Formability, microstructure and mechanical properties of Ti-6Al-4V deposited by wire and arc additive manufacturing with different deposition paths

Journal Pre-proof Formability, microstructure and mechanical properties of Ti-6Al-4V deposited by wire and arc additive manufacturing with different deposition paths Yefei Zhou, Guangkuo Qin, Lei Li, Xin Lu, Ran Jing, Xiaolei Xing, Qingxiang Yang PII:

S0921-5093(19)31440-6

DOI:

https://doi.org/10.1016/j.msea.2019.138654

Reference:

MSA 138654

To appear in:

Materials Science & Engineering A

Received Date: 31 July 2019 Revised Date:

6 November 2019

Accepted Date: 7 November 2019

Please cite this article as: Y. Zhou, G. Qin, L. Li, X. Lu, R. Jing, X. Xing, Q. Yang, Formability, microstructure and mechanical properties of Ti-6Al-4V deposited by wire and arc additive manufacturing with different deposition paths, Materials Science & Engineering A (2019), doi: https://doi.org/10.1016/ j.msea.2019.138654. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. 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. © 2019 Published by Elsevier B.V.

Formability, microstructure and mechanical properties of Ti-6Al-4V deposited by wire and arc additive manufacturing with different deposition paths Yefei Zhou

a, b

, Guangkuo Qin a, Lei Li c, Xin Lu d, Ran Jing e, Xiaolei Xing a, b*,

Qingxiang Yang b* a.

College of Mechanical Engineering, Yanshan University, Qinhuangdao 066004, P. R. China

b.

State Key Laboratory of Metastable Materials Science & Technology, Yanshan University,

Qinhuangdao 066004, P. R. China c.

Northwest Institute for Nonferrous Metal Research, Xi’an 710016, P. R. China

d.

Faculty of Materials Science and Engineering, Kunming University of Science and Technology,

Kunming, 650093, P.R. China. e.

National & Local Joint Engineering Laboratory for Slag Comprehensive Utilization and

Environmental Technology, Shaanxi University of Technology, Hanzhong 723000, P. R. China *.

Corresponding author, Tel: 86-335-8387471, Fax: 86-335-807-4545,

E-mail: [email protected] (X. L. Xing); [email protected] (Q. X. Yang)

Abstract By means of wire and arc additive manufacturing (WAAM) in an inert environment, the Ti-6Al-4V thin-wall parts were fabricated. To explore the effect of different deposition paths of WAAM on formability, microstructure and mechanical properties, the samples were analyzed by surface profiler, optical microscope (OM), scanning electron microscopy (SEM), X-ray diffraction (XRD), Vickers hardness and universal testing machine. The results indicate that the side surface waviness (SW) of the samples from samples by reciprocating deposition path in both ends are higher than those in the middle, but SW of the samples by co-directional

deposition path gradually rise from arc starting to arc stopping. Prior-β grains from the samples by co-directional deposition path are inclined relative to the building direction, and there are many winding prior-β grains in the samples by reciprocating deposition path. In the vertical direction, the strength of samples by co-directional deposition path is higher than those by reciprocating deposition path. Furthermore, the microstructure around the white band structure consists of fine basket-weave, colony α and coarse basket-weave in sequence from top to bottom. The white band region is a slight hardness drop (~20 HV0.2), but it has almost no effect on tensile strength. Keywords: WAAM; Deposition path; Formability; Microstructure; Mechanical properties.

1. Introduction Wire and arc additive manufacturing (WAAM) has increasingly drawn significant attention in the industrial manufacturing field, due to its advantages including high deposition rates, high material utilization and the potential for large part sizes compared to conventional manufacturing technologies [1-3]. In WAAM system, the heat source is usually Gas Tungsten Arc welding (GTAW) [4-6], Gas Metal Arc Welding (GMAW) [7-9], or Plasma Arc Welding (PAW) [10,11]. Generally, the GTAW-based WAAM is more stable and generates less weld fume and spatter [5].

