key: cord-1014706-9alkiafn
authors: Tseng, Jo-Chi; Huang, Wei-Chin; Chang, Wei; Jeromin, Arno; Keller, Thomas F.; Shen, Jun; Chuang, Andrew Chihpin; Wang, Chun-Chieh; Lin, Bi-Hsuan; Amalia, Lia; Tsou, Nien-Ti; Shih, Shao-Ju; Huang, E-Wen
title: Deformations of Ti-6Al-4 V additive-manufacturing-induced isotropic and anisotropic columnar structures: In situ measurements and underlying mechanisms
date: 2020-06-12
journal: Addit Manuf
DOI: 10.1016/j.addma.2020.101322
sha: b0a03e2b3750e14e9f64e009825f6d7e5e7f55ec
doc_id: 1014706
cord_uid: 9alkiafn

The deformations of isotropic and anisotropic Ti-6Al-4 V columnar structures fabricated by additive manufacturing were extensively examined. The distinct texture and microstructure distributions were characterised. In situ X-ray diffraction measurements show different lattice activities resulting from the different microstructure distributions. Spatially resolved mapping revealed manufacturing-induced crystallite-orientation distributions that determine the deformation mechanisms. We propose a self-consistent model to correlate the multi-scale characteristics, from the anisotropic-texture-distribution microstructure to the bulk mechanical properties. We determined that basal and pyramidal slip activities were activated by tension deformation. The underlying additive-manufacturing-induced crystal plasticity plays a major role. We find that the texture development of the columnar structures and the distribution of crystallite orientation achieved by different processing conditions during additive manufacturing have important effects on the mechanical properties. The dominant deformation mode for the anisotropic Ti-6Al-4 V columnar structure is basal slip, and that for the isotropic Ti-6Al-4 V columnar structure is pyramidal slip. The difference may be important for determining the fatigue behaviour.

The use of additive manufacturing (AM) in various applications has increased significantly in the past few years [1] . This emerging technology has even led to new business models for advanced manufacturing [2, 3] . For example, the recent COVID-19 situation may reshuffle the global supply chain, and AM could provide new solutions [4] . The unique feature of AM is the capability to integrate digitalised data from various sources, from high-throughput examinations to advanced photon sources for integrated computational materials engineering (ICME) [5] .

The advantages of AM for metallic and other systems can be found in the review by Tofail et al. [6] . Specifically, Ti-based alloys, in particular Ti-6Al-4V, have been developed for several decades for use in aerospace, automotive, and medical applications. Given its nontoxicity, anti-corrosive nature, and stable properties in the human body, Ti-6Al-4V is an J o u r n a l P r e -p r o o f important material for biomedical implants [7] . Furthermore, its mechanical properties are also well adapted to the growth of bones [8] . The high strength-to-weight ratio and fatigue resistance of Ti-6Al-4V also make it attractive for use in the components of turbine engines, gears, and wings in the aerospace industry. These properties make Ti-6Al-4V an in-demand material, with more than 50% of the metallurgy industry working towards its production [9] . However, certain bottlenecks have been found in the conventional approach of manufacturing Ti-6Al-4V, such as the processing of samples with complicated geometry, oxidation of the surface, and flexibility of sample sizes [10] . To solve these problems, efforts have been made towards developing coated and/or doped Ti-6Al-4V or using novel processing approaches, such as AM for titanium alloys [11] [12] [13] . Hence, additively manufactured Ti-6Al-4V has drawn significant attention [14, 15] .

Laser powder-bed fusion (LPBF), electron-beam melting, and directed energy deposition are three major AM technologies. The spot-sized beam and quick heating-sintering methodologies make AM an elegant approach for fabricating samples rapidly, manufacturing neck-like shapes, and adding dopants easily to Ti-6Al-4V.

However, it has been found that the temperature gradient, build direction, and other AM processing parameters all interact with each other to affect the microstructure of the printed component, changing its mechanical properties [16] [17] [18] [19] [20] . The grains, coexistence of α and β phases, and crystallite defects are reported to be crucial factors affecting the mechanical properties of Ti-6Al-4V [21] [22] [23] [24] . Moreover, porosity is a vital concern for Ti-6Al-4V [25] .

