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Effects of selective laser melting build orientations on the microstructure and tensile performance of Ti–6Al–4V alloy
Materials Science & Engineering A 776 (2020) 139001
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Effects of selective laser melting build orientations on the microstructure and tensile performance of Ti–6Al–4V alloy Zongyu Xie a, Yu Dai b, Xiaoqin Ou a, *, Song Ni a, Min Song a a b
State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, PR China School of Materials Science and Engineering, Nanchang University, Nanchang, 330031, PR China
A R T I C L E I N F O
A B S T R A C T
Keywords: Selective laser melting Ti–6Al–4V alloy Build orientation Phase composition Tensile performance
In this paper, Ti–6Al–4V alloy was fabricated by selective laser melting (SLM) with different build orientations, namely SLM-1 and SLM-2, respectively. Microstructures of SLM-1 and SLM-2 were characterized by X-ray diffraction, optical microscope, scanning electron microscope and transmission electron microscope. The results show that the α and α0 phases in the hexagonal close-packed (hcp) structure, as well as the residual β phase in the body centered cubic (bcc) structure, exist in both samples. However, different SLM build orientations led to varied cooling rates and α/α0 ratios in the final materials. The SLM-1 alloy with build orientation along its thickness direction had higher cooling rate and thus lower α/α0 ratio than the SLM-2 alloy with build orientation along the length direction. Besides, minor martensitic (α’’) phase was observed in the SLM-1 alloy and acted as an intermediate phase between β and α0 phases. The low α/α0 ratio and existence of α’’ phase in the SLM-1 alloy result in a relatively high tensile strength but deteriorated plasticity. Comparatively, the higher α/α0 ratio in the SLM-2 alloy without α’’ martensite leads to a much better plasticity regardless of the slightly reduced tensile strength.
1. Introduction As a near net shape manufacturing technology, additive manufacturing (AM) can save raw materials, improve efficiency and reduce post-processing [1]. Selective laser melting (SLM) is one of the most mature AM technologies due to its unique advantages, such as short lead time and high design freedom [2,3]. The SLM method has been successfully employed in industry to manufacture metallic mate rials, such as Ti alloys [4–6], Al alloys [7–9] and steels [10–13]. Over the last decades, titanium alloys were widely used in aerospace, ship building and biological bone field due to their outstanding properties including high strength-to-weight ratios, good corrosion and heat resistance, low elasticity modules and superior biocompatibility [14]. Nonetheless, the development of titanium alloys manufactured by traditional methods was severely restricted by their high post-processing cost, where the near net shape SLM technology shows special advantages. As the first commercial titanium alloy, Ti–6Al–4V alloy can be a good representative for SLM titanium alloys. Table 1 summarizes the micro structure and tensile properties of the Ti–6Al–4V alloy manufactured by different technologies, such as the select electron beam melting (SEBM)
method [15], solution treated and aged (STA) method [16], mill-annealed (MA) method [16] and SLM method [6,8,17–23]. It is found that the conventional Ti–6Al–4V alloy is composed of the hex agonal close-packed (hcp, α) and body centered cubic (bcc, β) structures [24,25]. However, martensite, such as orthorhombic-α’’ and HCP-α0 phases, may also coexist in the Ti–6Al–4V alloy when it is fabricated by the SLM method. The formation of those two martensitic phases results from the non-diffusional phase transition between the β and α structures. The existence of martensitic structure benefits the tensile strength but lowers the elongation of the SLM Ti–6Al–4V alloy [18]. Accordingly, it is seen from Table 1 that the SLM Ti–6Al–4V alloy shows high tensile strength but relatively low ductility. For example, Gong et al. [19] found that the tensile strength of SLM Ti–6Al–4V alloy reaches 1250 MPa, which is much higher than that of the wrought, STA and MA Ti–6Al–4V alloys. For the SLM Ti–6Al–4V alloy, part of the investigations [7,17,18] focused on the post-processing procedure to improve its combinational properties of strength and ductility, while the others [6,19–21] aimed to improve alloy performance by optimizing key SLM parameters, i.e. laser power (P), layer thickness (τ), laser scanning speed (v), hatch spacing (h) and focal offset distance (FOD). The traditional post-processing method
* Corresponding author. E-mail address: [email protected] (X. Ou). https://doi.org/10.1016/j.msea.2020.139001 Received 14 November 2019; Received in revised form 9 January 2020; Accepted 22 January 2020 Available online 23 January 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.
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w.t.% V) powder used in this research has an average particle size of 36 μm. The Ti–6Al–4V alloy was prepared using the selective laser melting (SLM) machine (SLM®280 2.0) manufactured by SLM Solution Group AG. Before machine working, the powder bed was preheated to about 200 � C, and the building chamber was filled with argon to ensure the content of oxygen below 80 ppm. The present study chooses the orthogonal band scanning mode. The scanning angle rotates by 37� around the Z-axis for each layer. The following SLM parameters are used: layer thickness (30 mm), laser power (Core: 950 W; Outer Hull: 350 W), laser scanning speed (Core: 230 mm/s; Outer Hull: 1100 mm/s), hatch spacing (0.06 mm), FOD (0 mm). These SLM parameters in present study are selected according to the Material Data Sheet from SLM so lutions Group AG [30], which includes proper combination schemes of key SLM parameters in practice to obtain optimal mechanical properties of Ti–6Al–4V alloys. The schematic picture in Fig. 1a and b shows the two samples with different SLM build directions in present study, namely SLM-1 and SLM-2, respectively. The SLM-1 sample in Fig. 1a was built along the thickness direction (L1), while the SLM-2 sample in Fig. 1b was built along the length direction (L2). Phase compositions in the printed samples were tested by X-ray diffraction (XRD) spectroscopy (Cu Kα radiation, 0.02� step size, 2� per min scanning rate). Phase fractions of the as-build alloys were calculated based on the XRD results using the MAUD software [31]. The as-build samples were machined into tensile specimens with the gauge size of 15 mm � Φ 4 mm. Tensile tests were performed with a constant loading speed of 1 mm/min at room temperature using an Instron-3369 uni versal testing instrument. Microstructures of samples were observed using the optical microscope (OM), FEI Titan G2 60–300 transmission electron microscope (TEM) equipped with an energy disperse spectro scope (EDS), and Nova Nano SEM320 scanning electron microscope (SEM) equipped with an electron backscatter diffraction (EBSD) detec tor. The samples for OM observation were polished and then etched using the Kroll’s reagent. The samples for the TEM and EBSD analysis were prepared as follows: slices with 0.5 mm in thickness were cut from the as-build samples, where the section planes of the slices were parallel to the XY plane for SLM-1 and the XZ plane for SLM-2, respectively, as illustrated in Fig. 1. Then, the slices were mechanically polished to a thickness of 60 μm, followed by twin-jet electropolishing in the elec trolyte composed of 35% normal butanol, 5% perchloric acid and 60% methyl alcohol solution at 25 � C under a voltage of 30 V.
Table 1 Summary of Ti–6Al–4V alloys manufactured by different technologies, where α, α0 and β phases represent the hcp, hcp and bcc structures, respectively. Process
Microstructure
Mill-annealed Solution treated þ aged SLM SLM SLM SLM SLM SLM þ HT SLM SLM þ HIP(920 � C,100 MPa) SLMþ400 � C/2h SLMþ800 C/2h �
SLMþin-situ heat treatment LMD SEBM
Tensile properties
Ref.