The WAAM technology is especially expected to apply in the aerospace field whose components are machined from costly wrought material at a low fly to-buy ratio [12,13]. Ti-6Al-4V alloy is the most commercially produced titanium alloy, which is the most commonly applied in the aerospace field [14]. Precisely because of this, WAAM Ti-6Al-4V has been attracted the attention of researchers. How to produce high-quality Ti-6Al-4V components with comparable mechanical properties as forgings is particularly important for the application of WAAM in the aerospace field [1]. Mechanical property is related to microstructural morphology. For a Ti-6Al-4V component which is produced by WAAM, the microstructures are complex, often varying spatially within the deposition. The microstructure of this α/β alloy is very sensitive to the thermal history, and different microstructures such as lamellar, equiaxed or bimodal microstructure can be obtained [15]. Baufeld et al. [16,17] studied microstructure characteristic of Ti-6Al-4V components where the Widmanstätten microstructure consists of fine α lamellae in the top and coarse lamellae in the bottom region, and the influence of different heat treatment on microstructure and mechanical properties. Wang et al. [2] and Baufeld et al. [18] studied the microstructure and mechanical properties of different locations. Wang et al. [19] studied grain morphology evolution and texture characterization of Ti-6Al-4V titanium

walls. Wu et al. [20] reported that the effects of various interpass temperatures on microstructural morphology, grain size and mechanical properties. Foster et al. [21] reported that increasing dwell time results in a slight decrease in the α lath widths and a much more noticeable decrease in the width of prior-β grains. Martina et al. [22,23] found the application of rolling can induce grain refinement and a modification of the microstructure from strongly columnar to equiaxed. To date, numerous investigations have been carried out to reveal microstructural formation mechanisms and explore the influence of different technological means on microstructure and mechanical properties, but more elaborate research of microstructural morphology is still essential. Furthermore, to the authors' best knowledge there is little literature available which investigates the effect of deposition path on the microstructure and mechanical properties of Ti-6Al-4V produced by WAAM. The objective of this study is to comprehensively investigate thermal cycle,

formability,

microstructural

characteristic

and

mechanical

properties of the Ti-6Al-4V thin-wall parts produced by GTAW-based WAAM. The difference of macrostructure, microstructure and mechanical properties in various regions of the thin-wall parts deposited by various paths and heat input are elaborately characterized and comparably analyzed.

2. Experimental procedures 2.1. Experiment setup The WAAM system employed for this investigation is shown in Fig. 1. The WAAM system consists of gas tungsten arc welding (GTAW) unit, water cooling unit, oxygen and water purification unit and motion control unit with three freedoms of X-, Y- and Z-axis linear movement. In reconstructive glove box with purification system, the GTAW torch is installed on Z-axis at a suitable position and Ti-6Al-4V wire with a diameter of φ1.0 mm is fed by a wire feeder.

Fig. 1. Schematic diagram of WAAM system

After filling with an inert gas (Ar, 99.9%), the reconstructive glove box can hold H2O and O2 content less than 1 ppm by gas recirculation and purification, respectively. Ti-6Al-4V substrate with the dimensions of 150 mm×50 mm×5 mm is cleared by ethyl alcohol and clamped on the

work table. The chiller outside the reconstructive glove box provides the cooling water of 6℃ by a water pipe for work table and welding torch constantly. The thin-wall parts with co-directional and reciprocating deposition paths, which can be seen in Fig. 2, were prepared under the different heat inputs for each deposition path, and the detailed deposition parameters of the thin-wall parts are list in Table 1.

Fig. 2. Diagrammatic drawing of the distribution with test locations Table 1. Deposition parameters of thin-wall parts and geometric data Component name Current(A) Travel speed(mm/min) Wire feed speed(mm/min) Flow rate of argon(L/min) Dwell time between layers(s) Number of deposited layers Wall height (mm) Wall length (mm) Wall width (mm) Heat input(J/mm)