Characterisation methods such as electron microscopy (EM), electron backscatter diffraction (EBSD), and X-ray diffraction (XRD) have been used to identify the homogeneity of phases and analyse the microstructural properties of samples [26] [27] [28] . Typically, the microstructural properties are determined by surface analysis (using EM/EBSD) and phase analysis (using XRD) [21, 29] . Moreover, by combining these methods, comprehensive details of grain boundaries can be revealed. Raabe et al. demonstrated that 3D EBSD characterisation can be used to correlate the microstructure of alloys with their bulk properties [30] [31] [32] . However, EBSD is a destructive analysis approach, and the quality of the data depends on the sampling methodology. Specifically, collecting data from a well-polished specimen surface is essential for a confident reconstruction process from the Kikuchi pattern with all the scattering events included.

Rai et al. noted the importance of the evolution of the grain structure during powder-bed AM via simulation [33] . Later, Liu et al. provided insight into the mechanisms of the columnarto-equiaxed grain transition during metallic AM [34] . Following their work, in this study, we J o u r n a l P r e -p r o o f attempt to reveal the microstructure evolution resulting from deformation, not fabrication, using XRD and tomography.

The use of an advanced photon source, such as a synchrotron X-ray or neutron beam with high penetration capability, combined with in situ tensile experiments, can reveal real-time lattice-to-macroscopic mechanical properties. Accordingly, the correlation of lattice strain with engineering strain can be obtained simultaneously. Studies by Wilkinson et al. have found a strong correlation between lattice strain and engineering strain with texture development during tensile experiments using high-energy synchrotron XRD analysis [35] [36] [37] . Huang et al.'s in situ neutron diffraction experiments also successfully captured the phase transformation of the AMinduced transient phase [17, 38] . However, the pores induced by the AM process play a crucial role in mechanical performance, as determined by in situ diffraction-related results reported by Qiu et al. [39, 40] . These results are important for the potential high-throughput experimental methodologies of the materials genome initiative (MGI) [41] and efficient AM development and production [6] . Therefore, information obtained from X-ray experiments can contribute to a comprehensive correlation between microstructure analysis and the bulk properties of materials.

The fabrication strategies [42, 43] associated with AM have been extensively summarised by the AM community for controlling the elongation, strength [44] [45] [46] [47] , residual stress [48] [49] [50] [51] , and fatigue [52, 53] properties. For example, Galarraga et al. already categorized the microstructure-dependent mechanical properties of the Ti-6Al-4V ELI alloy fabricated by electron beam melting (EBM) [54] . Galarraga et al. have found that equiaxed α+β, columnar α+ β, partially α′, and fully α′correspond to furnace-cooling, EBM-built, air-cooling, and watercooling condition, respectively, show different strengths and elongation levels accordingly [54] .

For large deformations, Galarraga et al. summarized the thickness of hcp lath and lamellar structure effects [54] . Wang et al. demonstrated the hierarchically heterogeneous microstructure [55] .

However, unlike the aforementioned comprehensive discussion on the overall strengths and elongation comparisons in the archived literature, we focus on small deformations to investigate the onset of plasticity in this work, which is essential for examination of the fatigue mechanisms [53] . For correlation with the bulk mechanical behaviour, we followed the protocols of in situ diffraction measurements because detectable lattice asymmetry can reveal the initiation of fatigue cracks [56] . Once there is a sufficient amount of fatigue pore formation, Considère's J o u r n a l P r e -p r o o f volume conservation for crystal plasticity will not be followed, which can be observed by the in situ lattice deformation [57, 58] .

In the current study, we applied various high-energy synchrotron X-ray diffraction (SXRD) experimental approaches to investigate LPBF-manufactured Ti-6Al-4V and the effect of the LPBF processing parameters, such as the build direction. Liu et al. [34] and Rai et al. [33] have already concluded that, together with the thermal gradient, the solidification rate can generate different microstructures, such as columnar and equiaxed grains. However, the direction correlation of such a columnar structure with the bulk mechanical properties of a material remains uncertain. In this work, we examined the deformation mechanisms for two additivemanufacturing-induced columnar structures: one has an isotropic texture, and the other has an anisotropic texture.

Moreover, we designed spatially resolved scans to map crystallite orientation distributions. This non-destructive mapping can disclose the local texture and microstructure, while the in situ diffraction analysis can provide kinetic information at the lattice scale and characterise the bulk properties of the material. Meanwhile, the average crystal structure of the prepared Ti-6Al-4V sample was also theoretically simulated (and compared with experimental results) for a quantitative analysis of crystal defects via whole powder pattern modelling (WPPM) [59, 60] .