UTS (MPa)
E (%)
Globular β in α matrix Fine (αþβ) lamellar
– –
– –
[16] [16]
Fine acicular α0 in columnar original-β grains Fine acicular α0 in columnar original-β grains Very fine (αþβ) lamellar – Fine acicular α0 in columnar original-β grains (αþβ) lamellar and globular
–
–
[6]
1240–1250
6.6–8.8
[19]
1180 1250 1095
11.4 6 8.1
[20] [21] [17]
990
11
[17]
Very fine acicular α0 α lamellas within β matrix
1090 960
10 14
[22] [8]
Fine (αþβ) lamellar in columnar original-β grains (αþβ) lamellar þ globular α at original-β grain boundaries Fine (αþβ) lamellar in columnar original-β grains
1051
13
[18]
1069
10
[18]
1210
11
[18]
–
–
[23]
–
–
[15]
α
Acicular α in columnar priorβ with grain boundary α (αþβ) lamellar
UTS: Ultra tensile strength E: Elongation.
can improve the performance to some extent, which is nonetheless negligible compared to the significant performance promotion resulted from the microstructural optimization by a slight change in SLM pa rameters [26,27]. Xu et al. [18] found that the laser energy density (E ¼ P 0 v⋅h⋅τ) are closely correlated to the formation of the acicular α phase and ultrafine lamellar (αþβ) structure. Specifically, a lower laser energy density produces a higher content of martensitic α0 phase in the SLM Ti–6Al–4V alloy [18]. Besides, it is reported by Zafari et al. [28] that the increase in FOD, which is a variable rarely used by the SLM community, would reduce the energy density during the SLM process. It should be noted that it is inadequate to adjust a single processing variable to control the laser energy density. Instead, it is necessary to design a proper combination of the key SLM processing variables in order to control the microstructure and further the mechanical performance of bulk materials. The SLM process is a rapid solidification process with the cooling rate ranging from 103 to 105 K/s [15], which is far beyond the critical cooling rates for β→α transformation (~20 � C/s) and for β→α’/α’’ martensitic transformation (~410 � C/s) [29]. Therefore, the fast cooling rate during the SLM process will induce martensitic transformation from β to α0 or α’’ phases in the bulk Ti–6Al–4V alloy. In general, the SLM Ti–6Al–4V alloy consists of acicular α0 phase and columnar β phase, whereas it may be composed of fine (αþβ) lamellae occasionally. The α’’ phase is rarely observed in SLM Ti–6Al–4V alloy due to the low content of β-phase stabilizing elements [14]. Except for the key parameters mentioned above, build orientation will also have influence on the microstructure and mechanical perfor mance of SLM titanium alloys. Thus, this paper focuses on the effects of different build orientations on the microstructure and mechanical per formance of SLM Ti–6Al–4V alloy, and the more superior build orien tation of the alloy is determined.
Fig. 1. Schematic illustrating: (a) SLM-1 sample printed along the thickness direction (L1), (b) SLM-2 sample printed along the length direction (L2), where L2 > L1. Section planes marked by dashed green color represent slices selected for TEM and EBSD tests for individual sample. (For interpretation of the ref erences to color in this figure legend, the reader is referred to the Web version of this article.)
2. Experimental methods Plasma torch atomized Ti–6Al–4V (89.7 w.t.% Ti, 5.8 w.t.% Al, 4.5 2
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3. Results
textured (αþα0 ) phase forming inside two neighboring β grains has different crystallographic orientations. By analyzing the volume frac tions of different phases based on the EBSD patterns in Fig. 4b and e, the contents of β and (αþα0 ) phases are 10% and 90%, respectively, in SLM-1 alloy, while the contents of β and (αþα0 ) phases are 13% and 87%, respectively, in SLM-2 alloy. These values match well with the XRD results in Table 2. The average grain size of SLM-2 is 3.56 μm, while SLM-1 has a relatively smaller grain size of 2.25 μm, as can be derived from Fig. 4c and f. The coarser microstructure in SLM-2 samples resulted from the slower cooling rate during the SLM process. Yang et al. [33] found that the distance from the molten pool to the substrate in one sample affected the microstructure, i.e. grain sizes and phase composi tion of Ti–6Al–4V alloy during the SLM process. Fig. 5a–c shows the bright field TEM images and corresponding selected area electron diffraction (SAED) patterns of SLM-1 alloy. It is seen that the alloy is mainly composed of α0 and α phases. The α phase corresponds to the needles surrounded by dotted red lines, while the α0 phase represents those needles surrounded by dotted blue lines. With the zone axis paralleling to the ½1213�α axis, the SAED pattern in Fig. 5b indicates that the α0 phase in the SLM-1 alloy is composed of (1011)α’ sub-twinning structure. The similar result was observed by Guo et al. [34]. In addition, there is also a small amount of α’’ phase with two different variants in the SLM-1 alloy, as shown by the SAED pattern in Fig. 5c. One represents the horizontally oriented α’’ grains marked by dotted yellow lines in Fig. 5a; the other corresponds to the vertically oriented α’’ grains marked by dotted purple lines. Additionally, the diffraction pattern of residual β phase is also detected. The orientation relationship between the β and α’’ phases is presented by: [100]β//[100]α’’, (011)β//(001)α’’, which conforms to the results re ported by Zhang et al. [35] and Wan et al. [36]. Fig. 5d–f shows the bright field TEM image and corresponding SAED patterns of the SLM-2 alloy, in which the α and α0 phases are also observed. The α0 phase is mainly composed of the {1011}α’ twinning substructure in a zigzag plate-like morphology. The similar results were found by Cao et al. [37] and Liu et al. [38]. Different from the SLM-1 alloy, no α’’ phase is observed in the SLM-2 sample. This may be resulted from the slower cooling rate in the SLM-2 alloy than in the SLM-1 alloy.