A 160 200 1300 8 180 20 25 90 9.660 708.4

B 180 180 1600 8 180 20 25 90 10.102 928.8

C

D

160 200 1300 8 180 20 24 90 9.504 708.4

180 180 1600 8 180 20 24 90 9.600 928.8

Deposition path

co-directional

reciprocating

2.2. Thermal cycle simulation methods The finite element software package, ANSYS 15.0, was used for the thermal analysis. Thermal cycle of the WAAM process was calculated based on the parameters in Table 1. The thermal model using eight nodes solid element (Solid 70) which can transfer heat in three directions, and the total number of element is 26910. The WAAM material was simulated using the ‘‘element birth technique’’. All the elements of the bead were deactivated at the first step of the analysis, and then the elements were activated sequentially following the heat source [24]. The Goldak double ellipsoidal heat source was used to apply the heat to the additive manufacture deposits. Newton’s Law for surface convection heat loss and Stefan-Boltzmann’s Law for radiation heat loss were considered as boundary conditions for thermal analysis. Finally, the full transient integral method was used to solve. 2.3. Material characterization methods Fig. 2 shows the distribution of each position for all analysis. Five side lines (L1-L5) of different positions evenly distributed in the length direction on each deposited thin-wall part were selected to measure surface waviness (SW) and total wall width (TWW). SW of the thin-wall parts deposited by WAAM was measured by a Mahr XC20 surface profiler, and TWW was measured by the micrometer. SW is evaluated by the maximum height of a waviness profile, and effective wall width

(EWW) can be calculated by SW= (TWW–EWW)/2. Maximum utilization ratio, the ratio of EWW and TWW, was calculated by these data. Microstructure, X-ray diffraction (XRD) and microhardness were all measured in cross-section, and the position of measuring side macrostructure next to it. It is well known that Ti-6Al-4V components produced by WAAM can be highly anisotropic, so in order to better highlight the differences among the samples, the orientation selected for testing were in the top horizontal direction and vertical direction, which were frequently reported to be the least ductile direction and the maximal ductile direction respectively in Ti-6Al-4V thin-wall parts manufactured by WAAM in studies previously conducted by Wang et al. [2], Baufeld et al. [18], and Bermingham et al. [25]. The different orientations are designated with H and V, standing for specimens with tensile direction horizontal (H) and vertical (V) to the deposition plane. Accordingly, label A-H denotes a specimen deposited by co-directional path with tensile direction horizontal to the deposition plane. Three dog bone-shaped tensile specimens were prepared in each group with a 1.0 mm×2.5 mm cross-section, as shown in Fig. 2. The samples for microstructure analyzing were ground with SiC papers (150, 400, 600, 1000, 1500, 2000, 2500 grit), and then polished with 2.5µm and 1.5µm diamond polishing fluid for about 3–5 min,

respectively. After polishing, the samples were etched for 15–20 s in Kroll reagent containing 2% vol HF, 6% vol HNO3 with balance H2O for microstructure observation by optical microscope (OM, Zeiss Axiovert 200 MAT) and scanning electron microscopy (SEM, Hitachi S4800). The phase orientations were identified by X-ray diffraction (XRD, Rigaku Smartlab) using CuKα radiation at an accelerating voltage of 40 kV and a current of 30 mA with a scanning step of 0.02° and a scanning speed of 1°/min in the range of 30°-80°. Vickers hardness of the deposited thin-wall parts was measured by microhardness tester (FM-ARS 9000) with the load of 200 g and dwell time of 10 s on polished cross-section. The tensile tests were carried out with universal testing machine (Instron 5848 micro-tester) at a displacement speed of 1 mm/min at room temperature. The fracture surfaces of tensile specimens were analyzed by scanning electron microscopy.

3. Results and discussion 3.1. Thermal cycle Fig. 3 shows the thermal history at the middle of the deposited thin-wall parts at the 4th layer, the 10th layer and the 17th layer based on the deposition parameters which list in Table 1. The deposited titanium alloy was experienced multiple thermal cycles with a reducing peak temperature during the WAAM process. The

finite element model results show that the deposited temperature of sample A and sample C exceeded the β phase transition temperature and thus underwent a complete transformation from liquid to β phase during the first four thermal cycles. However, sample B and sample D experienced five thermal cycles due to the higher heat input. Furthermore, the melted titanium alloy may not only consists of the deposited material in the current layer but includes one of the previous depositions. In sample A and sample C, the deposited material in the surface of the previous layer was re-melted by the next deposited layer, and the next deposited layer even can re-melt the whole last deposited layer in sample B and sample D.

Fig. 3. WAAM Ti-6Al-4V thermal history calculated from the finite element model at the 4th layer, the 10th layer and the 17th layer. (a) the thermal history from sample A; (b) the thermal history from sample B; (c) the thermal history from sample C; (d) the

thermal history from sample D.