In addition, we can also use the synchrotron X-ray to scan the specimen with varying penetration power to gauge information through the thickness of the samples. Subsequently, spatial scans allowed us to map the microstructural distribution for the whole testing area and to quantify the strain by modelling of the X-ray patterns. This makes it possible to build multi-scale character recognition from the lattice microstructure to bulk mechanical properties for LPBF specimens. This work focuses on examining the AM-induced texture and microstructure distribution effects of Ti-6Al-4V, which are important for fatigue characterisation.

Two Ti-6Al-4V samples were prepared via LPBF with build directions parallel (vt-anisoTAV) and perpendicular (hz-isoTAV) to the axis to form distinct textures for uniaxial tensile experiments (Figure 1(a) and (b) ). The zigzag approach was applied, and the vectors were changed by 67° from layer to layer [61] . Processing was done under a flowing, inert argon atmosphere with oxygen monitoring. All processing was completed at room temperature with no heat applied to the build plate. Samples were removed from the machine and cleaned of extra powder by sonicating in water. Parts were then dried with clean compressed dry air.

Through this process, multilayer-like Ti-6Al-4V samples can be obtained. Specifically, we controlled the fabrication parameters to ensure our samples were mainly in a hexagonal close packed (hcp) as α phase with very minor body centered cubic (bcc) phase as β phase. For overall bulk strengths and elongations, our samples have the same trend as Galarraga et al.

summarized [54] . horizontal (hz-isoTAV) build directions via the LPBF approach.

To examine the chemical distribution, the specimens were illuminated using the X-ray Nanoprobe beamline of the Taiwan Photon Source, TPS 23A, with 100 nm spatial resolution for complementary elemental distributions. The energy of the incident X-rays was 12.8 keV, well above the k-edge absorption energies of the constituent elements of the specimens. The excited fluorescence was collected simultaneously by a silicon drift detector with an energy resolution of approximately 150 eV [62] .

To reveal the pores, synchrotron transmission X-ray microscopy was used to examine beneath the material surface. Transmission X-ray microscopy (TXM) enables the characterisation of high-resolution X-ray radiography in three dimensions (3D). With the aid of a Fresnel zone plate and high-flux synchrotron hard X-ray source, TXM can achieve a spatial resolution of approximately 60 nm. The images of the tomography were reconstructed using Amira software [63] . The number and the geometry of the voids were estimated from the Amira algorithm analysis of the 3D TXM images.

To investigate the deformation mechanisms, in situ uniaxial tensile SXRD experiments were carried out at P02.1, DESY [64] . The samples were mechanically cut to the required shape J o u r n a l P r e -p r o o f according to the ASTM subsize of the E8M-04 Standard Test Methods for Tension Testing of Metallic Materials for the in situ uniaxial tensile SXRD experiments [65] . A high-energy X-ray beam of approximately 60 keV (wavelength ~0.207 Å) and a Perkin-Elmer EN1621 2D detector with a 0.2 × 0.2 μm 2 pixel size were used to collect the 2D XRD images. The sample-to-detector distance was set to 1,185 mm. A 1 × 1 mm 2 X-ray beam size was used for all experiments. For the tensile experiments, a 5 kN tensile/compression module machine (Kammrath-Weiss) was used, with an elongation rate of 20 μm/s applied in continuous mode. The continuous mode enables the correlation of engineering mechanical properties and lattice kinetics without creep.

During the uniaxial tensile experiments, both engineering strain and XRD patterns were simultaneously collected every second to study their correlation. The samples were also measured before and after the uniaxial tensile experiment with a 1 × 1 mm 2 X-ray beam spot over the entire rectangular area of the specimen to examine the local texture development.