3.1. XRD patterns Fig. 2 shows the XRD patterns of the SLM-1 and SLM-2 alloys. It can be seen that both samples consist of β and α/α0 phases. Both the α and α0 phases are in the HCP structure with similar lattice parameters and thus share the same diffraction peaks on the XRD patterns. Besides, the diffraction peak of ð110Þβ planes overlaps with that of ð002Þα’ planes. Therefore, it is difficult to distinguish the three phases according to their peak positions (diffraction angle). However, the content of each phase in SLM-1 and SLM-2 samples can be calculated by the MAUD software based on the intensity differences of β, α and α0 phases, as included in Table 2. It is seen that the SLM-1 and SLM-2 alloys are mainly composed of α and α0 phases as well as a small amount of retained β phase (less than 15%). The difference is that there are more α and β phases but less α0 phase in the SLM-2 alloy than in the SLM-1 alloy. 3.2. Microstructural observation Fig. 3a and b presents the OM micrographs of the SLM-1 sample. Numerous acicular α0 martensite phase forms inside the original columnar β grains, leaving minor residual β phase inside the SLM-1 alloy. The acicular α0 structure has a parallel alignment and seems to share the same crystallographic orientation (marked by the small blue circle) within one original β grain (labeled by dotted green line). Be sides, a small quantity of α phase can be observed, as labeled by the small red circle in Fig. 3b. Fig. 3c and d shows the OM micrographs of the SLM-2 sample. It can be found that the original columnar β grains grow paralleling to the building direction with the length up to milli meters and the averaged width being around 250 μm. On one side, one β grain may branch at the columnar tip. On the other side, two neigh boring β grains may merge into one coarser grain once contacted, as marked in Fig. 3c. It is noted that acicular α0 grains (marked by the blue circle) are in a paralleling alignment, while the α grains (marked by the red ellipse) are in a disordered alignment. In addition, remelting area is observed locally in the SLM-2 alloy, as indicated by the region between the two red lines in Fig. 3d. No α0 phase exists in the remelting areas, because the α0 phase decomposes into αþβ phases during the remelting process. Xu et al. [32] reported that changing SLM process parameters led to the formation of abundant remelting areas composed of fine αþβ lamellae in Ti–6Al–4V alloy. This remelting process is equivalent to a short aging treatment after quenching. Fig. 4a–f shows the EBSD patterns and grain size distribution of SLM1 and SLM-2 alloys. The EBSD patterns in Fig. 4a and d shows that the
3.3. Strength and ductility properties Fig. 6 presents the fractured SLM-1 and SLM-2 samples after tensile failure, from which the strength of both samples are obtained and included in Table 3. The tensile strength is 1236 MPa and 1065 MPa for the SLM-1 and SLM-2 alloys, respectively. By measuring the length change of individual sample after tensile fracture, the elongation values are 8.5% and 13.6% for the SLM-1 and SLM-2 alloys, respectively. The elongation of SLM-2 alloy is 5.1% higher than that of SLM-1 alloy, but the tensile strength is 171 MPa lower. The enhanced elongation of SLM2 sample can be further proved by the obvious necking after tensile test, which was not observed in the SLM-1 sample, as shown in Fig. 6. 4. Discussion The results in present study indicate that building orientation of the SLM process affects the microstructure and phase compositions of Ti–6Al–4V alloy and thus plays an important role in the mechanical Table 2 Phase contents in SLM-1 and SLM-2 alloys. Ti–6Al–4V SLM-1 SLM-2
Fig. 2. XRD patterns of SLM-1 and SLM-2 samples. 3
Volume Fraction (%)
α/α0 ratio
β phase
α phase
α0 phase
9.69 14.23
35.16 53.74
55.15 32.03
0.60 1.68
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Fig. 3. OM micrographs of (a–b) SLM-1 and (c–d) SLM-2 alloys.
Fig. 4. (a) EBSD micrograph, (b) phase analysis and (c) grain size distribution of SLM-1; (d) EBSD micrograph, (e) phase analysis and (f) grain size distribution of SLM-2. The dashed red lines in (a) represent molten pool along the scanning direction, while the dashed blue line in (d) represents grain boundary of original β grains. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
performances of the final bulk material. The studies by de Formanoir et al. [39] indicated that the formation of martensitic α0 phase increased the strength but decreased the plasticity of the Ti–6Al–4V alloy. Comparatively, the existence of α phase is also beneficial to the strength but has negligible effect on the plasticity of the Ti–6Al–4V alloy [40]. Therefore, it can be deduced that a higher α/α0 phase ratio (or a smaller content of α0 phase) will induce a better tensile plasticity for the Ti–6Al–4V alloy. In present study, the α/α0 phase ratio in SLM-2 is much higher than that in SLM-1, as shown in Table 2. Accordingly, the SLM-2 sample has a larger elongation value and thus a better plasticity than the
SLM-1 sample (Table 3). The similar results were also published in previous investigations [21,37,38]. Different α/α0 phase ratios of SLM-1 and SLM-2 samples with varied build orientations result from their different cooling rates during the SLM process. The β→α phase transformation takes place when the cooling rate is slow enough. Otherwise, the β phase may transform to martensitic phase, i.e. α0 phase, in the SLM Ti–6Al–4V alloy at high cooling rates (>106 K s 1) [28]. It is known that the cooling rate is closely correlated to the thermal gradient in the alloy [41,42]. A larger thermal gradient will induce a relatively higher cooling rate in the alloy, 4
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Fig. 5. (a) Bright field TEM image of SLM-1 alloy; (b–c) SAED patterns of SLM-1 alloy (via [1213]α axis and [100]β axis); (d) bright field TEM image of SLM-2 alloy; (e–f) SAED patterns of SLM-2 alloy (via [0111]α’ axis and [0001]α axis).
q¼
λ � gradT ¼
λ
∂T ⇀ n ∂n
(1)
where q is the heat flow per unit area, λ is the thermal conductivity of
the material, n is the normal unit vector, T is the temperature and n is the distance of heat flow. The negative sign represents that the heat flows towards a direction of decreasing temperature. In present study, in spite of a preheating process to the substrate, thermal gradient always exists between the substrate and molten pool due to their temperature dif ference [44]. By incorporating Eq. (1) into the present case, the thermal gradient gradTSLM between the molten pool to the substrate during the SLM process can be described by: ⇀
gradTSLM ¼
Table 3 Ultimate tensile strength (UTS) and elongation of SLM-1 and SLM-2 alloys. UTS (MPa)
Elongation (%)
SLM-1 SLM-2
1236 1065
8.5 13.6
S S
(2)
where ΔTM S is the temperature difference between the molten pool to the substrate and ΔnM S is the normal distance from the molten pool to the substrate. It should be noted that the value of ΔTM S almost remains unchanged with increasing ΔnM S . Therefore, the thermal gradient gradTSLM will decrease as the distance between the substrate and the molten pool increases. In this study, The SLM-1 and SLM-2 samples were built along their thickness (L1) and length (L2) directions, respectively, as illustrated in Fig. 1, where L2 > L1. Thus, it can be deduced the thermal gradient in the SLM-1 sample is larger than that in the SLM-2 sample. Correspondingly, the heat induced from the SLM process dissi pates through the substrate more efficiently in SLM-1 than in SLM-2. In other words, the cooling rate in the SLM-1 sample is higher than that the other one, which leads to a higher percentage of the martensitic α0 phase transforming from the β phase (or smaller α/α0 phase ratio) in the SLM-1 sample. Besides, melting and remelting behavior of the interlayers in the printed alloy also affects the α/α0 phase ratio in the printed alloy. During the SLM process, each interlayer corresponds to one remelting region, which experiences cyclic thermal treatment due to the latent heat
Fig. 6. SLM-1 and 2 samples after fracture.
Sample
ΔTM ΔnM
because less solidification latent heat per unit time is taken away during the cooling process [32]. According to the Fourier’s law, as described in the following Eq. (1), heat energy will flow from the high-temperature area to the lowtemperature area when a temperature gradient gradT (called as ther mal gradient) exists in the material [43]: 5
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(heating, melting and solidification) transferred from the successive deposition layers. Therefore, an alloy with more interlayers will cool down more slowly, because it experiences thermal effects from more successive deposition layers during SLM. According to Xu et al. [18,32], such a remelting behavior during SLM will promote the decomposition of acicular α0 martensite into lamellar (αþβ) structure in the Ti–6Al–4V alloy. In present study, the SLM-2 alloy printed along its length (L2) direction has much more interlayers than the SLM-1 alloy printed along its thickness (L1) direction. The remelting process occurs more frequently in the SLM-2 sample, and thus a smaller amount of α0 martensite is retained in the as-built alloy. Except for the α/α0 phase ratio, the α’’ phase, acting as another martensite other than the α0 phase, can cause great damage to plasticity of titanium alloys [45]. The serial β (bcc)→ α’’ (orthorhombic)→ α’ (hcp) phase transformations occur when the cooling rate is high enough. The orthorhombic (α’’) structure acts as an intermediate state between the hcp-α0 structure and the bcc-β phase. Specific orientation relation ships form between the β, α0 and α’’ phases as follows: [100]β//[100]α’’, (011)β//(001)α’’; [[111]β//[1120]α’, (110)β//(0001)α’ [14,34]. The incomplete lattice reorganization during the β→α0 phase transformation will lead to the generation of α’’ phase. In the current study, a small amount of α’’ phase was found in SLM-1 but not in SLM-2 alloy due to their different cooling rates, as discussed above. The similar result was also found by Donachie [16]. Therefore, the existence of α’’ phase further reduces the elongation of SLM-1. Generally, the α’’ phase is detrimental to the strength of alloys [46]. In present study, however, the content of α’’ phase is too small to influence the tensile strength of the SLM-1 alloy. Additionally, the average grain size of the SLM-2 alloy is larger than that of SLM-1. Therefore, the finer microstructure of the SLM-1 alloy leads to its slightly higher tensile strength than the SLM-2 alloy [47]. In overall, a SLM build orientation paralleling to the length direction contributes to producing Ti–6Al–4V alloy with better comprehensive mechanical performance.