It can be seen that the peak temperature of sample B and sample D is about 2000℃, which is higher than that of sample A and sample C. While, there is no great difference between co-directional path and reciprocating path.

3.2. Formability Fig. 4(a-e) shows the measuring results of the surface profile from four samples. Fig. 4(f) shows SW of five side lines of different positions uniformly distributed in the WAAM thin-wall parts. It can be clearly seen that SW in both ends are higher than that in the middle from sample C and sample D, and SW from sample A and sample B gradually rise from arc starting to arc stopping. Due to nodulation will be produced in arc starting and arc stopping, the position of arc starting is higher, and that of arc stopping is lower, which affect the appearance of outline and the surface roughness of profile. Different layer thicknesses affect the arc length, which can determine the metal transfer mode: no-droplet mode, tangent-droplet mode, or no-contact mode [26]. The no-contact mode in the position of arc stopping easily results in spatters and collapse, which lead to high roughness of side surface after multilayer co-directional deposition. And the roughness of the side surface of both ends from samples by reciprocating deposition path is high due to alternating the position of arc starting and arc stopping.

Fig. 4(g) shows EWW and the maximum utilization ratio of the WAAM thin-wall parts at different samples. The maximum utilization ratio of sample A and sample C are slightly higher. This difference can be attributed to the bigger waviness of sample B and sample D due that the whole last deposited layer was re-melted with the deposition of the next layer.

Fig. 4. The measuring and analyzing results of the surface profile. (a-e) L1-L5 profile curves of four samples; (f) SW in five positions of four samples; (g) Maximum utilization ratio and EWW of four samples.

Fig. 5 shows the cross-sectional macrostructure of the WAAM Ti-6Al-4V thin-wall parts in YZ plane and XZ plane. The cross-sectional macrostructure of the parts in YZ plane exhibits a layered structure in which large β grains homogeneous distributes pass through the layer. Prior-β grains from the thin-wall parts deposited by co-directional path grow epitaxially, which is inclined relative to the Z direction following by

the maximum thermal gradient resulting from the moving welding torch, as shown in Fig. 5(a) and (b). In the top and middle regions, features are obvious due to columnar large β grains. However, in Fig. 5(c) and (d), the growth direction of the prior-β grains reveals less apparent by comparison. Close inspection reveals that many grains are winding due to different arc moving direction in adjacent layers. Therefore, it can be considered that a competitive relationship exists among the different growing orientations of prior-β grains with additive layers deposition.

Fig. 5. The cross-sectional macrostructure of four samples in YZ plane and XZ plane.(a-d) samples A, B, C and D in YZ plane; (e-h) samples A, B, C and D in XZ plane.

Due to the influences of thermal gradient and solidification rate, the size of β grains gradually increases in the WAAM thin-wall parts from the bottom region to the top region. In the bottom region of the parts, refined β grains form due to the fast cooling rate. During the WAAM

process, the β grains at the edge of the fusion boundary acted as nucleation sites from which the solidification front grew back, epitaxially, into the weld pool, where each growing grain formed as a continuation of the grains that lie along the fusion boundary [2]. Consequently, coarser β grains formed in the subsequent layers. The cross-sectional macrostructure of WAAM Ti-6Al-4V thin-wall parts in XZ plane are shown in Fig. 5(e-h), which reveals that three distinct regions are discernible in the macrostructure of four samples. In the top region of the thin-wall parts, there is no noticeable layered structure. While the macrostructure of the thin-wall parts show horizontal white bands in the middle region and concave white bands in the bottom region. Compared with the surface waviness profiles of the WAAM thin-wall parts, it can be concluded that the position of the parallel bands, which is not directly equal to the position of deposition layers, depends on the temperature field of the WAAM process. During the initial deposition process, the Ti-6Al-4V substrate at the fusion boundary was heated to the liquidus temperature. Heat on the substrate can be transmitted to both sides, not just downward. Therefore, the bands in the bottom region of the part are concave. With the thin-wall part deposited layer by layer, the repeated thermal cycling effect is equivalent to the repeated heat treatment process of the subsequent layers on the