Prior to the analysis of the 2D XRD images, Cr2O3 powder was measured to calibrate the sample-to-detector distance, the beam centre, and the tile of the detector. Based on the instrumental parameters obtained by the Cr2O3 measurement, the collected 2D XRD images for samples could be integrated into 1D XRD patterns with the Data Analysis Workbench (DAWN) [66, 67] . Two strategies for the integration of 2D XRD images were used. (1) Each 2D XRD image was integrated for 36 distinct orientation regions along the azimuthal angle to the 1D XRD patterns to study the texture properties of samples. Each peak in the 1D XRD patterns was fitted with pseudo-Voigt functions by a Python program, and the peak intensity, peak position, and peak width were obtained for subsequent analysis. (2) The 2D XRD image was integrated using full Debye-Scherrer rings to obtain the average peak positions and lattice parameters. The determination of d0 for each sample was performed to address the starting parameters for the asprepared samples with residual stresses induced from the AM process. Subsequently, calculation of the change in lattice strain during the tensile experiment was used to determine the deformation behaviour of samples. In addition to comparing the XRD patterns obtained from the two sampling techniques, the theoretical static strain and growth fault were simulated via the whole powder pattern modelling (WPPM) approach [59, 68, 69] . The broadening of peaks in an XRD pattern can be regarded as an effect of the microstructure characteristics, such as crystallite size and strain. One of the modelling approaches, WPPM, allowed us to simulate the amount of growth faults for hcp Ti-6Al-4V. WPPM is implemented in the PM2K v2.1 software package, and an XRD pattern can be generated by the Fourier transform method using a defined crystal model with the desired defects, such as dislocations, twins, or growth faults. By changing the probability of the mentioned defects, a series of XRD patterns can be obtained to understand the number of defects in correlation with the change in the XRD pattern. Because the size of crystallites in the studied case was not the dominant factor, the parameters controlling crystallite size were fixed during the simulation process.

A high-resolution field-emission scanning electron microscope (FEI Nova Nano SEM 450) was used for the EBSD experiments with a 5 keV electron beam [70] . The samples were fixed on copper tape.

Prior to the uniaxial tensile experiments, the samples were mechanically polished and characterised to investigate the physical properties related to mechanics, such as the concentration of pores, the distribution of elements, and the homogeneity of phases. We followed the protocols reported by Stef et al. [71] to examine the pores in our samples. The 3D transmission X-ray tomography (TXM) high-spatial-resolution images are shown in SI 1(a) and (b), respectively. Although the pore volume fractions of our samples are much lower, Stef et al. have confirmed the influence of voids on the mechanical properties of ultimate tensile stress (UTS) and ultimate elongation for a volume fraction as low as 0.48 vol.%. Hence, in this study, we limited our scope to only small deformations within 3.8%, which is much lower than the UTS and the ultimate elongation. For small deformations, Considère's construction shows that in deformed crystalline metals, volume conservation occurs during plastic deformation subjected to slip and dislocation activities without void nucleation [72] . Specifically, the onset of void effects can be revealed by the deviation of the Poisson's ratio [56, 58] , which describes the expansion of a material in directions perpendicular to the direction of compression. We applied angle-resolved diffraction to monitor the two orthogonal diffractions to examine the evolution of the diffraction pattern in both the tension and transverse directions simultaneously. The residual strains present in the as-prepared samples after removal from the LPBF bed were analysed via scanning XRD, as shown in SI 3. The range of the residual strain was similar for both as-prepared samples, between approximately 10e -2 and 10e -3 %.

Although only the hcp phase was identified for both samples, the peak intensities in both XRD patterns are slightly different, in particular for the (0002) Therefore, to resolve the convoluted information in the XRD patterns from both samples, the raw 2D XRD images of as-prepared vt-anisoTAV and hz-isoTAV were analysed (Figure 2 (a) and (c)). The azimuthally resolved Debye-Scherrer rings in Figure 2 (a) show that (0002), the second ring away from the beam centre, is partially discontinuous for vt-anisoTAV. The intensities of the (0002) Debye-Scherrer ring are much higher at the azimuthal angles of 50°, 140°, 180°, 230°, and 320°. In contrast, no regularity of the intensity changes was found for the (0002) ring of iso-TAV (Figure 2(b) ). Considering that the thermal history of the build layers in AM is the same for both samples, the stacking structure along with the X-ray viewing direction and tensile direction can be distinct if the build direction is different (Figure 2(b) and (d) ). The yellow cubes represent the X-ray transmitted volume with a 1 × 1 mm 2 beam size, which corresponds to the measuring area for each resulting XRD pattern. It is clear that vt-anisoTAV, with the build direction perpendicular to the tensile direction, has a layer-like structure with highly preferred orientated crystallites (Figure 2(b) ). When X-rays are transmitted through such a sample, the regular changes in diffraction intensities can be illustrated. In contrast, the crystallite structure of the hz-isoTAV is randomly stacked along the X-ray path if its build direction is parallel to the tensile J o u r n a l P r e -p r o o f direction (Figure 2(d) ). Consequently, the plot of the (0002) Debye-Scherrer ring appears to have no intensive regularity with the change in intensity along the azimuthal angles. To exclude the randomness from measuring with a 1 × 1 mm 2 X-ray beam, spatially resolved scans over the entire rectangular area of both samples were performed (Figure 3 As shown from the static plots in Figure 3 (a) and 3(b), six strong peaks were obtained for vt-anisoTAV, at azimuthal angles of 40°, 130°, 170°, 220°, 310°, and ~0°, whilst no strong peak was found for hz-isoTAV. It is expected that the peaks can be paired as 40°/220°, 170°/~0°, and 130°/310° from the perspective of the multiplicity of {0002}. Interestingly, in the non-normalised data set (SI 4), the same trend of strong peaks was found in the third Debye-Scherrer ring, assigned to (101 ̅ 1). The crystallographic angle between (0002) and (101 ̅ 1) is nearly 120° in hcp Ti-6-4V. Using tan -1 (c/a * 2/√3) for calculation [74] , the same scattering angle along φ could occur when the crystallites are oriented such that the angle between the incident beam and these two crystal planes is approximately 60°. The ratio of the X and Y components of the rotated (hkl) planes satisfies the following equation [36] : 