CRediT authorship contribution statement Zongyu Xie: Conceptualization, Methodology, Formal analysis, Investigation, Data curation, Writing - original draft. Yu Dai: Concep tualization, Methodology, Resources. Xiaoqin Ou: Validation, Super vision, Project administration, Funding acquisition, Writing - review & editing. Song Ni: Supervision, Project administration, Funding acqui sition. Min Song: Conceptualization, Validation, Writing - review & editing, Supervision, Project administration, Funding acquisition. Acknowledgements The financial support from National Natural Science Foundation of China (No. 51901248, 51828102) is appreciated. The Advanced Research Center of Central South University is sincerely appreciated for TEM technical support. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.msea.2020.139001. References [1] H. Azizi, H. Zurob, B. Bose, S. Reza Ghiaasiaan, X. Wang, S. Coulson, V. Duz, A. B. Phillion, Additive manufacturing of a novel Ti-Al-V-Fe alloy using selective laser melting, Addit. Manuf. 21 (2018) 529–535, https://doi.org/10.1016/j. addma.2018.04.006. [2] A. Haleem, M. Javaid, 3D printed medical parts with different materials using additive manufacturing, Clin. Epidemiol. Glob. Health. (2019), https://doi.org/ 10.1016/j.cegh.2019.08.002. In press. [3] K. Rajaguru, T. Karthikeyan, V. Vijayan, Additive manufacturing – state of art, Mater. Today Proc. (2019), https://doi.org/10.1016/j.matpr.2019.06.728. In press. [4] K. Wang, R. Bao, T. Zhang, B. Liu, Z. Yang, B. Jiang, Fatigue crack branching in laser melting deposited Ti–55511 alloy, Int. J. Fatig. 124 (2019) 217–226, https:// doi.org/10.1016/j.ijfatigue.2019.03.006. [5] L. Zhou, T. Yuan, R. Li, J. Tang, M. Wang, F. Mei, Microstructure and mechanical properties of selective laser melted biomaterial Ti-13Nb-13Zr compared to hotforging, Mater. Sci. Eng. 725 (2018) 329–340, https://doi.org/10.1016/j. msea.2018.04.001. [6] L. Thijs, F. Verhaeghe, T. Craeghs, J.V. Humbeeck, J.-P. Kruth, A study of the microstructural evolution during selective laser melting of Ti–6Al–4V, Acta Mater. 58 (2010) 3303–3312, https://doi.org/10.1016/j.actamat.2010.02.004. [7] Y.D. Jia, P. Ma, K.G. Prashanth, G. Wang, J. Yi, S. Scudino, F.Y. Cao, J.F. Sun, J. Eckert, Microstructure and thermal expansion behavior of Al-50Si synthesized by selective laser melting, J. Alloys Compd. 699 (2017) 548–553, https://doi.org/ 10.1016/j.jallcom.2016.12.429. [8] A.B. Spierings, K. Dawson, P. Dumitraschkewitz, S. Pogatscher, K. Wegener, Microstructure characterization of SLM-processed Al-Mg-Sc-Zr alloy in the heat treated and HIPed condition, Addit. Manuf. 20 (2018) 173–181, https://doi.org/ 10.1016/j.addma.2017.12.011. [9] N. Kang, M. EL Mansori, A new insight on induced-tribological behaviour of hypereutectic Al-Si alloys manufactured by selective laser melting, Tribol. Int. (2019), https://doi.org/10.1016/j.triboint.2019.04.035. In press. [10] K. Zhang, X. Guo, L. Sun, X. Meng, Y. Xing, Fabrication of coated tool with femtosecond laser pretreatment and its cutting performance in dry machining SLMproduced stainless steel, J. Manuf. Process. 42 (2019) 28–40, https://doi.org/ 10.1016/j.jmapro.2019.04.009. [11] C. Hardes, F. P€ ohl, A. R€ ottger, M. Thiele, W. Theisen, C. Esen, Cavitation erosion resistance of 316L austenitic steel processed by selective laser melting (SLM), Addit. Manuf. 29 (2019), https://doi.org/10.1016/j.addma.2019.100786, 100786. [12] R. Mertens, B. Vrancken, N. Holmstock, Y. Kinds, J.-P. Kruth, J. Van Humbeeck, Influence of powder bed preheating on microstructure and mechanical properties of H13 tool steel SLM parts, Phys. Procedia. 83 (2016) 882–890, https://doi.org/ 10.1016/j.phpro.2016.08.092. [13] G.-A. Tilita, W. Chen, C.K.L. Leung, C.C.F. Kwan, R.L.W. Ma, M.M.F. Yuen, Influence of ultrasonic excitation on the mechanical characteristics of SLM 304L stainless steel, Procedia Eng 216 (2017) 18–27, https://doi.org/10.1016/j. proeng.2018.02.084. [14] R.I. Jaffee, The Science, Technology and Application of Titanium, Pergamon Press, 1970. [15] L. Murr, E.V. Esquivel, S.A. Quinones, S. Gaytan, M.I. Lopez, E.Y. Martinez, F. Medina, D. Hernandez, J.L. Martinez, S. Stafford, D.K. Brown, T. Hoppe, W. Meyers, U. Lindhe, R.B. Wicker, Microstructures and mechanical properties of electron beam-rapid manufactured Ti–6Al–4V biomedical prototypes compared to wrought Ti–6Al–4V, Mater. Char. 60 (2009) 96–105, https://doi.org/10.1016/j. matchar.2008.07.006. [16] M.J. Donachie, Titanium: a Technical Guide, ASM international, 2000.