previous ones, which in turn leads the α phase band to coarsen. There are no obvious bands on the last few layers in the top region due that the layers at this position have experienced relatively incomplete thermal histories. 3.3. Microstructure Due that there is no obvious difference in microstructure among four thin-wall parts, Fig. 6 only shows the microstructural distributions of well-formed sample A in XZ plane along the building direction. The microstructures in three regions of the thin-wall part consist of Widmanstätten structures, colony structure, basket-weave structure and martensite α′ in the β matrix. Due to the influences of thermal gradient and solidification rate, in the top region of the thin-wall part, the acicular α interwoven with a basket-weave structure and martensite α′ in the β matrix were generated (Fig. 6(a)). The lamellar α interwoven with a basket-weave structure was produced in the middle region (Fig. 6(c)) and the Widmanstätten structure (Fig. 6(e)) was formed in the bottom region.

Fig. 6. The microstructural distributions of thin-wall part. (a) the top region of low-magnification; (b) the top region of high-magnification; (c) the middle region of low-magnification; (d) the middle region of high-magnification; (e) the bottom region of low-magnification; (f) the bottom region of high-magnification.

The sizes of α phase lamellae, which are distributed in the β phase matrix, are obviously different in three regions, as shown in Fig. 6(b), (d) and (f). It can be seen that the width of α plates in the middle region is much wider than that in other regions. The produce of fine or coarse microstructure is mostly related to the thermal history, which mainly depends on the heat treatment process. The formation of coarse lamellae in the middle region of the thin-wall part is related to the repeated heat treatment during the WAAM process. In the top region, the

microstructure consists of very fine lamellae with incomplete thermal histories relative to previous layers. The microstructure in the middle region of the thin-wall part was subjected to subsequent heat treatments in the α/β phase field allowing diffusive element partition, which results in a coarsening α lamellae [16,27]. Although the bottom region of the thin-wall part experienced completed thermal histories, it can be obtained a faster cooling rate due to the heat transmitting boundary conditions, which only resulted in part coarsen lamellae microstructure. The α lamellae grew coarsening in the form of colony firstly (Fig. 6(f)), and there was no enough time for a complete transformation due to the fast cooling rate. The mixture of martensitic and basket-weave structure (Fig. 6(a)) can be observed in the top region. Usually, the martensite α′ structure is characterized by a rectangular grid structure [28]. Due to the high cooling rate in the first few layers, the heat can be dissipated from the substrate by direct contact, and martensite α′ was formed into lathlike matrix structure. However, as the WAAM deposition process, the martensite α′ structure decomposed gradually with multiple thermal cycles. It has been reported that martensite α′ can be changed to fine lamellar α+β microstructures by annealing in the temperature region of around 700-850 ℃ [28]. In the last few layers, the martensite α′ structure was retained on account of incomplete thermal histories relative to previous

layers. The microstructures in every parallel region of the WAAM thin-wall parts are not significantly different due to the similar heat dissipation behavior at the same location within each of the produced parts. As can be seen from the optical microscope (Fig. 6(c)), the color of each layer changes from light to dark along with the building direction. That is related to the microstructure characteristic. The microstructures around the white band region with samples by co-directional and reciprocating deposition paths (Fig. 7 (a) and (b)) are all consisted of fine basket-weave, colony α and coarse basket-weave in sequence from top to bottom. The diagrammatic drawing of that is shown in Fig. 7 (c). The white band structure, which is about 130 µm in width, derives from the β/α transus lines (Fig. 3) during subsequent depositing. The heat treatment below α transus lines resulted in a thickening of lamellar α. However, the heat treatment in the α+β phase field produced a structure with bundles of α laths (colony structure), and the heat treatment above β transus lines produced a fine structure with α laths. The changes are not obviously in the white band structure between sample A and sample C. Therefore, the deposition path contributes a minimal effect to the white band structure markedly.

Fig. 7. The microstructures around the white band region (a) the sample by co-directional deposition path; (b) the sample by reciprocating deposition path; (c) diagrammatic drawing.

3.4. XRD The preferred orientation and the crystal phase of the deposited Ti-6Al-4V alloy are detected and estimated by XRD measurement at the three regions of the WAAM thin-wall parts under different deposited process. In order to make the XRD analysis more differentiated, the XRD samples were selected at the near center of the additive manufacturing samples. It is observed that the diffraction pattern of the thin-wall parts presents six obvious peaks in Fig. 8.