where rhkl is the vector of the (hkl) plane normal to its crystal structure configuration, X and Y are in-plane components of the crystallite orientations, and φ is the angle between X and Y (Figures 2   and 3 ). The two conditions discussed above are only satisfied when (101 ̅ 1) is almost parallel to the build direction and (0002) is at an angle of approximately 30° to the sample surface ( Figure   4 (a)).

Accordingly, the spatial mapping of the distribution of (0002) preferred orientations over the entire sample with respect to transmitted X-ray cross-sections is plotted in Figure 4 (b) and

(c). In the vt-anisoTAV, Figure 3 (b), more than 80% of the area was in the same orientation.

Conversely, the hz-isoTAV sample had less than 35% of crystallites in the same orientation.

Thermodynamically, such preferred orientations of (0002) in the Ti-6Al-4V system could be caused by the texture formation of an hcp structure. When the crystallites grow rapidly, the growth of an hcp structure along the c axis is likely preferable, leading to (101 ̅ 1) facets lying on the substrate [47, [75] [76] [77] [78] [79] . The map of the microstructure from columnar to equiaxed grains subjected to both various thermal gradients and solidification rates has been reported by Liu et al. [34] . With different building strategies, the crystallites of our samples have different effective times and regions for grain growth during the manufacturing process. Furthermore, successive build layers can be regarded as substrates for the following layers in the AM approach.

Accordingly, our samples manufactured via LPBF show different partial preferred orientations for the crystallites. For more details of how and why the hcp texture is formed, please see the following representative references [80] [81] [82] .

The nature of the growth direction of an hcp structure leads to the assumed stacking structure for the two samples with respect to the LPBF build strategy, as shown in Figure 2 (b) and 2(d). The crystallites were rotated 67° layer-by-layer along the build direction (for ease of understanding, the crystallites are shown with a 3D structure). The 3D crystal structure with face (0002) facets and side (101 ̅ 1) facets is shown based on an hcp structure configuration.

Therefore, the crystallites were arranged in the same way for both samples. However, depending on the viewing direction of the X-ray beam, the scattering events differed according to the orientation of crystallites in the area measured by the 1 × 1 mm transmitted X-ray beam. For the anisoTAV sample, the (0002) was facing or at an angle of less than 30° to the surface of the specimens (Figure 2(b) ). Consequently, a highly preferred orientation of (0002) could be expected and observed in the X-ray pattern; the (0002) was parallel to the tensile direction in this configuration. In contrast, for the iso-TAV sample, a highly random crystallite orientation was evident in the X-ray transmitted area. However, crystallites with a growth direction along (101 ̅ 1)

were mostly observed on the surfaces of the samples, which was parallel to the tensile direction, as shown in Figure 2 (d). Based on the two presumed directional-distinct stacking models, the Xray scattering events of d(0002) were expected to be azimuthally dependent in both cases [36] .