5. Conclusion The present paper studies the effects of SLM build orientations on the microstructure and tensile performance of Ti–6Al–4V alloys. The SLM alloys are mainly composed of α, α0 and residual β phases. Different build orientations induce different cooling rates during the SLM process. The SLM-2 sample with build orientation along its length direction has smaller cooling rate, more interlayered remelting area and thus higher α/α0 ratio than the SLM-1 sample with build orientation along the thickness direction. In addition, there is also a small amount of α’’ phase, which is an intermediate phase between β and α0 phases, in the SLM-1 alloy due to the high cooling rate during SLM process. The smaller α/α0 phase ratio and existence of α’’ phase induces a slightly higher tensile strength but a much worse plasticity of the SLM-1 alloy. Comparatively, the SLM-2 sample has a slightly lower tensile strength but a much better plasticity due to the higher α/α0 phase ratio and higher fraction of residual β phase. Therefore, the results in present study indicate that SLM build orientation plays an important role in deter mining the microstructure and mechanical properties of alloys. It is suggested that a SLM build orientation along the length direction of Ti–6Al–4V alloy benefits the production of a final material with better comprehensive mechanical performance. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Effects of selective laser melting build orientations on the microstructure and tensile performance of Ti–6Al–4V alloy Zongyu Xie a, Yu Dai b, Xiaoqin Ou a, *, Song Ni a, Min Song a a b
State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, PR China School of Materials Science and Engineering, Nanchang University, Nanchang, 330031, PR China
A R T I C L E I N F O
A B S T R A C T
Keywords: Selective laser melting Ti–6Al–4V alloy Build orientation Phase composition Tensile performance
In this paper, Ti–6Al–4V alloy was fabricated by selective laser melting (SLM) with different build orientations, namely SLM-1 and SLM-2, respectively. Microstructures of SLM-1 and SLM-2 were characterized by X-ray diffraction, optical microscope, scanning electron microscope and transmission electron microscope. The results show that the α and α0 phases in the hexagonal close-packed (hcp) structure, as well as the residual β phase in the body centered cubic (bcc) structure, exist in both samples. However, different SLM build orientations led to varied cooling rates and α/α0 ratios in the final materials. The SLM-1 alloy with build orientation along its thickness direction had higher cooling rate and thus lower α/α0 ratio than the SLM-2 alloy with build orientation along the length direction. Besides, minor martensitic (α’’) phase was observed in the SLM-1 alloy and acted as an intermediate phase between β and α0 phases. The low α/α0 ratio and existence of α’’ phase in the SLM-1 alloy result in a relatively high tensile strength but deteriorated plasticity. Comparatively, the higher α/α0 ratio in the SLM-2 alloy without α’’ martensite leads to a much better plasticity regardless of the slightly reduced tensile strength.
1. Introduction As a near net shape manufacturing technology, additive manufacturing (AM) can save raw materials, improve efficiency and reduce post-processing [1]. Selective laser melting (SLM) is one of the most mature AM technologies due to its unique advantages, such as short lead time and high design freedom [2,3]. The SLM method has been successfully employed in industry to manufacture metallic mate rials, such as Ti alloys [4–6], Al alloys [7–9] and steels [10–13]. Over the last decades, titanium alloys were widely used in aerospace, ship building and biological bone field due to their outstanding properties including high strength-to-weight ratios, good corrosion and heat resistance, low elasticity modules and superior biocompatibility [14]. Nonetheless, the development of titanium alloys manufactured by traditional methods was severely restricted by their high post-processing cost, where the near net shape SLM technology shows special advantages. As the first commercial titanium alloy, Ti–6Al–4V alloy can be a good representative for SLM titanium alloys. Table 1 summarizes the micro structure and tensile properties of the Ti–6Al–4V alloy manufactured by different technologies, such as the select electron beam melting (SEBM)
method [15], solution treated and aged (STA) method [16], mill-annealed (MA) method [16] and SLM method [6,8,17–23]. It is found that the conventional Ti–6Al–4V alloy is composed of the hex agonal close-packed (hcp, α) and body centered cubic (bcc, β) structures [24,25]. However, martensite, such as orthorhombic-α’’ and HCP-α0 phases, may also coexist in the Ti–6Al–4V alloy when it is fabricated by the SLM method. The formation of those two martensitic phases results from the non-diffusional phase transition between the β and α structures. The existence of martensitic structure benefits the tensile strength but lowers the elongation of the SLM Ti–6Al–4V alloy [18]. Accordingly, it is seen from Table 1 that the SLM Ti–6Al–4V alloy shows high tensile strength but relatively low ductility. For example, Gong et al. [19] found that the tensile strength of SLM Ti–6Al–4V alloy reaches 1250 MPa, which is much higher than that of the wrought, STA and MA Ti–6Al–4V alloys. For the SLM Ti–6Al–4V alloy, part of the investigations [7,17,18] focused on the post-processing procedure to improve its combinational properties of strength and ductility, while the others [6,19–21] aimed to improve alloy performance by optimizing key SLM parameters, i.e. laser power (P), layer thickness (τ), laser scanning speed (v), hatch spacing (h) and focal offset distance (FOD). The traditional post-processing method
* Corresponding author. E-mail address: [email protected] (X. Ou). https://doi.org/10.1016/j.msea.2020.139001 Received 14 November 2019; Received in revised form 9 January 2020; Accepted 22 January 2020 Available online 23 January 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.
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Materials Science & Engineering A 776 (2020) 139001
w.t.% V) powder used in this research has an average particle size of 36 μm. The Ti–6Al–4V alloy was prepared using the selective laser melting (SLM) machine (SLM®280 2.0) manufactured by SLM Solution Group AG. Before machine working, the powder bed was preheated to about 200 � C, and the building chamber was filled with argon to ensure the content of oxygen below 80 ppm. The present study chooses the orthogonal band scanning mode. The scanning angle rotates by 37� around the Z-axis for each layer. The following SLM parameters are used: layer thickness (30 mm), laser power (Core: 950 W; Outer Hull: 350 W), laser scanning speed (Core: 230 mm/s; Outer Hull: 1100 mm/s), hatch spacing (0.06 mm), FOD (0 mm). These SLM parameters in present study are selected according to the Material Data Sheet from SLM so lutions Group AG [30], which includes proper combination schemes of key SLM parameters in practice to obtain optimal mechanical properties of Ti–6Al–4V alloys. The schematic picture in Fig. 1a and b shows the two samples with different SLM build directions in present study, namely SLM-1 and SLM-2, respectively. The SLM-1 sample in Fig. 1a was built along the thickness direction (L1), while the SLM-2 sample in Fig. 1b was built along the length direction (L2). Phase compositions in the printed samples were tested by X-ray diffraction (XRD) spectroscopy (Cu Kα radiation, 0.02� step size, 2� per min scanning rate). Phase fractions of the as-build alloys were calculated based on the XRD results using the MAUD software [31]. The as-build samples were machined into tensile specimens with the gauge size of 15 mm � Φ 4 mm. Tensile tests were performed with a constant loading speed of 1 mm/min at room temperature using an Instron-3369 uni versal testing instrument. Microstructures of samples were observed using the optical microscope (OM), FEI Titan G2 60–300 transmission electron microscope (TEM) equipped with an energy disperse spectro scope (EDS), and Nova Nano SEM320 scanning electron microscope (SEM) equipped with an electron backscatter diffraction (EBSD) detec tor. The samples for OM observation were polished and then etched using the Kroll’s reagent. The samples for the TEM and EBSD analysis were prepared as follows: slices with 0.5 mm in thickness were cut from the as-build samples, where the section planes of the slices were parallel to the XY plane for SLM-1 and the XZ plane for SLM-2, respectively, as illustrated in Fig. 1. Then, the slices were mechanically polished to a thickness of 60 μm, followed by twin-jet electropolishing in the elec trolyte composed of 35% normal butanol, 5% perchloric acid and 60% methyl alcohol solution at 25 � C under a voltage of 30 V.
Table 1 Summary of Ti–6Al–4V alloys manufactured by different technologies, where α, α0 and β phases represent the hcp, hcp and bcc structures, respectively. Process
Microstructure
Mill-annealed Solution treated þ aged SLM SLM SLM SLM SLM SLM þ HT SLM SLM þ HIP(920 � C,100 MPa) SLMþ400 � C/2h SLMþ800 C/2h �
SLMþin-situ heat treatment LMD SEBM
Tensile properties
Ref.