Fig. 8. XRD patterns of three regions from samples A, B, C and D

As can be seen from in the XRD patterns, the peaks of β {110} plane are almost invisible compared to α {101} peaks. That is attributed to the fact that the β grain sizes are too large to obtain sufficient data, but there are more β phases that can be detected by X-ray due to relatively small grain size in bottom regions. However, there are still tiny compared to α {101} peaks. There are 12 possible α orientations (each satisfying the Burgers orientation relationship) of which can be transformed into a single parent β grain [11,19]. In the bottom regions of the thin-wall parts, the results of XRD are similar to that of TC4 alloy due that a considerable part of the materials which are not re-melted and heated above β transus lines repeatedly. Some reports showed preferred crystallographic orientation of TC4 [5,29]. During the phase transition process, the Burgers relationship

between the two phases of {0001}α∥{110}β and〈110〉α∥〈111〉β was strictly followed. Therefore, α phase has a preferred crystal orientation perpendicular to the α {101} planes in this region. In the top regions of the thin-wall parts, α diffraction peaks in this region are different from the bottom regions, and the strongest peaks appear in different positions. As shown in Fig. 5, there are no white band structure due that the reheat temperature always exceed the β transus line. Lonardelli et al. [30] has found that the nucleation of new β grains from the matrix followed Burgers orientation relations, and the new β grains favors the grain growth of residual β phase orientations. Different α orientations can occur during the heating and the cooling phase transformations. The proportion of variants which can be modified during nucleation and growth determines the phase characteristics. However, the middle regions of the thin-wall parts experienced more thermal cycles. After repeated remelting and reheating, the proportion of different α diffraction peaks became more similar. Careful observation can be found that the XRD patterns of sample C, D in the top and middle regions have slight peak shifting to the high angle , and it can be indicated that the lattice constant of the samples are decreased. The heat concentration at the near center of the samples by reciprocating deposition path is much higher than that by co-directional deposition path due to the reciprocating thermal cycle, which in turn

increases the temperature of the samples at the near center and prolongs the cooling process. The defects, such as dislocation, could be slipped sufficiently, which eventually leads the defects reduction and shifts the XRD peaks to high angle slightly during the thermal cycle of WAAM process. 3.5. Hardness Fig. 9 shows the microhardness distribution along the vertical centerline of the cross-sectional Ti-6Al-4V thin-wall parts deposited under different process conditions. It can be clearly seen that the microhardness in the bottom regions and top regions is slightly higher than that in the middle regions. The amount of impurities, especially oxygen, can influence obviously the hardness of Ti-6Al-4V [28,31]. But the samples were deposited in the environment with less than 1 ppm of H2O and O2. So the microhardness distribution variation of the thin-wall parts may be produced by other causes. Furthermore, the hardness of WAAM Ti-6Al-4V is mainly determined by solid solution and grain boundaries due to some segregated elements interacting between grain boundaries existing in α/β microstructures [20]. The alloys in the bottom region have a faster cooling rate, which brings more grain boundaries and dislocations

to

produce

higher

microhardness.

Moreover,

the

microstructure of the WAAM thin-wall parts in the top region consists of large amounts of martensite α′ structure, which are usually harder and

have higher strength than that in the middle region. The results indicate that there is no obvious change in average microhardness values in both depositing paths. While, the average microhardness of thin-wall parts with high heat input is slightly higher than that with low heat input, which may depend on the solid solution hardening. The alloy of the thin-wall parts with high heat input experienced more thermal cycles above the β transus temperature, which caused more alloying elements to melt into the β phase. The solid solution hardening has a more significant impact on microhardness within a certain peak temperature and cooling rate.

Fig. 9. The microhardness distribution along the vertical centerline of the cross-sectional Ti-6Al-4V thin-wall parts. (a) the samples by co-directional deposition path; (b) the samples by reciprocating deposition path.

Fig. 10 shows the microhardness distribution around the white band

structure. There is a notable change in microhardness around the white band structure due to markedly different appearances. Obviously, the microhardness in the white band region is lower than that of surrounding regions. No clear dependence on the deposition paths can be reported. There are a large number of α colonies in the white band region, which leads to the increase of size of actual slip on strain [11]. So the microhardness relatively decreases. There is no obvious change on the microhardness despite different lamellae widths above and below the white band. Therefore, it is reasonable to conclude that the lamellae width is almost not responsible for the hardness, but phase morphology can affect hardness.