By calculating the diffraction angles of the aforementioned crystallite orientations, the strong peaks, from three grouped φ angles, obtained from experimental data (Figure 3(a) ) were matched to the proposed stacking crystallite model shown in Figure 2 (b). It should be noted that for the case assuming (0002) to be almost perpendicular to the X-ray beam, the scattering events on (0002) are difficult to detect because of the use of a planar detector. Conversely, no preferred orientation was found from the 2D XRD images for hz-isoTAV, as shown in Figure 3 Sometimes, the preferred orientation of crystallites is not significantly related to the texture formation of crystallites, and vice versa. In this particular case, strong preferred orientations of crystallites were observed, based on the scattering spots shown in the 2D XRD images.

The stacking structure of the crystallites for vt-anisoTAV could be considered a layer-like structure with respect to the direction of the uniaxial tensile experiments, and the (0002) directions were oriented at an angle of less than 30°. In contrast, for the hz-isoTAV sample, no specific (0002) stacking structure was observed; however, the (101 ̅ 1) is in-plane with the applied tension direction.

J o u r n a l P r e -p r o o f Interestingly, the presumed stacking models also agree with the results of EBSD tests ( Figure 5 ). The crystallites in the vt-anisoTAV sample have a layer-like arrangement along the build direction, as shown in Figure 5 (a). Looking at the X-ray viewing direction in Figure 4 (c), the crystallites are regularly oriented, which leads to the discontinuity in the Debye-Scherrer rings in the 2D XRD image shown in Figure 2 (a). Such texture properties are specific to the tensile direction for vt-anisoTAV. Alternatively, on the surface of the iso-TAV, an in-plane stacking structure along (101 ̅ 1) was observed for the majority of the sample test area ( Figure   5 (b)). Also, the crystallite orientation along the X-ray viewing direction was highly random, resulting in the good continuity of the Debye-Scherrer rings in the 2D XRD image shown in 

The correlation between engineering mechanical properties and the proposed stacking structure models is discussed below focusing on small deformation. Both samples were examined using in situ XRD uniaxial experiments under continuous mode at the high-energy Synchrotron X-ray beamline, P02.1, DESY. This allowed us to collect the 2D XRD images rapidly (1 frame/s) and illustrate the dynamic changes in engineering strain and lattice strain simultaneously, without creep. By integrating the 2D XRD images to the 1D XRD patterns and subsequently fitting the diffraction peaks, the lattice strain on different crystal planes can be calculated using:

where d is the d-spacing between two lattice planes. The XRD patterns were collected for both samples before the tensile experiment to define the d(0).

The results in Figure 6 show that the lattice strain averaged over all of the azimuthal angles and engineering strain are highly correlated in both the elastic and plastic deformation regions during the uniaxial tensile experiments [37] . However, the individual lattice strain evolution on different crystal planes was very different between the two cases. The lattice strains calculated from (1010), (0002), and (101 ̅ 1) are similar for vt-anisoTAV, although the lattice strain of the (0002) crystallite plane slightly decreases in the plastic deformation region ( Figure   6 (a)). In contrast, the lattice strain of (101 ̅ 1) is much more pronounced in both the elastic and plastic deformation regions for the hz-isoTAV sample (Figure 6(b) ). Such distinctive directionally related deformation can be correlated with the proposed stacking structure. For hz-isoTAV, (101 ̅ 1) was presumably in-plane with the uniaxial tensile experiment, which played a crucial role during the deformation, as observed from the highest strain accumulated [22] . In contrast, the lattice strain seems to have developed nearly equally for the vt- For the vt-anisoTAV, the more layer-like stacking structure caused by its columnar texture led to presumably distinct layers perpendicular to the experiment's tensile axis. Most of the (0002) facets were at an angle of less than 30° to the tensile direction. Such a stacking structure resulted in a strong correlation between the (0002) and (101 ̅ 1) planes and the major deformation of (0002) along the tensile direction. In addition, between distinct layers or within one build layer, a twist-like distortion of the crystallites occurred (Figure 7(d) ). Not until the lattice strain of ε22 became zero and the sliding of the crystallites occurred did the rupture of the sample, either in the pore-joint area or in the crystallite-twisted area, take place. It has been reported that larger amounts of plastic strain in the basal planes, or the (0002) facets, can lead to less critical stress normal to the (0002) facet and result in a quicker occurrence of rupture [22, 83] . This could explain the weaker mechanical properties in the vt-anisoTAV sample. Based on the proposed model depicted in Figure 3 (a) and (b) and the discussion above, the slip systems between two samples were found to be distinct because of the different microstructures resulting from the different orientations of the columnar texture. Basal and pyramidal slip are the dominant systems for the vt-aniso-and hz-isoTAV samples, respectively.