UTS (MPa)
E (%)
Globular β in α matrix Fine (αþβ) lamellar
– –
– –
[16] [16]
Fine acicular α0 in columnar original-β grains Fine acicular α0 in columnar original-β grains Very fine (αþβ) lamellar – Fine acicular α0 in columnar original-β grains (αþβ) lamellar and globular
–
–
[6]
1240–1250
6.6–8.8
[19]
1180 1250 1095
11.4 6 8.1
[20] [21] [17]
990
11
[17]
Very fine acicular α0 α lamellas within β matrix
1090 960
10 14
[22] [8]
Fine (αþβ) lamellar in columnar original-β grains (αþβ) lamellar þ globular α at original-β grain boundaries Fine (αþβ) lamellar in columnar original-β grains
1051
13
[18]
1069
10
[18]
1210
11
[18]
–
–
[23]
–
–
[15]
α
Acicular α in columnar priorβ with grain boundary α (αþβ) lamellar
UTS: Ultra tensile strength E: Elongation.
can improve the performance to some extent, which is nonetheless negligible compared to the significant performance promotion resulted from the microstructural optimization by a slight change in SLM pa rameters [26,27]. Xu et al. [18] found that the laser energy density (E ¼ P 0 v⋅h⋅τ) are closely correlated to the formation of the acicular α phase and ultrafine lamellar (αþβ) structure. Specifically, a lower laser energy density produces a higher content of martensitic α0 phase in the SLM Ti–6Al–4V alloy [18]. Besides, it is reported by Zafari et al. [28] that the increase in FOD, which is a variable rarely used by the SLM community, would reduce the energy density during the SLM process. It should be noted that it is inadequate to adjust a single processing variable to control the laser energy density. Instead, it is necessary to design a proper combination of the key SLM processing variables in order to control the microstructure and further the mechanical performance of bulk materials. The SLM process is a rapid solidification process with the cooling rate ranging from 103 to 105 K/s [15], which is far beyond the critical cooling rates for β→α transformation (~20 � C/s) and for β→α’/α’’ martensitic transformation (~410 � C/s) [29]. Therefore, the fast cooling rate during the SLM process will induce martensitic transformation from β to α0 or α’’ phases in the bulk Ti–6Al–4V alloy. In general, the SLM Ti–6Al–4V alloy consists of acicular α0 phase and columnar β phase, whereas it may be composed of fine (αþβ) lamellae occasionally. The α’’ phase is rarely observed in SLM Ti–6Al–4V alloy due to the low content of β-phase stabilizing elements [14]. Except for the key parameters mentioned above, build orientation will also have influence on the microstructure and mechanical perfor mance of SLM titanium alloys. Thus, this paper focuses on the effects of different build orientations on the microstructure and mechanical per formance of SLM Ti–6Al–4V alloy, and the more superior build orien tation of the alloy is determined.
Fig. 1. Schematic illustrating: (a) SLM-1 sample printed along the thickness direction (L1), (b) SLM-2 sample printed along the length direction (L2), where L2 > L1. Section planes marked by dashed green color represent slices selected for TEM and EBSD tests for individual sample. (For interpretation of the ref erences to color in this figure legend, the reader is referred to the Web version of this article.)
2. Experimental methods Plasma torch atomized Ti–6Al–4V (89.7 w.t.% Ti, 5.8 w.t.% Al, 4.5 2
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3. Results
textured (αþα0 ) phase forming inside two neighboring β grains has different crystallographic orientations. By analyzing the volume frac tions of different phases based on the EBSD patterns in Fig. 4b and e, the contents of β and (αþα0 ) phases are 10% and 90%, respectively, in SLM-1 alloy, while the contents of β and (αþα0 ) phases are 13% and 87%, respectively, in SLM-2 alloy. These values match well with the XRD results in Table 2. The average grain size of SLM-2 is 3.56 μm, while SLM-1 has a relatively smaller grain size of 2.25 μm, as can be derived from Fig. 4c and f. The coarser microstructure in SLM-2 samples resulted from the slower cooling rate during the SLM process. Yang et al. [33] found that the distance from the molten pool to the substrate in one sample affected the microstructure, i.e. grain sizes and phase composi tion of Ti–6Al–4V alloy during the SLM process. Fig. 5a–c shows the bright field TEM images and corresponding selected area electron diffraction (SAED) patterns of SLM-1 alloy. It is seen that the alloy is mainly composed of α0 and α phases. The α phase corresponds to the needles surrounded by dotted red lines, while the α0 phase represents those needles surrounded by dotted blue lines. With the zone axis paralleling to the ½1213�α axis, the SAED pattern in Fig. 5b indicates that the α0 phase in the SLM-1 alloy is composed of (1011)α’ sub-twinning structure. The similar result was observed by Guo et al. [34]. In addition, there is also a small amount of α’’ phase with two different variants in the SLM-1 alloy, as shown by the SAED pattern in Fig. 5c. One represents the horizontally oriented α’’ grains marked by dotted yellow lines in Fig. 5a; the other corresponds to the vertically oriented α’’ grains marked by dotted purple lines. Additionally, the diffraction pattern of residual β phase is also detected. The orientation relationship between the β and α’’ phases is presented by: [100]β//[100]α’’, (011)β//(001)α’’, which conforms to the results re ported by Zhang et al. [35] and Wan et al. [36]. Fig. 5d–f shows the bright field TEM image and corresponding SAED patterns of the SLM-2 alloy, in which the α and α0 phases are also observed. The α0 phase is mainly composed of the {1011}α’ twinning substructure in a zigzag plate-like morphology. The similar results were found by Cao et al. [37] and Liu et al. [38]. Different from the SLM-1 alloy, no α’’ phase is observed in the SLM-2 sample. This may be resulted from the slower cooling rate in the SLM-2 alloy than in the SLM-1 alloy.
3.1. XRD patterns Fig. 2 shows the XRD patterns of the SLM-1 and SLM-2 alloys. It can be seen that both samples consist of β and α/α0 phases. Both the α and α0 phases are in the HCP structure with similar lattice parameters and thus share the same diffraction peaks on the XRD patterns. Besides, the diffraction peak of ð110Þβ planes overlaps with that of ð002Þα’ planes. Therefore, it is difficult to distinguish the three phases according to their peak positions (diffraction angle). However, the content of each phase in SLM-1 and SLM-2 samples can be calculated by the MAUD software based on the intensity differences of β, α and α0 phases, as included in Table 2. It is seen that the SLM-1 and SLM-2 alloys are mainly composed of α and α0 phases as well as a small amount of retained β phase (less than 15%). The difference is that there are more α and β phases but less α0 phase in the SLM-2 alloy than in the SLM-1 alloy. 3.2. Microstructural observation Fig. 3a and b presents the OM micrographs of the SLM-1 sample. Numerous acicular α0 martensite phase forms inside the original columnar β grains, leaving minor residual β phase inside the SLM-1 alloy. The acicular α0 structure has a parallel alignment and seems to share the same crystallographic orientation (marked by the small blue circle) within one original β grain (labeled by dotted green line). Be sides, a small quantity of α phase can be observed, as labeled by the small red circle in Fig. 3b. Fig. 3c and d shows the OM micrographs of the SLM-2 sample. It can be found that the original columnar β grains grow paralleling to the building direction with the length up to milli meters and the averaged width being around 250 μm. On one side, one β grain may branch at the columnar tip. On the other side, two neigh boring β grains may merge into one coarser grain once contacted, as marked in Fig. 3c. It is noted that acicular α0 grains (marked by the blue circle) are in a paralleling alignment, while the α grains (marked by the red ellipse) are in a disordered alignment. In addition, remelting area is observed locally in the SLM-2 alloy, as indicated by the region between the two red lines in Fig. 3d. No α0 phase exists in the remelting areas, because the α0 phase decomposes into αþβ phases during the remelting process. Xu et al. [32] reported that changing SLM process parameters led to the formation of abundant remelting areas composed of fine αþβ lamellae in Ti–6Al–4V alloy. This remelting process is equivalent to a short aging treatment after quenching. Fig. 4a–f shows the EBSD patterns and grain size distribution of SLM1 and SLM-2 alloys. The EBSD patterns in Fig. 4a and d shows that the
3.3. Strength and ductility properties Fig. 6 presents the fractured SLM-1 and SLM-2 samples after tensile failure, from which the strength of both samples are obtained and included in Table 3. The tensile strength is 1236 MPa and 1065 MPa for the SLM-1 and SLM-2 alloys, respectively. By measuring the length change of individual sample after tensile fracture, the elongation values are 8.5% and 13.6% for the SLM-1 and SLM-2 alloys, respectively. The elongation of SLM-2 alloy is 5.1% higher than that of SLM-1 alloy, but the tensile strength is 171 MPa lower. The enhanced elongation of SLM2 sample can be further proved by the obvious necking after tensile test, which was not observed in the SLM-1 sample, as shown in Fig. 6. 4. Discussion The results in present study indicate that building orientation of the SLM process affects the microstructure and phase compositions of Ti–6Al–4V alloy and thus plays an important role in the mechanical Table 2 Phase contents in SLM-1 and SLM-2 alloys. Ti–6Al–4V SLM-1 SLM-2
Fig. 2. XRD patterns of SLM-1 and SLM-2 samples. 3
Volume Fraction (%)
α/α0 ratio
β phase
α phase
α0 phase
9.69 14.23
35.16 53.74
55.15 32.03
0.60 1.68
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Fig. 3. OM micrographs of (a–b) SLM-1 and (c–d) SLM-2 alloys.
Fig. 4. (a) EBSD micrograph, (b) phase analysis and (c) grain size distribution of SLM-1; (d) EBSD micrograph, (e) phase analysis and (f) grain size distribution of SLM-2. The dashed red lines in (a) represent molten pool along the scanning direction, while the dashed blue line in (d) represents grain boundary of original β grains. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
performances of the final bulk material. The studies by de Formanoir et al. [39] indicated that the formation of martensitic α0 phase increased the strength but decreased the plasticity of the Ti–6Al–4V alloy. Comparatively, the existence of α phase is also beneficial to the strength but has negligible effect on the plasticity of the Ti–6Al–4V alloy [40]. Therefore, it can be deduced that a higher α/α0 phase ratio (or a smaller content of α0 phase) will induce a better tensile plasticity for the Ti–6Al–4V alloy. In present study, the α/α0 phase ratio in SLM-2 is much higher than that in SLM-1, as shown in Table 2. Accordingly, the SLM-2 sample has a larger elongation value and thus a better plasticity than the
SLM-1 sample (Table 3). The similar results were also published in previous investigations [21,37,38]. Different α/α0 phase ratios of SLM-1 and SLM-2 samples with varied build orientations result from their different cooling rates during the SLM process. The β→α phase transformation takes place when the cooling rate is slow enough. Otherwise, the β phase may transform to martensitic phase, i.e. α0 phase, in the SLM Ti–6Al–4V alloy at high cooling rates (>106 K s 1) [28]. It is known that the cooling rate is closely correlated to the thermal gradient in the alloy [41,42]. A larger thermal gradient will induce a relatively higher cooling rate in the alloy, 4
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Fig. 5. (a) Bright field TEM image of SLM-1 alloy; (b–c) SAED patterns of SLM-1 alloy (via [1213]α axis and [100]β axis); (d) bright field TEM image of SLM-2 alloy; (e–f) SAED patterns of SLM-2 alloy (via [0111]α’ axis and [0001]α axis).
q¼
λ � gradT ¼
λ
∂T ⇀ n ∂n
(1)
where q is the heat flow per unit area, λ is the thermal conductivity of
the material, n is the normal unit vector, T is the temperature and n is the distance of heat flow. The negative sign represents that the heat flows towards a direction of decreasing temperature. In present study, in spite of a preheating process to the substrate, thermal gradient always exists between the substrate and molten pool due to their temperature dif ference [44]. By incorporating Eq. (1) into the present case, the thermal gradient gradTSLM between the molten pool to the substrate during the SLM process can be described by: ⇀
gradTSLM ¼
Table 3 Ultimate tensile strength (UTS) and elongation of SLM-1 and SLM-2 alloys. UTS (MPa)
Elongation (%)
SLM-1 SLM-2
1236 1065
8.5 13.6
S S
(2)
where ΔTM S is the temperature difference between the molten pool to the substrate and ΔnM S is the normal distance from the molten pool to the substrate. It should be noted that the value of ΔTM S almost remains unchanged with increasing ΔnM S . Therefore, the thermal gradient gradTSLM will decrease as the distance between the substrate and the molten pool increases. In this study, The SLM-1 and SLM-2 samples were built along their thickness (L1) and length (L2) directions, respectively, as illustrated in Fig. 1, where L2 > L1. Thus, it can be deduced the thermal gradient in the SLM-1 sample is larger than that in the SLM-2 sample. Correspondingly, the heat induced from the SLM process dissi pates through the substrate more efficiently in SLM-1 than in SLM-2. In other words, the cooling rate in the SLM-1 sample is higher than that the other one, which leads to a higher percentage of the martensitic α0 phase transforming from the β phase (or smaller α/α0 phase ratio) in the SLM-1 sample. Besides, melting and remelting behavior of the interlayers in the printed alloy also affects the α/α0 phase ratio in the printed alloy. During the SLM process, each interlayer corresponds to one remelting region, which experiences cyclic thermal treatment due to the latent heat
Fig. 6. SLM-1 and 2 samples after fracture.
Sample
ΔTM ΔnM
because less solidification latent heat per unit time is taken away during the cooling process [32]. According to the Fourier’s law, as described in the following Eq. (1), heat energy will flow from the high-temperature area to the lowtemperature area when a temperature gradient gradT (called as ther mal gradient) exists in the material [43]: 5
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(heating, melting and solidification) transferred from the successive deposition layers. Therefore, an alloy with more interlayers will cool down more slowly, because it experiences thermal effects from more successive deposition layers during SLM. According to Xu et al. [18,32], such a remelting behavior during SLM will promote the decomposition of acicular α0 martensite into lamellar (αþβ) structure in the Ti–6Al–4V alloy. In present study, the SLM-2 alloy printed along its length (L2) direction has much more interlayers than the SLM-1 alloy printed along its thickness (L1) direction. The remelting process occurs more frequently in the SLM-2 sample, and thus a smaller amount of α0 martensite is retained in the as-built alloy. Except for the α/α0 phase ratio, the α’’ phase, acting as another martensite other than the α0 phase, can cause great damage to plasticity of titanium alloys [45]. The serial β (bcc)→ α’’ (orthorhombic)→ α’ (hcp) phase transformations occur when the cooling rate is high enough. The orthorhombic (α’’) structure acts as an intermediate state between the hcp-α0 structure and the bcc-β phase. Specific orientation relation ships form between the β, α0 and α’’ phases as follows: [100]β//[100]α’’, (011)β//(001)α’’; [[111]β//[1120]α’, (110)β//(0001)α’ [14,34]. The incomplete lattice reorganization during the β→α0 phase transformation will lead to the generation of α’’ phase. In the current study, a small amount of α’’ phase was found in SLM-1 but not in SLM-2 alloy due to their different cooling rates, as discussed above. The similar result was also found by Donachie [16]. Therefore, the existence of α’’ phase further reduces the elongation of SLM-1. Generally, the α’’ phase is detrimental to the strength of alloys [46]. In present study, however, the content of α’’ phase is too small to influence the tensile strength of the SLM-1 alloy. Additionally, the average grain size of the SLM-2 alloy is larger than that of SLM-1. Therefore, the finer microstructure of the SLM-1 alloy leads to its slightly higher tensile strength than the SLM-2 alloy [47]. In overall, a SLM build orientation paralleling to the length direction contributes to producing Ti–6Al–4V alloy with better comprehensive mechanical performance.
CRediT authorship contribution statement Zongyu Xie: Conceptualization, Methodology, Formal analysis, Investigation, Data curation, Writing - original draft. Yu Dai: Concep tualization, Methodology, Resources. Xiaoqin Ou: Validation, Super vision, Project administration, Funding acquisition, Writing - review & editing. Song Ni: Supervision, Project administration, Funding acqui sition. Min Song: Conceptualization, Validation, Writing - review & editing, Supervision, Project administration, Funding acquisition. Acknowledgements The financial support from National Natural Science Foundation of China (No. 51901248, 51828102) is appreciated. The Advanced Research Center of Central South University is sincerely appreciated for TEM technical support. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.msea.2020.139001. References [1] H. Azizi, H. Zurob, B. Bose, S. Reza Ghiaasiaan, X. Wang, S. Coulson, V. Duz, A. B. Phillion, Additive manufacturing of a novel Ti-Al-V-Fe alloy using selective laser melting, Addit. Manuf. 21 (2018) 529–535, https://doi.org/10.1016/j. addma.2018.04.006. [2] A. Haleem, M. Javaid, 3D printed medical parts with different materials using additive manufacturing, Clin. Epidemiol. Glob. Health. (2019), https://doi.org/ 10.1016/j.cegh.2019.08.002. In press. [3] K. Rajaguru, T. Karthikeyan, V. Vijayan, Additive manufacturing – state of art, Mater. Today Proc. (2019), https://doi.org/10.1016/j.matpr.2019.06.728. In press. [4] K. Wang, R. Bao, T. Zhang, B. Liu, Z. Yang, B. Jiang, Fatigue crack branching in laser melting deposited Ti–55511 alloy, Int. J. Fatig. 124 (2019) 217–226, https:// doi.org/10.1016/j.ijfatigue.2019.03.006. [5] L. Zhou, T. Yuan, R. Li, J. Tang, M. Wang, F. Mei, Microstructure and mechanical properties of selective laser melted biomaterial Ti-13Nb-13Zr compared to hotforging, Mater. Sci. Eng. 725 (2018) 329–340, https://doi.org/10.1016/j. msea.2018.04.001. [6] L. Thijs, F. Verhaeghe, T. Craeghs, J.V. Humbeeck, J.-P. Kruth, A study of the microstructural evolution during selective laser melting of Ti–6Al–4V, Acta Mater. 58 (2010) 3303–3312, https://doi.org/10.1016/j.actamat.2010.02.004. [7] Y.D. Jia, P. Ma, K.G. Prashanth, G. Wang, J. Yi, S. Scudino, F.Y. Cao, J.F. Sun, J. Eckert, Microstructure and thermal expansion behavior of Al-50Si synthesized by selective laser melting, J. Alloys Compd. 699 (2017) 548–553, https://doi.org/ 10.1016/j.jallcom.2016.12.429. [8] A.B. Spierings, K. Dawson, P. Dumitraschkewitz, S. Pogatscher, K. Wegener, Microstructure characterization of SLM-processed Al-Mg-Sc-Zr alloy in the heat treated and HIPed condition, Addit. Manuf. 20 (2018) 173–181, https://doi.org/ 10.1016/j.addma.2017.12.011. [9] N. Kang, M. EL Mansori, A new insight on induced-tribological behaviour of hypereutectic Al-Si alloys manufactured by selective laser melting, Tribol. Int. (2019), https://doi.org/10.1016/j.triboint.2019.04.035. In press. [10] K. Zhang, X. Guo, L. Sun, X. Meng, Y. Xing, Fabrication of coated tool with femtosecond laser pretreatment and its cutting performance in dry machining SLMproduced stainless steel, J. Manuf. Process. 42 (2019) 28–40, https://doi.org/ 10.1016/j.jmapro.2019.04.009. [11] C. Hardes, F. P€ ohl, A. R€ ottger, M. Thiele, W. Theisen, C. Esen, Cavitation erosion resistance of 316L austenitic steel processed by selective laser melting (SLM), Addit. Manuf. 29 (2019), https://doi.org/10.1016/j.addma.2019.100786, 100786. [12] R. Mertens, B. Vrancken, N. Holmstock, Y. Kinds, J.-P. Kruth, J. Van Humbeeck, Influence of powder bed preheating on microstructure and mechanical properties of H13 tool steel SLM parts, Phys. Procedia. 83 (2016) 882–890, https://doi.org/ 10.1016/j.phpro.2016.08.092. [13] G.-A. Tilita, W. Chen, C.K.L. Leung, C.C.F. Kwan, R.L.W. Ma, M.M.F. Yuen, Influence of ultrasonic excitation on the mechanical characteristics of SLM 304L stainless steel, Procedia Eng 216 (2017) 18–27, https://doi.org/10.1016/j. proeng.2018.02.084. [14] R.I. Jaffee, The Science, Technology and Application of Titanium, Pergamon Press, 1970. [15] L. Murr, E.V. Esquivel, S.A. Quinones, S. Gaytan, M.I. Lopez, E.Y. Martinez, F. Medina, D. Hernandez, J.L. Martinez, S. Stafford, D.K. Brown, T. Hoppe, W. Meyers, U. Lindhe, R.B. Wicker, Microstructures and mechanical properties of electron beam-rapid manufactured Ti–6Al–4V biomedical prototypes compared to wrought Ti–6Al–4V, Mater. Char. 60 (2009) 96–105, https://doi.org/10.1016/j. matchar.2008.07.006. [16] M.J. Donachie, Titanium: a Technical Guide, ASM international, 2000.
5. Conclusion The present paper studies the effects of SLM build orientations on the microstructure and tensile performance of Ti–6Al–4V alloys. The SLM alloys are mainly composed of α, α0 and residual β phases. Different build orientations induce different cooling rates during the SLM process. The SLM-2 sample with build orientation along its length direction has smaller cooling rate, more interlayered remelting area and thus higher α/α0 ratio than the SLM-1 sample with build orientation along the thickness direction. In addition, there is also a small amount of α’’ phase, which is an intermediate phase between β and α0 phases, in the SLM-1 alloy due to the high cooling rate during SLM process. The smaller α/α0 phase ratio and existence of α’’ phase induces a slightly higher tensile strength but a much worse plasticity of the SLM-1 alloy. Comparatively, the SLM-2 sample has a slightly lower tensile strength but a much better plasticity due to the higher α/α0 phase ratio and higher fraction of residual β phase. Therefore, the results in present study indicate that SLM build orientation plays an important role in deter mining the microstructure and mechanical properties of alloys. It is suggested that a SLM build orientation along the length direction of Ti–6Al–4V alloy benefits the production of a final material with better comprehensive mechanical performance. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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