Fig. 10. The microhardness distribution around the white band region. (a) sample A; (b) sample B; (c) sample C; (d) sample D.

3.6. Tensile properties The tensile test results of all samples are summarized in Fig. 11. The

ultimate tensile strength (UTS) varies between 826 and 948 MPa, depending mainly on the orientation and heat input of the tensile specimens. The elongation of the tensile specimens generally varies between 10 and 15%. In the comparison of the specimens tested from the horizontal direction, those from the vertical direction exhibited a higher ductility, with a mean tensile elongation of 13.83%, and slightly lower strength, with a mean UTS of 859 MPa. A significant difference in the ductility data is only seen between horizontal direction and vertical direction. Ductility is much worse when tested in the horizontal direction, transverse to the direction of β grain growth. This orientation relationship of the ductility mainly results from the anisotropic property of the WAAM process. The columnar β grains grow parallel similarly to the building direction, it is loaded transverse to the β grain boundaries when measured in the horizontal direction, this would be expected to promote premature failure through the grain boundary [32].

Fig. 11. Tensile properties of samples. (a) tensile strength; (b) elongation.

Fig. 12 shows the engineering stress-strain curves of WAAM specimens for the different process parameters and orientations. It can be seen that the curves of the vertical specimens in the plastic deformation stage are relatively stable. However, the plastic deformation section of the horizontal specimens shows a fast decreasing trend. The result of hardness shows that the white band regions have lower hardness, and it can be deduced that the plastic deformation of white band regions takes precedence over other regions. The result corresponds well to the low yield strength of the vertical specimens. Meanwhile, the working hardening occurs in the white band regions during the plastic deformation process. When the strength of the white band regions exceed the other regions, the plastic deformation will be dispersed to other regions. In

contrast, the white band regions of the horizontal specimen are restrained by non-white band regions during the stretching process, and the plastic deformation of white band regions cannot be occurred preferentially. The complicated microstructure of additive manufacturing results in obvious fluctuation of stress-strain curve of horizontal specimens, and the error range increases obviously.

Fig. 12. Stress–strain curves of vertical specimens from samples A-D (a-d) and horizontal specimens from samples A-D (e-h).

This is not that surprising as when loaded transverse to the prior-β grain boundaries, which would be expected to promote premature failure through the grain-boundary a layer [2]. The reason of different UTS between samples of two heat input agrees with the hardness in this study. No significant difference of UTS can be observed between samples of two deposition paths except tensile specimens extracted from the vertical direction. That could be attributable to different fracture pattern.

Representative fractographs of the tensile fracture surface of specimens C-H, D-H, B-V and D-V are shown in Fig. 13. The fracture surfaces of specimens are basically characterized by dimple-like structures, which are indicative of ductile rupture. The difference of fracture surfaces is the feature of dimples, which can reflect their mechanical properties. Many researches indicated that the small size of dimples leads to high tensile strength and low tensile ductility, and the shallow dimples reflect the drop in elongation for the specimens [32].

Fig. 13. SEM micrographs of tensile fracture surfaces of (a) specimens C-H; (b) specimens D-H; (c) specimens B-V; (d) specimens D-V.

Specimen C-H (Fig. 13(a)) and specimen D-H (Fig. 13(b)) show shallower dimple compared to specimen B-V (Fig. 13(c)) and specimen D-V (Fig. 13(d)), which is indicative of its relatively low ductility. From the specimen D-H, it is noticeable that some torn edges can be observed, showing more brittle cleavage morphology. That is consistent with the result of high tensile strength and low ductility, as shown in Fig. 11.

Fig. 14. Optical macrostructure image of the fracture plate tensile specimens and corresponding overall view for fracture surfaces of SEM micrographs. (a, b) specimens B-V; (c, d) specimens D-V.