In addition, the slip system evolution could lead to the different mechanical behaviour. In summary, the stacking structures caused by the isotropic and anisotropic columnar structures subjected to different build strategies led to distinct underlying deformation mechanisms.

Ti-6Al-4V was manufactured using LPBF, and the characteristics at different scales, including the microstructure, texture, and engineering mechanical properties, were investigated.

Synchrotron 2D XRD analysis suggests that the orientations of the stacking crystallites differed with respect to the local columnar texture distributions, which resulted from control of the thermal gradient and solidification rate, as a result of the nature of the preferred growth direction of the hcp structure. For the vertically built Ti-6Al-4V (vt-anisoTAV), the crystallites were highly preferentially orientated in an anisotropic manner, where, in most of the grains, the (0002) planes were distributed with orientations less than 30° from the direction of the tensile experiments. In contrast, the crystallites in the horizontally built Ti-6Al-4V (hz-isoTAV) were randomly oriented. Similar evidence was obtained from the complementary EBSD analysis. The 

The rise of 3-D printing: The advantages of additive manufacturing over traditional manufacturing

How Additive Manufacturing Technology Changes Business Models? -Review of Literature

Additive manufacturing in construction: A review on processes, applications, and digital planning methods

Supply chain implications of additive manufacturing: a holistic synopsis through a collection of case studies

Overview of Materials Qualification Needs for Metal Additive Manufacturing

Additive manufacturing: scientific and technological challenges, market uptake and opportunities

Microstructure and mechanical behavior of Ti-6Al-4V produced by rapid-layer manufacturing, for biomedical applications

New Developments of Ti-Based Alloys for Biomedical Applications

Titanium: A Technical Guide

Energy and emissions saving potential of additive manufacturing: the case of lightweight aircraft components

Problems on the development of hard and low friction in-situ coatings on Ti-6Al-4V using laser cladding

Additive manufacturing of Ti6Al4V alloy: A review

Effect of Ti additions on mechanical properties of NbMoTaW and VNbMoTaW refractory high entropy alloys

Solidification microstructure simulation of Ti-6Al-4V in metal additive manufacturing: A review

As-fabricated surface morphologies of Ti-6Al-4 V samples fabricated by different laser processing parameters in Selective Laser Melting

Controlling phase composition in selective laser melted stainless steels

Hardening steels by the generation of transient phase using additive manufacturing

Additive manufacturing of metals

Transient Phase-Induced Nucleation

Imaging of Transient Structures Using Nanosecond in Situ TEM

Strong grain neighbour effects in polycrystals

Evolution of lattice strain in Ti-6Al-4V during tensile loading at room temperature

Experimental and theoretical studies of intergranular strain in an alpha titanium alloy during plastic deformation

Effect of texture on load partitioning in Ti-6Al-4V

Investigation of Porosity Changes in Cast Ti6Al4V Rods After Hot Isostatic Pressing

A study of the microstructural evolution during selective laser melting of Ti-6Al-4V

Anisotropic tensile behavior of Ti-6Al-4V components fabricated with directed energy deposition additive manufacturing

Effect of build geometry on the β-grain structure and texture in additive manufacture of Ti6Al4V by selective electron beam melting

Strains, planes, and EBSD in materials science

Investigation of orientation gradients around a hard Laves particle in a warm-rolled Fe3Al-based alloy using a 3D EBSD-FIB technique

Three-dimensional investigation of the texture and microstructure below a nanoindent in a Cu single crystal using 3D EBSD and crystal plasticity finite element simulations

Orientation gradients and geometrically necessary dislocations in ultrafine grained dual-phase steels studied by 2D and 3D EBSD

Simulation of grain structure evolution during powder bed based additive manufacturing

Insight into the mechanisms of columnar to equiaxed grain transition during metallic additive manufacturing

A synchrotron X-ray diffraction study of in situ biaxial deformation

Assessment of X-ray diffraction and crystal plasticity lattice strain evolutions under biaxial loading

Shih-Min Chen, Hung-Sheng Choug, Hardening steels by the generation of transient phase using additive manufacturing

Plastic anisotropy and deformation-induced phase transformation of additive manufactured stainless steel

Microstructure and tensile properties of selectively laser-melted and of HIPed laser-melted Ti-6Al-4V

On the role of melt flow into the surface structure and porosity development during selective laser melting

Fulfilling the promise of the materials genome initiative with high-throughput experimental methodologies

Correlation between process parameters, microstructure and properties of 316 L stainless steel processed by selective laser melting

Process-property-geometry correlations for additivelymanufactured Ti-6Al-4V sheets

Additive manufacturing of metallic components -Process, structure and properties

Refinement of the grain structure of additive manufactured titanium alloys via epitaxial recrystallization enabled by rapid heat treatment

Anisotropic polycrystal plasticity due to microstructural heterogeneity: A multi-scale experimental and numerical study on additively manufactured metallic materials

Additive manufacturing of Ti6Al4V alloy: A review

Thermomechanical Modeling of Additive Manufacturing Large Parts

Mitigation of thermal distortion during additive manufacturing

Linking process, structure, property, and performance for metal-based additive manufacturing: computational approaches with experimental support

An improved prediction of residual stresses and distortion in additive manufacturing

Study on the fatigue behaviour of selective laser melted AlSi10Mg alloy

Fatigue life of additively manufactured Ti-6Al-4V in the very high cycle fatigue regime

Effects of heat treatments on microstructure and properties of Ti-6Al-4V ELI alloy fabricated by electron beam melting (EBM)

Additively manufactured hierarchical stainless steels with high strength and ductility

Cyclic-Loading Induced Lattice-Strain Asymmetry in Loading and Transverse Directions

Fatigue-induced reversible/irreversible structural-transformations in a Ni-based superalloy

Fatigue induced deformation and thermodynamics evolution in a nano particle strengthened nickel base superalloy

PM2K: a flexible program implementing Whole Powder Pattern Modelling

Whole powder pattern modelling

Method and device for manufacturing a three-dimensional object United States

Element Effects on High-Entropy Alloy Vacancy and Heterogeneous Lattice Distortion Subjected to Quasi-equilibrium Heating

Microstructure Investigation for a PM2.5 Air-Filter Efficiency Study of Non-Woven Polypropylene

Beamline P02.1 at PETRA III for high-resolution and high-energy powder diffraction

Materials, E8M-04 Standard Test Methods for Tension Testing of Metallic Materials (Metric) 1, ASTM international2004

Processing two-dimensional X-ray diffraction and small-angle scattering data in DAWN 2

Data Analysis WorkbeNch (DAWN)

Dislocation effects in powder diffraction

On the powder diffraction pattern of crystals with stacking faults

Journal of largescale research facilities

Mechanism of porosity formation and influence on mechanical properties in selective laser melting of Ti-6Al-4V parts

A model of ductile fracture based on the nucleation and growth of voids

Significance of mechanical twinning in hexagonal metals at high pressure

Deviatoric deformation kinetics in high entropy alloy under hydrostatic compression

An Automated Visual Inspection System for the Classification of the Phases of Ti-6Al-4V Titanium Alloy

Phase transformations during cooling in α+β titanium alloys

Effect of energy density on the microstructure and texture evolution of Ti-6Al-4V manufactured by laser powder bed fusion

Phase-field model study of the crystal morphological evolution of hcp metals

Modified embedded-atom method interatomic potentials for Ti and Zr

hcp-bcc structural phase transformation of titanium: Analytic model calculations

Classical potential describes martensitic phase transformations between the α, β, and ω titanium phases

Identifying the multiplicity of crystallographically equivalent variants generated by iterative phase transformations in Ti

Failure of Metals I -Brittle and Ductile Fracture

Supporting Information SI 2. XRD analysis with whole ring integration from 2D-XRD images for as-prepared vt-anisoTAV (black line) and hz-isoTAV

Strain mapping of as-prepared (a) vt-anisoTAV and (b) hz-isoTAV samples calculated by scanning X-ray diffraction method and the static information of residual strain over the whole

Theoretical calculation of XRD patterns using whole powder pattern modelling (WPPM) for an hcp phase with different probabilities of {101 ̅ 1} growth faults. The intensity changes in (101 ̅ 1) during the uniaxial tensile experiment are plotted for (b) vt-anisoTAV and (c) hz-isoTAV

We are deeply grateful to Dr