direct metal deposition Ti6Al4V alloy, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2017.11.106 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
In-situ investigation of the anisotropic mechanical properties of laser direct metal deposition Ti6Al4V alloy Junxia Lua*, Ling Changa, Jin Wangb, Lijun Sangb, Shikai Wua, Yuefei Zhangb** a Institute of Laser Engineering, Beijing University of Technology, Chao Yang District, Ping Le Yuan 100#, Beijing, 100124, China b Institute of Microstructure and Property of Advanced Materials, Chao Yang District, Ping Le Yuan 100#, Beijing, 100124, China E-mail: [email protected] E-mail: [email protected] *Corresponding author: Junxia Lu, Tel.: +86 010-67396557 Fax: +86-10-67392635 **Corresponding author: Yuefei Zhang, Tel: +86-10-67392635 Fax: +86-10-67392635 Abstract This study compares the microstructure and tensile properties of Ti6Al4V components fabricated by laser direct metal deposition (LDMD) additive manufacturing (AM) in the transverse and longitudinal directions. The results show anisotropic tensile properties with the transverse direction having high tensile and fracture strengths and the longitudinal direction having a high elongation and reduction of cross section. The anisotropic mechanical properties are attributed to the anisotropic microstructural distribution. The transverse tensile specimen is composed of short columnar prior-β grains which grow perpendicular to the tensile direction, and have a lamellar structure. Along the β grain boundary, αGB and large α colonies were identified. However, the longitudinal specimen shows that the long β structure is parallel to the tensile axis and that the microstructure is composed of basket-woven α phases with shorter α plates and smaller colony sizes compared with those in the transverse specimen. The fracture mechanism induced by the anisotropic microstructure along the transverse and longitudinal directions was compared by examining the fracture process in real-time using uniaxial in-situ scanning electron microscopy (SEM) tensile testing. The results show that shear fracture, which is caused by the vertical β grain boundaries and large α colonies with long α plates, occurs in the transverse specimen. The shear mode is the main reason behind the enhanced tensile strength and fracture strength due to the high resistance to microcrack propagation. However, in the longitudinal specimens, symmetric necking behavior due to the fine α grains resulted in uniform deformation of the grains on both sides of the grain boundaries, inducing greater elongation. Key words: Additive manufacturing; Titanium alloy; Microstructure; Mechanical properties; In-situ tensile
1. Introduction The Ti6Al4V (wt.%) alloy represents about 50% of the titanium alloy market, with very large quantities being used by the aerospace industry because of its exceptionally high specific strength, good cryogenic properties, and excellent corrosion resistance [1-4]. Parts made of the alloy are
traditionally machined by die forging, hand forging or plate stock. However, these methods cause substantial wastage of material, especially when the geometries involved are complex, resulting in high buy-to-fly ratios [1]. For these types of components, laser near-net-shape additive manufacturing (AM) offers several potential benefits, including an increased raw material utilization ratio and a reduced number of machining operations with a consequential decrease in the cost and lead time [5]. Laser direct metal deposition (LDMD) is a near-net-shape AM technology in which successive layers of metal powder, wires or ribbon are melted onto a workpiece. In this process, a focused laser heat source is used to melt the metallic feedstock and produce a desired shape. Materials can be delivered to the workpiece either by first spreading the individual power layers and then, melting the individual layers selectively in a powder-bed fusion process, or by continuous direction through coaxial nozzles to the melt pool in a directed-energy deposition process. LDMD offers the ability to fabricate fully dense three-dimensional parts with highly complex geometric features and lightweight structures using low-density alloys, thereby resulting in a shorter manufacturing time, less raw material waste, and lower cost [6, 7]. Conventional manufacturing techniques, such as rolling and forging are associated with material deformation and rely on process control and material properties to produce metallic components. The mechanical properties of the product are determined by the process parameters. In comparison, components produced using LDMD AM differ substantially from those produced using conventional manufacturing methods. The final mechanical properties are affected not only by the process parameters but also the solidification behavior of the material. The LDMD AM process involves rapid melting and solidification coupled with high supercooling at the solid/liquid interface. Repetitive thermal cycling by the reheating of the previously deposited layer induces the phase transformation of Ti6Al4V alloy from the body centered cubic (bcc) β-phase to face centered cubic (fcc) α-phase. The steep temperature gradients and repeated heating lead to macroscopically epitaxial grain growth [8-10]. A general observation in earlier studies was the presence of large columnar prior-β grain structures parallel to the build direction when selective laser melting (SLM) deposition is used [11-14]; these structures were also observed when the electron-beam (EB) AM [15-17] and wire + arc AM [18] techniques were used. The width of these prior-β grains depend on the cooling rate during solidification. For instance, the width of the columnar prior-β grain in Ti6Al4V alloy when EB melting was used is reported to be about 25-30 μm [19] while it is 220-2450 μm for laser beam deposition [20]. The coarse columnar β grains that grow from the substrate and extend across multiple layers often tend to favor intergranular failure [11] and the vertical columnar growth along the build direction causes anisotropic mechanical properties in the deposition direction and build direction [12，21]. Within the prior-β grains, a significant amount of acicular α' martensite or lamellar α structure can be observed; the quantities depend on the cooling rate [22]. Some authors also observed grain boundary α phases which distributed along the prior-β grain boundary [3, 18]. However, Bermingham et al [23] report that the grain boundary α was eliminated and the prior-β size decreased by the trace additions of boron to the Ti6Al4V feedstock material. Such nonhomogeneous microstructural distribution within a single component often leads to difficulties in characterizing and predicting the final structural performance. This is particularly relevant when designing parts for aircraft assembly application using AM technique [24]. The anisotropic tensile properties of AM Ti6Al4V were tested and the results showed greater
elongation and reduction of cross section in the perpendicular direction while higher values of yield strength and ultimate tensile strength were observed in the deposition direction [8, 11]. Some reports partially attribute the high elongation in the perpendicular orientation to the less porous components [25, 26]. In addition, the anisotropy is less pronounced for the powder-based processes compared with that for wire-based processes because of the lower heat input and the consequent higher cooling rate that occurs during powder-based, causing the formation of a finer microstructure. Furthermore, it has been reported that the mechanical properties vary along the height in the build [3]. Higher values of hardness have been measured at the top and bottom of the build, which is explained by the different temperature gradients and cooling rates at different positions in the build [27]. A post-build heat treatment was usually used to improve the comprehensive mechanical properties of AM Ti6Al4V alloy by transforming the acicular α′ martensite into equilibrium (α + β) microstructures or reducing the thermal stresses [1, 28-30]. However, the relationship between the microstructure and anisotropic mechanical properties in LDMD AM Ti6Al4V is not fully understood. In-situ SEM tensile testing plays an important role in studying the mechanical properties and deformation behavior by placing a tensile stage into SEM equipment. It can directly reveal the deformation process via real-time monitoring of the microcrack initiation, propagation, and final fracture. Understanding in-depth the real-time relationship between the microstructure and mechanical properties during tensile loading are very helpful for manipulating the mechanical properties and microstructure of the alloy by adjusting the processing parameters. Studies on forged titanium alloys have shown that during tensile loading, shear bands, slip bands, and dislocation movements are significant in the early stages of plastic deformation. By comparison, in the case of LDMD AM titanium alloy, which is composed of large columnar β grains interspersed with α phase, the fracture mechanism differ along the transverse and longitudinal directions when subjected to tensile loading. In this work, the evolution of the microstructure and the tensile fracture processes induced in LDMD Ti6Al4V was examined using in-situ tensile experiments by placing a tensile loading stage within a field-emission SEM (FE-SEM). This study aims to compare the anisotropic tensile properties by real-time observation of the crack initiation and propagation and the concomitant evolution of microstructure using in-situ FE-SEM tensile testing. On the basis of the in-situ FE-SEM tensile results, the relationship between the microstructure and mechanical properties as a function of the location and direction in an AM Ti6Al4V component was constructed. A mechanism that gives rise to the anisotropic mechanical properties is proposed.
2. Experimental details 2.1 Laser direct metal deposition The specimen was fabricated using the LDMD technique directly onto a Ti6Al4V base plate at room temperature of 23℃. An EOS M280 apparatus made by Electro Optical Systems, Germany, was used at the maximum power of 340 W. The apparatus was operated on the basis of the coaxial powder feeding principle. The pre-alloyed Ti6Al4V powder (in wt.%: 6.17 Al, 3.98 V, 0.173 O, 0.018 C, 0.0022 H, 0.009 N, balance Ti) with the particle size of approximately 40 μm was used during LDMD. The coupons were built using an alternating scanning strategy post the deposition
of each layer . In this way, the laser path was rotated by 180° for each consecutive pass within the same layer. The experiment was performed in a protective argon atmosphere to avoid oxidation. The experimental parameters are listed in Table 1. According to the parameters, a cube-shaped component with dimensions of 180 mm × 42 mm × 130 mm was fabricated, as shown in Fig. 1a. The staircase effect demonstrates the layer thickness. Table 1. LDMD deposition parameters used to fabricate Ti6Al4V AM specimens. Deposition parameters Laser power Scan speed Powder bed layer thickness Beam diameter Hatch space Layer thickness Protecting atmosphere flow
170 W 1250 mm/s 30 μm 0.2 mm 5 mm 700 μm 10 L/min
2.2 Characterization The microstructure of the as-deposited samples was examined using Olympus optical microscope (OM) and scanning electron microscopy (SEM, FEI QUANTA FEG650). The specimens used for metallographic analysis and in-situ study were prepared using standard polishing techniques such as grinding and polishing with SiC abrasive papers (from 400# to 2000#) until mirror smooth. The samples were then cleaned with ethanol and isopropyl alcohol and dried at room temperature for 24 h. A solution of Keller’s reagent (HNO3:HCl:HF = 2.5 mL:1.5 mL:1.0 mL) was used as the etchant for the OM and SEM metallography studies. 2.3 Tensile testing To compare the anisotropic mechanical properties in the as-deposited condition, five dog-bone-shaped specimens, each with the long axis along the transverse (X-direction) and longitudinal directions (Z-direction), were fabricated to avoid error. The specimens were machined according to GB/145-2001 by electrical discharge machining as the dimensions shown in Fig. 1b. In order to study the fracture mechanisms in the X- and Z-directions, in-situ tensile tests were performed in an FEI QUANTA FEG650 SEM at room temperature of 23℃. Three small specimens with dimensions of 32 mm × 5 mm × 0.42 mm (shown in the inset in Fig.5) were extracted along the X- and Z-directions respectively. Before testing, the specimens were polished to a mirror-smooth finish to be able to observe the fracture phenomenon clearly. The in-situ tensile stage shown in Figs. 2a and b was developed by the authors. The left- and right-side clamps on the tensile stage are controlled by a multistep gear-drive system, which ensures biaxial stretching along the central axis (Fig. 2c). The tensile loading rate is 1 μm/s. The time interval at which the images were captured is 4 frames/s. The SEM images were taken at pauses in the tensile process.
Fig. 1. (a) Schematic of the LDMD AM Ti6Al4V component and (b) dimensions of the tensile test specimens.
Fig. 2. In-situ SEM tensile system: (a) tensile testing stage mounted on the SEM sample stage, (b) side view of the tensile testing stage inside the microscope and (c) the in-situ setup of the tensile testing.
3. Results 3.1 As-built microstructure Figure 3 shows the OM micrographs of the tensile specimens (Fig. 3a for the X-direction and Fig. 3b for the Z-direction). Coarse columnar primary β grains having lengths much larger than the layer thickness were observed in both the specimens. The formation of large β grains is due to the good wettability between the neighboring layers or between the liquid and substrate because of the similarity in their compositions that, theoretically, reduces the wetting angle to 0º. This implies that there is no nucleation barrier during solidification. The growth direction is along the build direction itself as there is a steep thermal gradient in the liquid at the growth front and the predominant direction of the heat flow is perpendicular to the base plate [16]. Heat is mainly dissipated away along the Z-direction by conduction through any material that has been previously deposited as well as the build plate [24]. Primary β grains decorated with the fine α grains grow approximately parallel to the build direction and through the deposition layers. The distance between two adjacent interfaces is approximately 700 μm, which is consistent with the adjacent
layer thicknesses obtained from the experiments. The LDMD process characterized by the layer-by-layer additive results in different microstructures between neighboring layers since each layer is subjected to its own unique thermal history [11, 31].
Fig. 3. Optical micrographs of the as-built LDMD Ti6Al4V alloy: (a) X-direction, and (b) Z-direction. Figure 4 shows SEM images of the as-received microstructures in the X- and Z-directions. The low-magnification image of the X-direction in Fig. 4a shows a lamellar structure interspersed with the transformation products composed of grain boundary α and colonies of primary α. At the β grain boundaries, the primary α plates have a low angle grain boundary to satisfy the Burgers orientation relationship (BOR) with the prior β phase. An approximately 45° tilting of the α plates toward the deposition direction X was also identified to be a result of epitaxial growth along the steepest temperature gradient. During solidification, the growth rates are also influenced by the crystallography through specific preferred growth directions. For the Ti6Al4V alloy, the favorable growth directions of β phase are the <100> directions since, kinetically, much faster growth occurs here during the solidification [32]. Thus, the nucleation along the <100> direction will outgrow the less favorably oriented sites along the temperature gradient. The α plates parallel to each other constitute an α colony, which is visible in the high-magnification image in Fig. 4b. The colony size is limited by the prior-β grain size because α phases are initially nucleated from the prior-β grain boundaries and subsequently, grow in colonies with the parallel α plates. With increasing cooling rates, the growth of the α colony from the prior-β grain boundaries is not fast enough and therefore, α plates start nucleating at the boundaries of other colonies giving rise to new α colonies. The growth direction of each α colony is in accordance with the BOR. The average thickness of the α plates is approximately 500 nm. The size of the α plates vary randomly within the specimen; thus, a correlation between the colony size and location cannot be established. This result indicates that the cooling rate in the part fabricated for this study was insufficient to form acicular α′ martensite. The absence of martensite may imply an increase in the ductility of the material [33]. The top regions in the specimen oriented in the Z-direction have a different microstructure from that prevalent in regions that are farther down. The five tensile specimens all exhibit a failure point that appears nearly at the middle of the specimen at room temperature. In order to observe the fracture process and deduce the fracture mechanism by in situ tensile testing, therefore, the microstructure was taken from the middle position in the Z-direction specimen. Fig. 4c exhibits a secondary electron morphology composed of the basket weave of α phases. The thickness of the single α plate is comparable to that shown in Fig. 4b at the same magnification because of the identical cooling rates. However, the size of the individual α plates and the colony size are both
much smaller than those of X-direction in Fig. 4b. The distribution with a smaller α plate size and smaller colony arrangement is the result of imperfect sectioning; i.e., the growth direction of the majority of the α plates is not coherent, but show an inclination towards the direction of deposition. As a result, the side view shows the formation of shorter α plates with a smaller colony distribution. Each α plate becomes a separate grain in the final solidified structure.
Fig. 4. SEM microstructures of as-deposited LDMD Ti6Al4V components: (a) low-magnification image of the front view showing the grain boundary α nucleating along the β grain boundary, (b) local magnification showing the large α colony and (c) side view showing the basket-weave structure. 3.2 As-built tensile properties Five specimens each, oriented along the X- and Z-directions respectively, were subjected to tensile testing. The measured mechanical properties at 23℃ are tabulated in Table 2. Differences in the data are observed to be a function of specimen location and orientation. The average tensile and yield strengths of the five specimens are approximately 993 and 923 MPa in the X-direction and approximately 957 and 868 MPa in the Z-direction, respectively. The average elongation and reduction of cross section along the Z-direction are measured as approximately 14 and 34%, respectively. Along the X-direction, the respective lower values of 8 and 18% indicate the anisotropy in mechanical properties. Leuders et al conjectured that the low ductility in the X-direction can be attributed to the presence of lack-of-fusion porosity [34]. In the present study, porosity was observed neither in the X-direction specimens nor in Z-direction specimens. It may be concluded that the underlying mechanism for the anisotropy in mechanical properties between X- and Z-directions in this study is directly related to the anisotropic microstructure of the LDMD AM build. Table 2. Mechanical properties of the as-built LDMD Ti6Al4V measured along the X- and Z-directions. Specimens
Tensile strength, MPa
Yield strength, MPa
Elongation, %
Reduction of cross section, %
X1 X2 X3 X4 X5 Average Z1 Z2
993 993 996 986 996 992.8 958 930
925 924 922 918 924 922.6 871 847
8.5 6.0 10.0 7.5 8.0 8 12.0 15.0
18 13 20 24 16 18.2 27 38
Z3 Z4 Z5 Average
962 966 968 956.8
867 872 883 868
13.0 16.0 13.0 13.8
38 40 28 34.2
3.3 In-situ tensile testing To elucidate the tensile fracture behavior of the as-built LDMD Ti6Al4V alloy along the X- and Z-directions, we investigated the crack initiation and growth process using in-situ tensile testing. Figure 5 shows the force-displacement curves along the X- and Z-directions, and the inset shows the in-situ tensile specimen size. The stepwise appearance of the curves is mainly due to the relaxation of the specimen and due to the intermittent starting and stopping of the experiment to enable acquisition of SEM images at various stages. From the tensile curves, the X-direction specimen exhibits a maximum force of 716 N and a maximum displacement of 1.368 mm. The maximum force and displacement in the Z-direction specimen are 707 N and 1.440 mm, respectively. Notably, the tendency of high displacement in the Z-direction and the high force in the X-direction agree with that of the high elongation in the Z-direction and the high strength in the X-direction. The difference in the values is caused by the different specimen size.
Fig. 5. Force-displacement curves from the in-situ tensile experiments (the inset shows the in-situ tensile specimen size). 3.4 In-situ tensile deformation process Figure 6 shows the in-situ tensile deformation process for an X-direction specimen with tensile axis parallel to the scanning direction. When the displacement reached 1.011 mm (point a on the red line in Fig. 5), some phenomena are observed along the grain boundary α, as shown in Fig. 6(a). However, the deformation in the specimen was relatively uniform. As the displacement increased to 1.165 mm (point b on the red line in Fig. 5), necking occurred, accompanied by shear deformation, as shown in Fig. 6(b), due to the plastic deformation resulting from exceeding the yield strength. At the free surface of the top part, microcracks appeared along the α plates due to stress concentrations. Notably, the relative shear motion along the cracks on the top was also
observed at an angle of approximately 45°. The plastic flow and shear motion induce chaotic grain-boundary tortuosity within the plastic deformation zone. As the displacement further increased to 1.309 mm (point c on the red line in Fig. 5), the severe shear motion led to fracture along the shear direction, as shown in Fig. 6(c). The fracture mode is trans-crystalline with an inclination of approximately 45° between the fractured surface and tensile axis, which is consistent with the direction of the shear motion. Ductile fracture characteristics with tortuous grain boundary deformation (Fig. 6c) and deep equi-axed dimples (Fig. 6d) were also identified. In addition, the fracture surface seems to follow the prior β grain boundary, as visible in Fig. 6c. The elongation is believed to stem from the traces of prior β grain boundaries, which would constitute the crack propagation path as the material fractures [35]. From the fracture surface profile in Fig. 6c, elongated features resembling the shape of the columnar prior β grains can be observed. The fractured structure that was composed of the fully dense material did not reveal the presence of pores or unmolten areas, suggesting suitable overlap between the passes and sufficient filling of the melted material. Therefore, it can be concluded that porosity did not play a substantial role in determining tensile fracture properties in the transverse specimen. The shear motion of the fractured surface and the elongated features in the profile suggest lower values of elongation.
Fig. 6. In-situ tensile fracture process for an X-direction specimen: (a) a displacement of 1.011 mm, (b) a displacement of 1.165 mm, (c) a displacement of 1.309 mm and (d) fracture morphology. To explain why the Z-direction specimen has a higher elongation and lower tensile strength than the X-direction specimen, in-situ tensile testing of the Z-direction specimen with the tensile axis aligned along the build direction was conducted. The results are shown in Fig. 7. At a displacement of 1.067 mm (point a on the Z-direction force-displacement curve in Fig. 5), initial necking was observed, with the necking zone showing homogeneous deformation (Fig. 7a). As the displacement increased to 1.279 mm (point b on the Z-direction curve in Fig. 5), homogenous deformation along with necking and microstructural changes were observed as shown in Fig. 7b. The crack was also first initiated along the top edge. The ambient α plates caused the extension of homogeneously tortuous grain boundary and plastic flow, showing ductile characteristics. As the
displacement further increased to 1.342 mm (point c on the Z-direction curve in Fig. 5), a trans-crystalline fracture occurred with smooth cup-cone shaped fracture morphology (Fig. 7c) and deep dimples (Fig. 7d). In addition, the cross section of the fracture surface was very flat and smooth (Fig. 7c). This is because the orientation of the prior β grains is along the horizontal direction; therefore, the fracture surface did not show any cracks or microstructural elongation. The observations of the fracture in the Z-direction, with deep dimples and cup-cone fracture surfaces, support the higher values of elongation.
Fig. 7. In-situ tensile fracture process for a Z-direction specimen: (a) a displacement of 1.067 mm, (b) a displacement of 1.279 mm, (c) a displacement of 1.342 mm and (d) fracture morphology.
4. Discussion Observations made in this study show that the location and direction of specimens extracted from the AM build strongly affect the microstructure and tensile properties of laser-processed Ti6Al4V. The effect of pores, as mentioned in many other studies, often results in limited ductility; these were not observed in the present study and hence, are not discussed. 4.1 Effect of columnar β-grains’ direction on mechanical properties Ti6Al4V is an (α+β) alloy and is especially sensitive to its thermal history. Layer-by-layer manufacturing introduces the columnar prior-β grains with a length of several millimeters growing in the build direction. The transverse-extracted specimens in this study were tested with the columnar β grains approximatively perpendicular to the tensile axis; the results showed high strength and low elongation. When the longitudinal specimens were tested, the stress axis was approximately parallel to the direction of growth of the columnar β grains, and the result was high elongation and low strength. This result is consistent with many other studies, that recorded high elongation and reduction of cross section in the perpendicular specimens, and a high yield strength and ultimate tensile strength in parallel specimens [12, 33, 36]. The difference in tensile properties is mainly because the growth direction of large columnar β grains is approximately perpendicular to the tensile axis in the transverse specimen, which is a result of the short axes of the prior-β
grains boundaries and the grain boundary α phases being subjected to the tensile load. Grain-boundary α is a soft, thin continuous layer that grows along the prior-β grain boundary. As schematically explained by Carroll et al. [12], it will be more detrimental to the mechanical properties of X-direction specimens. This is because, in this direction, the tensile stress acts as the separating force on the grain boundary α. More prior-β grains boundaries in the fracture surface (see in Fig. 6c) could also explain the resistance to fracture. However, in the longitudinal specimen where the tensile direction is along the growth direction for the large columnar β grains, the long axes of the prior-β grain boundaries and grain-boundary α phases are subjected to the tensile load. The number of grain boundaries in a cross-section of the specimens will differ from that in the X-direction specimens. Therefore, the long axes of the prior-β grain boundaries and grain boundary α serve as a path where damage can accumulate preferentially, leading to failure at a low stress and high elongation. 4.2 Effect of α colonies and their crystal orientation on the mechanical properties According to Lütjering [37], in the case of titanium alloys with a lamellar microstructure, the α colony size is the most important factor that influences the mechanical properties because it determines the effective slip length [1, 23, 37, 38]. However, the α colony size is similar for the two specimen orientations in the same LDMD Ti6Al4V build and therefore, the anisotropy in mechanical properties cannot be attributed to the α colony size. Due to the effect of thermal gradient, the growing direction of prior β grains and thereby, the grain boundary α, will deflect from the vertical direction [12]. The large columnar prior β grains cause a significant difference in the number of grain boundaries in a cross-section of the specimen in X- and Z-direction because of the orientation effect. As per the result, the size of α colonies and single α plate present the difference along the X- and Z-direction. Therefore, the α colonies and their crystal orientation can be assumed to be the main reason for the anisotropy in mechanical properties [39]. In the X-direction specimen, big α-colony size with long α plates (see Fig. 4b) and more prior β grains (see the fracture profile in Fig. 6c) cause the initial microcrack formation along the α plate boundaries at a displacement of 1.011 mm as shown in Fig. 8a. Microcracks usually nucleate and grow relatively easily along the primary α phase boundary because the homogenous microstructure there presents very few obstacles. The observation of crack initiation and propagation along the primary α phase boundary also corroborates the detrimental influence of grain boundary α on tensile elongation. However, Lütjering [1, 37] speculated that the colony boundaries are strong barriers that inhibit cracks from passing through and propagating. Therefore, a large α colony is beneficial for improving the strength by resisting the propagation of microcracks due to increased crack roughness and crack closure [37]. At a displacement of 1.309 mm (Fig. 8b), the severe shear motion causes the boundaries of the α plates to offer little resistance to deformation. A large number of microcracks appear across one or several α plates. Some of the fractured α plates exhibit stepped staircase-like features. This behavior is analogous to the cleavage fracture usually observed for materials with a higher grain-boundary fraction. The fracture surface in Fig. 8c also shows trans-crystalline fracture characteristics.
Fig. 8 Microstructures of the X-direction specimen during in-situ tensile loading: (a) microcrack initiation at a displacement of 1.011 mm, (b) α-plate shear fracturing at a displacement of 1.309 mm, and (c) transcrytalline fracture. In the Z-direction specimen, because of α colonies and their crystal orientation, the fine basket-woven structure (see Fig. 4c) with smaller α colonies decreases the slip length. Therefore, more uniform deformation is observed at a displacement of 1.067 mm (Fig. 9a). The deformation during tensile testing leads to substantial changes in the microstructure. These α plates are elongated and bent, indicating high plasticity. As the displacement increases to 1.279 mm, a crack is observed at the α-plate boundaries because of uniaxial tension (Fig. 9b). The sharp angles of the microcrack cause a local stress concentration and play substantial roles in the initial fracture by promoting crack initiation under dynamic tensile loading, particularly for Z-direction tensile loading [29]. The fracture in Fig. 9c also shows the shearing of unfavorably oriented α plates.
Fig. 9 Microstructures of the Z-direction specimen during in-situ tensile loading: (a) uniform deformation at a displacement of 1.067 mm, (b) microcracking along the α plates at a displacement of 1.279 mm, and (c) transcrystalline fracture.
5. Conclusions The anisotropic mechanical properties of LDMD AM Ti6Al4V components were compared for transverse and longitudinal orientations. The fracture mechanisms were clarified on the basis of microstructural distributions and fracture processes, as tested by in-situ SEM tensile loading. The conclusions are summarized as follows: 1) LDMD AM produced compact and defect-free Ti6Al4V components. The anisotropic mechanical properties were characterized with high strength and low elongation in the transverse direction and high elongation and reduction of area in the perpendicular direction. The anisotropic mechanical properties are caused by the anisotropic microstructural distribution. 2) The microstructure in the transverse tensile specimen was composed of α phases having soft grain boundary, and short and thin columnar prior-β grains that grew approximately perpendicular to the tensile axis, rendering the tensile strength and fracture strength significantly higher than in the longitudinal direction. However, the longitudinally oriented tensile specimen consists of long columnar prior-β grains parallel to the tensile axis and small α colonies with short α plates, which
result in the preferential accumulation of damage. The results of the study indicate that the longitudinal specimen shows a low stress and high elongation. 3) The anisotropic tensile properties were also examined by in-situ tensile testing. The transverse specimen fractured along the prior-β grain boundary with plenty of shear bands. The vertical β-grain boundaries and large α colonies contribute to the high strength by resisting microcrack propagation. The longitudinal specimen fractured with symmetric necking due to the fine α grains, which results in uniform deformation of the grains on both sides of the grain boundaries.
Acknowledgements The authors wish to thank the National Natural Science Foundation of China (Grant number: 51505010) and Beijing Natural Science Foundation (Grant number: 2152007)
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