It is found in Fig. 14(a) and (c) that the cracks of specimen B-V and specimen D-V do not extend along the white band regions. Therefore, the white band regions will not have a significant negative impact on the tensile properties of thin-wall parts. It is noted that the white band region in the thin-wall parts is very narrow (~130 µm) and a slight hardness drop (~20 HV0.2) at this region is attained. During loading, the deformation of this softened band region was restrained by the fine basket-weave region, resulting in a stress triaxiality region produced at the softened region next to the fine basket-weave region [33]. Therefore, in case of such stress triaxiality, this softened band region during the plastic deformation had a higher work hardening rate. The difference between specimen B-V and specimen D-V is attributed to many winding grain boundaries in sample

D, which resulted in the stress concentration around grain boundary tips in regions of coarse lamellae, promoting crack initiation. Compared to Fig. 14(b), numerous crack sources can be observed in Fig. 14(d) which reduce the tensile strength.

4. Conclusions In this study, Ti-6Al-4V thin-wall parts were fabricated by WAAM in an inert environment. Through the analysis of thermal cycle, formability,

microstructure

and

mechanical properties, following

conclusions can be drawn: (1) SW in both ends are higher than those in the middle from samples by reciprocating deposition path, and SW from samples by co-directional deposition path gradually rise from arc starting to arc stopping. Prior-β grains from samples by co-directional deposition path are inclined relative to the building direction, and there are many winding prior-β grains in samples by reciprocating deposition path due to different direction of the maximum thermal gradient. (2)

The microstructures of different deposition paths are almost basket-weave structure, the Widmanstätten structure and martensite α′ in the β matrix. The microstructures around the white band structure consist of fine basket-weave, colony α and coarse basket-weave in sequence from top to bottom.

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Figure captions Fig. 1. Schematic diagram of WAAM system Fig. 2. Diagrammatic drawing of the distribution with test locations Fig. 3. WAAM Ti-6Al-4V thermal history calculated from the finite element model at the 4th layer, the 10th layer and the 17th layer. (a) the thermal history from sample A; (b) the thermal history from sample B; (c) the thermal history from sample C; (d) the thermal history from sample D. Fig. 4. The measuring and analyzing results of the surface profile. (a-e) L1-L5 profile curves of four samples; (f) SW in five positions of four samples; (g) Maximum utilization ratio and EWW of four samples. Fig. 5. The cross-sectional macrostructure of four samples in YZ plane and XZ plane. (a-d) samples A, B, C and D in YZ plane; (e-h) samples A, B, C and D in XZ plane. Fig. 6. The microstructural distributions of thin-wall part. (a) the top region of low-magnification; (b) the top region of high-magnification; (c) the middle region of low-magnification; (d) the middle region of high-magnification; (e) the bottom region of low-magnification; (f) the bottom region of high-magnification. Fig. 7. The microstructures around the white band region (a) the sample by co-directional deposition path; (b) the sample by reciprocating deposition path; (c) diagrammatic drawing. Fig. 8. XRD patterns of three regions from samples A, B, C and D Fig. 9. The microhardness distribution along the vertical centerline of the cross-sectional Ti-6Al-4V thin-wall parts. (a) the samples by co-directional deposition path; (b) the samples by reciprocating deposition path. Fig. 10. The microhardness distribution around the white band structure. (a) sample A; (b) sample B; (c) sample C; (d) sample D. Fig. 11. Tensile properties of samples. (a) tensile strength; (b) elongation. Fig. 12. Stress–strain curves of vertical specimens from samples A-D (a-d) and horizontal specimens from samples A-D (e-h). Fig. 13. SEM micrographs of tensile fracture surfaces of (a) specimens C-H; (b) specimens D-H; (c) specimens B-V; (d) specimens D-V. Fig. 14. Optical macrostructure image of the fracture plate tensile specimens and corresponding overall view for fracture surfaces of SEM micrographs. (a, b) specimens B-V; (c, d) specimens D-V.

Table 1. Deposition parameters of thin-wall parts and geometric data

Component name Current(A) Travel speed(mm/min) Wire feed speed(mm/min) Flow rate of argon(L/min) Dwell time between layers(s) Number of deposited layers Wall height (mm) Wall length (mm) Wall width (mm) Heat input(J/mm) Deposition path

A

B

160 180 200 180 1300 1600 8 8 180 180 20 20 25 25 90 90 9.660 10.102 708.4 928.8 co-directional

C

D

160 180 200 180 1300 1600 8 8 180 180 20 20 24 24 90 90 9.504 9.600 708.4 928.8 reciprocating

Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: