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\title{Advances in additively manufactured titanium alloys by powder bed fusion and directed energy deposition: Microstructure, defects, and mechanical behavior }

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Review Article

\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-01} \\
 ${ }^{a}$ Centre for Advanced Materials and Manufacturing, School of Engineering, Edith Cowan University, 270 Joondalup Drive, Joondalup, Perth, WA 6027, SA \\
 ${ }^{\mathrm{b}}$ School of Engineering, M050, The University of Western Australia, 35 Stirling Highway, Crawley, Perth, WA 6009, SA \\
 ${ }^{\mathrm{C}}$ Institute of Metals, College of Material Science and Engineering, Changsha University of Science \& Technology, Changsha 410004, China \\
 ${ }^{\mathrm{d}}$ School of Material Science and Engineering, Jiangsu University of Science and Technology, Zhenjiang 212100, China \\
 e State Key Laboratory of Metal Matrix Composites, School of Material Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China\\
\}

\section*{A R T I C L E I N F O}
\section*{Article history:}
Received 5 August 2023

Revised 5 November 2023

Accepted 8 November 2023

Available online 25 November 2023

\section*{Keywords:}
Powder bed fusion

Directed energy deposition

Titanium alloys

Phase transformation

Defects

Mechanical property

\begin{abstract}
A B S T R A C T $\mathrm{Ti}$ and its alloys have been broadly adopted across various industries owing to their outstanding properties, such as high strength-to-weight ratio, excellent fatigue performance, exceptional corrosion resistance and so on. Additive manufacturing (AM) is a complement to, rather than a replacement for, traditional manufacturing processes. It enhances flexibility in fabricating complex components and resolves machining challenges, resulting in reduced lead times for custom designs. However, owing to distinctions among various AM technologies, Ti alloys fabricated by different AM methods usually present differences in microstructure and defects, which can significantly influence the mechanical performance of built parts. Therefore, having an in-depth knowledge of the scientific aspects of fabrication and material properties is crucial to achieving high-performance Ti alloys through different AM methods. This article reviews the mechanical properties of Ti alloys fabricated by two mainstream powder-type AM techniques: powder bed fusion (PBF) and directed energy deposition (DED). The review examines several key aspects, encompassing phase formation, grain size and morphology, and defects, and provides an in-depth analysis of their influence on the mechanical behaviors of Ti alloys. This review can aid researchers and engineers in selecting appropriate PBF or DED methods and optimizing their process parameters to fabricate high-performance Ti alloys for a wide range of industrial applications.
\end{abstract}

(C) 2024 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science \&

Technology.

This is an open access article under the CC BY license (\href{http://creativecommons.org/licenses/by/4.0/}{http://creativecommons.org/licenses/by/4.0/})

\section*{1. Introduction}
Titanium (Ti) and its alloys are widely utilized in various industrial applications due to their exceptional attributes, including excellent fatigue performance, outstanding strength-to-weight ratio, exceptional corrosion resistance, relatively low elastic modulus, and superior biocompatibility [1-8]. Traditional manufacturing techniques, such as casting [9-13], wrought [14], space holder technique [15], powder metallurgy [16-18], foaming [19], and rolling [20] are commonly used for producing Ti alloy parts. However, these methods still require improvement for produc-
\footnotetext{\begin{itemize}
  \item Corresponding authors
\end{itemize}

E-mail addresses: \href{mailto:yjliu@csust.edu.cn}{yjliu@csust.edu.cn} (Y.J. Liu), \href{mailto:lychen@just.edu.cn}{lychen@just.edu.cn} (L.Y. Chen), \href{mailto:wang_liqiang@sjtu.edu.cn}{wang\_liqiang@sjtu.edu.cn} (L.Q. Wang), \href{mailto:l.zhang@ecu.edu.au}{l.zhang@ecu.edu.au}, \href{mailto:lczhangimr@gmail.com}{lczhangimr@gmail.com} (L.C. Zhang).
}

ing Ti alloys due to challenges related to shape complexity, machining difficulties, and susceptibility to oxidation, etc. [21-23]. As such, emerging additive manufacturing (AM) offers a potential supplementary approach to address these challenges in Ti alloy fabrication. AM techniques apply a layer-by-layer strategy to fabricate products from computer-aided design (CAD) models by selectively melting and solidifying raw materials [21,24], offering advantages in Ti alloy production [21,25-29]. Although AM is a well-established advanced manufacturing technology that has attracted significant attention over the past two decades and has been widely applied in various industrial sectors, its industrial adoption has been slower than initially anticipated [28,30,31]. Consequently, it is increasingly regarded as a complementary approach rather than a complete replacement for traditional manufacturing processes. Recently, considerable efforts have concentrated on AM of Ti alloys. In particular, the utilization of multi-laser AM systems,

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-02}
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Fig. 1. Typical types of metal AM techniques (Reproduced with permission from Ref. [47]. Copyright (2022), Elsevier).

ranging from two to eight lasers, has demonstrated the potential to substantially reduce the production time for large-sized $\mathrm{Ti}$ components, such as Ti-6Al-4V [32] and Ti-6.5Al-2Zr-Mo-V alloys [33]. Moreover, multi-material AM has exhibited remarkable versatility in meeting the demands for multifunctional parts. Nevertheless, the disparity in thermal and physical properties has given rise to challenges related to material inhomogeneity and the occurrence of interfacial defects [34-36]. For instance, Wang et al. [37] have utilized the benefits of AM technique by employing elemental mixed powder to investigate $\beta$-type Ti-35Nb alloys. Their study demonstrated a lower cost and greater availability of multipowder feedstock, resulting in a relatively low Young's modulus and moderate corrosion resistance. However, the chemical inhomogeneity in the alloy prepared using powder mixture remains a cause for concern. A recent study of Wang et al. [38] has examined the influence of microstructural and chemical inhomogeneity on the mechanical behaviors of Ti-35Nb alloys prepared using powder mixture and found that the presence of undissolved $\mathrm{Nb}$ induced inhomogeneity, which had a negative impact on the tensile ductility $(3.9 \% \pm 1.1 \%)$, while maintaining a high YS $(648 \pm 13 \mathrm{MPa})$. Using prealloyed Ti-35Nb powder can effectively address the issue of inhomogeneity, due to its stable melting process, leading to improved tensile ductility of $23.5 \% \pm 2.2 \%$, although with a slightly lower YS of $485 \pm 28 \mathrm{MPa}$ [39]. Some researchers have also investigated the corrosion behavior of AM-built Ti alloys. Dai et al. [40] have explored the corrosion resistance of AM-fabricated Ti-6Al-4V alloys in $\mathrm{NaCl}$ solution, where they have reported that AM-fabricated Ti-6Al-4V alloys exhibit poorer corrosion behavior than commercial Ti-6Al-4V (Grade 5) alloys because of the dominant $\alpha^{\prime}$ and less $\beta$ phases. Lu et al. [41] also reported that the AM-fabricated Ti-6Al-4V alloy displayed reduced film corrosion resistance compared to its conventional as-cast counterpart. However, it demonstrated enhanced passivation film stability in dynamic potential polarization curves than the as-cast alloy. Further surface laser remelting significantly improved the surface corrosion resistance of both Ti-6Al-4 V materials. Moreover, Qin et al. [42] found that the corrosion performance of the AM-fabricated $\beta$-type Ti-24Nb-4Zr-8Sn (Ti-2448) alloy was comparable to that of the wrought counterpart. Accordingly, many challenges remain in achieving a well-balanced combination of mechanical and corrosion performance in AM-fabricated Ti alloys.\\
Recently, review papers have highlighted the growing interest in the utilization of Ti alloys for the biomedical industry via AM, such as powder bed fusion (PBF) [43-46]. For instance, Zhang and Chen [44] have conducted a review of $\mathrm{L}$ (laser)-PBF-prepared Ti alloys, which have demonstrated progress in microstructure design, mechanical and corrosion behaviors. Similarly, EB (electron beam)PBF-fabricated Ti alloys with porous and solid structures have also been reviewed, focusing on their microstructure, fatigue behavior and corrosion properties [45]. Moreover, Wang et al. [21] have provided a review of the characteristics of single and multi-melt pool tracks using L-PBF, emphasizing the importance of process parameter optimization and Ti alloy powder qualification. Their review also highlights the significance of precise in-situ measurements and reliable modeling for optimizing parameters. However, there is a notable absence of detailed comparisons of the mechanical behaviors of Ti alloys produced through L-PBF, EB-PBF and directed energy deposition (DED) in the existing literature. As such, this review aims to fill this research gap by thoroughly examining the mechanical behavior of Ti alloys fabricated by two mainstream AM techniques: PBF and DED. By exploring various aspects, such as phase transformation, grain size/morphology and defects, this article provides an in-depth analysis of current advances in the mechanical behaviors of AM-fabricated Ti alloys.

\section*{2. Additive manufacturing}
As shown in Fig. 1, AM techniques are categorized into seven types based on ASTM F2792-12a [48]. Among those AM techniques, sheet lamination (SHL), binder jetting (BJT), directed energy deposition (DED) and powder bed fusion (PBF) are widely employed for manufacturing parts using metallic materials as the feedstock. However, compared with BJT and SHL, PBF and DED are more commonly used for metal fabrication due to two main reasons. Firstly, PBF, DED, and BJT can use metallic powder feedstock to fabricate products, while SHL only uses metal sheets as feedstock. In general, AM techniques using powder feedstock generally result in parts with higher accuracy and surface finish [49], while utilizing wires and metallic sheets as printing feedstock material usually leads to more defects, lower geometry precision, high surface roughness and limitations for the production of complex shapes [50]. Secondly, PBF and DED can directly produce a net-\\
shaped part from a computer model without additional intermediate processing steps to achieve designed shapes in contrast to BJT and SHL [28]. Hence, PBF and DED are more suitable for manufacturing metallic products with high mechanical performance, whereby they have received extensive investigation. Accordingly, this section will provide a further overview of PBF and DED.

\subsection*{2.1. Powder bed fusion}
PBF techniques apply electron or laser beam power to melt a selected area on each thin layer of a pre-deposited powder bed, in order to fabricate products [51-53]. Based on the types of heat sources used, PBF can be categorized into electron beam PBF (EBPBF) and laser PBF (L-PBF). As shown in Fig. 2(a), an L-PBF machine consists of the laser power section, galvanometer-driven mirrors, recoating arm, building platform, shield gas chamber, and other parts [28]. In contrast, a typical EB-PBF system (Fig. 2(b)) is similar to the L-PBF system, except for its special electron beam power section, electromagnetic coils, powder hopper and a vacuum chamber. As illustrated in Fig. 2(a), a 3D computer-aided design (CAD) model needs to be created first in STL format, which will then be sliced into printable 2D layers with a preset layer thickness $[52,54]$. In the EB-PBF process, an extra preheating step is required to lightly sinter the metallic powder feedstock, in order to avoid repulsion and electrostatic charging before the manufacturing process [28]. In contrast, preheating powder feedstock is optional in the L-PBF production process.

The major processing parameters of the EB-PBF and L-PBF techniques include heating source power, scanning rate, layer thickness and hatch spacing [57]. These processing parameters directly affect the energy density, which is of vital guidance and can significantly influence the final performance of products produced by EBPBF and L-PBF. The volume energy density can be determined by Eq. (1) [58]:


\begin{equation*}
E_{\mathrm{V}}=\frac{P}{v \cdot t \cdot s} \tag{1}
\end{equation*}


where $E_{\mathrm{V}}$ is the laser or electron volumetric energy density (J $\left.\mathrm{mm}^{-3}\right), P$ is the input power $(\mathrm{W}), v$ is the scanning rate $\left(\mathrm{mm} \mathrm{s}^{-1}\right)$, $t$ is the layer thickness $(\mathrm{mm})$ and $s$ is the hatching space $(\mathrm{mm})$.

Due to differences in heating sources, there are also some distinctions between the L-PBF and EB-PBF techniques. Firstly, due to the smaller laser spot in size (therefore higher manufacturing accuracy), L-PBF-fabricated products can usually achieve a higher asbuilt surface finish and precision than the products fabricated by EB-PBF [59-62]. Secondly, L-PBF technology can adopt a wider variety of engineering materials as feedstock, while only electrically conductive powder can be used as the feedstock of EB-PBF, due to the heating mode of the electron beam [28,63]. Thirdly, because of the higher input power of the electron beam and energy absorption of metallic powder, EB-PBF has a higher production rate than L-PBF [61].

\subsection*{2.2. Directed energy deposition}
In addition to PBF technology, a great deal of attention has also been paid to the DED technique over the past few years. The DED system generally includes a laser heating source, a multi-axis control system and a feeding system. A typical DED system is shown in Fig. 2(c). DED uses the laser to create an active molten pool, whereby the feeding system utilizes inert gas to transfer metallic powder to the melt pool to layer-by-layer manufacture products from CAD models [28,64]. During the feedstock-adding process, inert gas is also transferred to create a shield gas atmosphere for the molten pool to prevent oxidation influences [65].

Because of differences in manufacturing principles, there are notable distinctions between DED and PBF. Firstly, DED can fabricate larger scaled parts compared to PBF techniques, where there are no limitations on the build chamber size due to the extrusionbased deposition of material [66]. Furthermore, due to its special feeding system, DED can efficiently fabricate functionally graded materials by adjusting feedstock material during the production process [67]. In addition, multi-axis nozzles used in the DED technique offer greater flexibility for material deposition from various angles, where they have become an effective tool for repairing damaged parts [65]. However, unlike the PBF system, DED lacks recoaters to level the powder feedstock before manufacturing each layer. This can lead to the accumulation of build errors in each layer, resulting in lower fabrication accuracy of DED-fabricated products compared to those produced through PBF [66,68]. A comparison of common information for these $3 \mathrm{AM}$ techniques is shown in Table 1.

\section*{3. Phase transformation in AM-fabricated Ti alloys}
\subsection*{3.1. Characterizing phase transformation}
The crystal structure of pure Ti remains hexagonal close-packed (hcp, $\alpha$ phase) at low temperatures, but transitions to the bodycentered cubic (bcc, $\beta$ phase) structure when the temperature exceeds the allotropic transition temperature $\left(882.5^{\circ} \mathrm{C}\right)$ [76]. It has been shown that the addition of various alloying elements can significantly affect the phase transformation temperature of Ti [77]. $\alpha$-stabilizers, such as $\mathrm{Al}, \mathrm{C}$, and $\mathrm{O}$, usually can increase the $\beta / \alpha$ transit temperature, while $\beta$-stabilizers, such as Ta, Mo, and $\mathrm{Nb}$, can decrease the $\beta / \alpha$ transit temperature [78]. Ti alloys can generally be classified into three types based on the remaining phases, namely $\alpha$-type, $(\alpha+\beta)$-type and $\beta$-type Ti alloys $[21,79]$.

$\alpha$-type Ti alloys consist of a single solid solution of $\alpha$ phase and usually contain various grades of commercially pure Ti (CPTi) alloys [80]. $\alpha$-type Ti alloys usually exhibit good weldability, excellent creep resistance, and outstanding corrosion resistance, making them suitable for applications in chemical engineering and high-temperature fields $[81,82]$. However, because of the brittleness and stability of the hcp structure, $\alpha$-type Ti alloys usually present low strength at room temperature, and commonly cannot be enhanced by microstructural modifications from heat treatment $[81,83]$. In contrast, $(\alpha+\beta)$-type Ti alloys contain more $\beta$ stabilizers $(\sim 4 \%-16 \%)$ and can maintain about $5-30$ vol.\% of $\beta$ phase $[78,79]$. This leads to $(\alpha+\beta)$-type Ti alloys exhibiting excellent corrosion resistance and superior strength at room temperature [84-86]. The Ti-6Al-4V alloy is one of the most widely used $\alpha+\beta$ dual-phase Ti alloys, with over half of Ti alloy usage being attributed to this alloy. As a result, it has been extensively applied in biomedical and energy industries [80,87]. Nevertheless, despite its popularity, the alloy's $\mathrm{Al}$ and $\mathrm{V}$ elements can have negative impacts on human health when used as a medical implant. Excessive $\mathrm{V}$ intake can cause dehydration, diarrhea and reduced weight gain [88], while an excess of Al can lead to Alzheimer's disease [89]. In addition, Young's modulus of Ti-6Al-4V ( 110-120 GPa) is still significantly higher than that of human bones, where the different Young's modulus between human bone and implant can induce the stress-shielding effect [90]. Therefore, novel $\beta$-Ti alloys consisting of nontoxic compositions have been considered as alternative materials of Ti-6Al-4V for medical implants. $\beta$-type Ti alloys contain the highest proportion of $\beta$-stabilizers in Ti alloys, and mainly include the $\beta$ phase at room temperature [91]. Due to the higher proportion of $\beta$ phase in $\beta$-type Ti alloys, properties such as higher toughness, plasticity and heat treatment capability can be achieved, along with a lower elastic modulus [81]. Thus, the stress shield effect can effectively be decreased by using $\beta$-type Ti alloys. Moreover, $\beta$-type Ti alloys also present better corrosion resistance, because the micro-galvanic effect between various phases can be decreased in $\beta$-type Ti alloys [92].\\
(a) 3D designed model

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Computer-aided design software

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Scan direction\\
Sliced 2D layer

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Laser Powder Bed Fusion\\
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\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-04(4)}

Fig. 2. Schematic diagram of additive manufacturing: (a) L-PBF (Reproduced with permission from Ref. [21]. Copyright (2023), Elsevier), (b) EB-PBF systems (Reproduced with permission from Ref. [55]. Copyright (2021), Elsevier), and (c) DED powder feed system (Reproduced with permission from Ref. [56]. Copyright (2014), Springer Nature).

Table 1

\begin{center}
\begin{tabular}{llll}
Comparison of L-PBF, EB-PBF and DED techniques. &  &  &  \\
\hline
AM techniques & L-PBF & EB-PBF & DED (powder based) \\
\hline
Heating source & Laser beam & Electron beam & Laser beam \\
Typical Power $(\mathrm{W})$ & $\sim 100-1000$ & $\sim 3500$ & $\sim 500$ \\
Beam spot size $(\mu \mathrm{m})$ & $40-100$ & $100-200$ & $380-900$ \\
Scan speed $\left(\mathrm{mm} \mathrm{s}^{-1}\right)$ & $\sim 100-2000$ & $>1000$ & $3-15$ \\
Typical layer thickness $(\mu \mathrm{m})$ & $10-50$ & 50 & $100-250$ \\
Typical cooling rate $\left(\mathrm{K} \mathrm{s}^{-1}\right)$ & $10^{3}-10^{8}$ & $10^{3}-10^{5}$ & $10^{4}-10^{6}$ \\
Build plate temperature $\left({ }^{\circ} \mathrm{C}\right)$ & $\sim 80-200$ & $\sim 500-750$ & - \\
Building environment & Argon/Nitrogen & Vacuum & Argon \\
Preheating of powder $\left({ }^{\circ} \mathrm{C}\right)$ & - & $\sim 600-750$ & - \\
Refs. & $[69-72]$ & $[69,72,73]$ & $[70,74,75]$ \\
\end{tabular}
\end{center}

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Fig. 3. Microstructure of L-PBF-produced CP-Ti alloys observed by (a) optical microscopy (OM) (Reproduced with permission from Ref. [94]. Copyright (2017), Elsevier) and (b) scanning electron microscopy (SEM) (Reproduced with permission from Ref. [95]. Copyright (2016), Elsevier), (c) transmission electron microscopy (TEM) micrograph of CP-Ti fabricated by EB-PBF (Reproduced with permission from Ref. [96]. Copyright (2015), Elsevier), and (d) OM micrograph of DED-fabricated CP-Ti alloys (Reproduced with permission from Ref. [94]. Copyright (2017), Elsevier).

As mentioned above, the presence of different phases plays a critical role in determining the mechanical properties of Ti alloys [37]. In addition to alloy composition, process conditions during fabrication, such as cooling rate and build plate temperature, also have a significant impact on the phase formation and transformation of Ti alloys [21]. As a result, the phase formation of Ti alloys produced by different AM techniques can vary due to differences in the production process. In particular, the amount of nonequilibrium martensite phases is distinct in Ti alloys fabricated by different AM techniques. The formation of the martensite usually requires two essential conditions: (i) the build plate temperature is lower than the martensite start temperature; (ii) an appropriate low cooling temperature and high cooling rate must be achieved during the AM process. The required cooling rate and martensite start temperature vary depending on the chemical composition of the Ti alloy. For CP-Ti alloys, the martensite start temperature is around $850{ }^{\circ} \mathrm{C}$, where the required cooling rate should be more than 300-1000 $\mathrm{K} \mathrm{s}^{-1}$ for martensite formation [93]. As shown in Table 1, the cooling rate in all $3 \mathrm{AM}$ techniques satisfies the formation conditions of martensite in the CP-Ti alloy. However, the final phase constituents vary in CP-Ti alloys fabricated using these 3 AM techniques. As shown in Fig. 3(a), Attar et al. [94] have reported that L-PBF-fabricated CP-Ti alloy contains acicular and lathtype martensitic $\alpha^{\prime}$ phase. Sing et al. [95] have also shown that\\
L-PBF-fabricated CP-Ti alloys exhibit a microstructure consisting of mixed platelet $\alpha$ and acicular $\alpha^{\prime}$ phases, as shown in Fig. 3(b). By contrast, as shown in Fig. 3(c), only $\alpha$-lath phase is formed in CP-Ti alloy fabricated by EB-PBF [96]. This is because the manufacturing process of EB-PBF usually takes several hours at high build plate temperatures, which is almost the same as the annealing/aging process [62]. This type of phenomenon in the EB-PBF process facilitates the $\alpha^{\prime}$ phase to transform into the $\alpha$ phase [96]. Fig. 3(d) shows the DED-fabricated CP-Ti alloy, which also only contains the plate-like $\alpha$ phase [94,97]. DED requires a higher energy density than that of L-PBF in order to achieve high-density fabrication, resulting in a higher penetration (melt pool) depth during the producing process of DED [94]. Thus, compared to L-PBF, DED involves more repeated heat cycles in the melt layers, resulting in the transfer of more heat to the deposited layers, and causing the $\alpha^{\prime}$ phase to dissolve into the $\alpha$ phase.

Compared with the CP-Ti alloy, the martensite start temperature of the Ti-6Al-4V is around $575^{\circ} \mathrm{C}$ [98]. Additionally, it should be noted here that for the Ti-6Al- $\mathrm{V}$, the formation of $\alpha^{\prime}$ martensite is also dependent on the cooling rate during the fabrication process. More specifically, cooling rates exceeding $683 \mathrm{~K} \mathrm{~s}^{-1}$ would result in complete $\alpha^{\prime}$ martensite formation, while cooling rates ranging from $683 \mathrm{~K} \mathrm{~s}^{-1}$ to $293 \mathrm{~K} \mathrm{~s}^{-1}$ result in incomplete $\alpha^{\prime}$ martensite formation. Cooling rates lower than $293 \mathrm{~K} \mathrm{~s}^{-1}$ would not result\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-06}

Fig. 4. Microstructures of Ti-6Al-4V alloys fabricated by (a) L-PBF (Reproduced with permission from Ref. [113]. Copyright (2017), Elsevier), (b) EB-PBF (Reproduced with permission from Ref. [114]. Copyright (2015), Springer Nature), and (c, d) DED techniques (Reproduced with permission from Ref. [112]. Copyright (2019), John Wiley and Sons).

in the formation of $\alpha^{\prime}$ martensite [98]. Although the cooling rate in the 3 AM techniques met the above conditions for the complete martensite formation of the Ti-6Al-4V alloy, the results were not as expected. As seen from Fig. 4(a), due to the relatively highest cooling rate, the Ti-6Al-4V alloy fabricated by L-PBF is mainly composed of fine $\alpha^{\prime}$ phase [72,95,99-103]. Moreover, in addition to the $\alpha^{\prime}$ phase, Ren et al. [104] have reported 0.1 vol.\% $\beta$ phase in L-PBF-fabricated Ti-6Al-4V samples. By contrast, as shown in Fig. 4(b), the EB-PBF-fabricated Ti-6Al-4V alloy contained dominant $\alpha$ phase and trace $\beta$ phase, where the $\alpha^{\prime}$ phase has not been noted in EB-PBF-fabricated Ti-6Al-4V samples in the literature, e.g. Refs. [72,99,101,105,106]. Neikter et al. [107] have reported on the basket-weave microstructure of mixed $(\alpha+\beta)$ phases in EBPBF-fabricated Ti-6Al-4V samples. Further works have also shown mixed $\alpha+\alpha^{\prime}$ phases occurring in the top surface of the EB-PBFfabricated Ti-6Al-4V sample, due to this area having a higher cooling rate [108]. The distinction in the phase constituent in Ti-6Al$4 \mathrm{~V}$ alloys fabricated by L-PBF and EB-PBF technologies can be attributed to different building conditions in L-PBF and EB-PBF techniques. As shown in Table 1, the elevated build plate temperature in the EB-PBF fabrication chamber, typically exceeding the martensite start temperature of the Ti-6Al- $4 \mathrm{~V}$ alloy, encourages the formation of $\alpha+\beta$ phases instead of the martensitic $\alpha^{\prime}$ phase transformation [109], and can even trigger the transformation of $\alpha^{\prime}$ phase into $\alpha$ phases, especially during prolonged printing [110].

For the DED-fabricated Ti-6Al-4V samples, the presence of the dominant fine needle-shaped $\alpha^{\prime}$ phase microstructure has been observed by Hao et al. [111]. In contrast to this, Razavi et al. [112] have reported on the coexistence of both $\alpha^{\prime}$ phase and $\alpha+\beta$ phases in Ti-6Al-4V samples fabricated by the DED method, as shown in Fig. 4(c) and (d). In addition, Zhai et al. [105] have indicated that the Ti-6Al-4V alloy fabricated by DED with low power\\
( $330 \mathrm{~W})$ mainly contained the $\alpha^{\prime}$ phase, while a mixed microstructure of $\alpha^{\prime}$ and $\alpha+\beta$ phases occurred in the Ti-6Al-4V alloy fabricated with high power DED ( $780 \mathrm{~W})$. Moreover, Rashid et al. [106] also noted that DED-fabricated Ti-6Al-4V alloys mainly consisted of $\alpha$ and $\beta$ phases. As indicated in Table 1, DED exhibits an intermediate cooling rate, resulting in higher $\alpha^{\prime}$ martensite content in Ti-6Al-4V alloys compared to EB-PBF and lower content compared to L-PBF. Moreover, substantial thermal accumulation during DED fabrication can induce partial $\alpha^{\prime}$ phase decomposition.

Due to the properties of nontoxic elements and lower elastic modulus, $\beta$-type Ti alloys have been considered as potential alternative materials for medical implants in place of Ti-6Al-4V alloys. The Ti-2448 [62], Ti-25Nb-3Zr-3Mo-2Sn (TLM) [115], Ti-35Nb-7Zr5Ta [116], Ti-35Nb-2Ta-3Zr [117-119], and Ti-35Nb [39] alloys are examples of typical $\beta$-type Ti alloys that are known for their low elastic modulus, whereby the formation of distinct phases within these alloys is also influenced by the specific AM techniques employed. For instance, it has been reported that Ti-2448 alloy fabricated using L-PBF exhibits a single $\beta$ phase $[42,62,120]$. In addition, Hernandez et al. [121] have shown that the Ti-2448 alloy fabricated by EB-PBF using a powder preheating of $250{ }^{\circ} \mathrm{C}$ contained both $\beta$ and $\alpha^{\prime \prime}$ phases, as shown in Fig. 5(a). Furthermore, Fig. 5(b) shows that the EB-PBF-fabricated Ti-2448 alloys exhibited both $\alpha$ and $\beta$ phases when the build plate heating temperature was set at $500{ }^{\circ} \mathrm{C}[62,122]$.

With regard to the Ti-35Nb-7Zr-5Ta alloy, some studies have reported that the L-PBF-fabricated Ti-35Nb-7Zr-5Ta alloy only exhibits a single $\beta$ phase in the X-ray diffraction (XRD) pattern [123125]. However, as shown in Fig. 5(c) and (d), some studies have also found that the $\omega$ phase occurred in the lamellar mechanical twins in the L-PBF-fabricated Ti-35Nb-7Zr-5Ta alloy, as evident via further observation of the corresponding selected area electron\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-07}

Fig. 5. Microstructures of Ti-2448 alloy produced by EB-PBF with (a) powder preheating of $250{ }^{\circ} \mathrm{C}$ (Reproduced with permission from Ref. [121]. Copyright (2013), Elsevier) and (b) build plate temperature of $500^{\circ} \mathrm{C}$ (Reproduced with permission from Ref. [122]. Copyright (2016), Elsevier); (c) lamellar mechanical twins and (d) its SAED patterns of the L-PBF-built Ti-35Nb-7Zr-5Ta alloy (Reproduced with permission from Ref. [116]. Copyright (2022), Elsevier); (e) bright field TEM image and (f) relevant SAED patterns in the DED-produced Ti-35Nb-7Zr-5Ta alloy (Reproduced with permission from Ref. [127]. Copyright (2006), John Wiley and Sons).

diffraction (SAED) pattern and transmission electron microscopy (TEM) [116,126]. Compared with the Ti-2448 alloy, the Ti-35Nb7Zr-5Ta alloy has a higher amount of $\beta$ stabilizers (Nb), which hinders the $\alpha^{\prime \prime}$ phase formation due to the high percentage of $\beta$ stabilizers. Accordingly, only the $\omega$ phase formed in L-PBF-fabricated Ti-35Nb-7Zr-5Ta alloys. Similarly, Wang et al. [39] have reported that Ti-35Nb alloys fabricated using L-PBF with prealloyed powder primarily exhibit metastable $\beta$ phase, $\alpha^{\prime \prime}$ phase and trace nano $\omega$ phase, while the alloys fabricated with mixed powder showed a combination of $\beta+\alpha$ phases and trace amounts of $\alpha^{\prime \prime}$ phase. This difference in microstructure can be attributed to the presence of many undissolved $\mathrm{Nb}$ particles in the mixed powder sample, resulting in their distinct microstructure during the fabrication process. In addition, as shown in Fig. 5(e) and (f), Banerjee et al. [127] have claimed that the DED-fabricated Ti-35Nb-7Zr-5Ta alloy contains the $\beta$ phase and nanometer-scale $\omega+\alpha$ phases. Similarly, Nartu et al. [128] have investigated the Ti-35Nb-7Zr-5Ta alloy fabricated by DED with different laser powers, finding that only a single $\beta$ phase occurred in the XRD pattern of Ti-35Nb-7Zr-5Ta alloy fabricated by DED under $400 \mathrm{~W}$ laser power, while both $\alpha$ and $\beta$ phases presented in the XRD pattern of DED-fabricated Ti-35Nb7Zr-5Ta alloy with 500 and $600 \mathrm{~W}$ laser power [128]. According to Gu et al. [129], linear energy density (LED), which is calculated as the laser power $(P)$ divided by the scanning speed $(v)$, is a critical factor in determining heat input during the melting process. In a study conducted by Nartu et al. [128], the DED process with different laser powers of 400, 500 and $600 \mathrm{~W}$ resulted in LED values of $31.5,39.4$ and $47.2 \mathrm{~J} / \mathrm{mm}$ respectively, which were much higher than the LED of L-PBF $(0.16 \mathrm{~J} / \mathrm{mm})$ reported in the study by Luo et al. [116]. This suggests that there is a larger heat accumulation during the process of DED. Moreover, the higher heat accumulation of DED at 500 and $600 \mathrm{~W}$ laser powers, compared to DED at $400 \mathrm{~W}$, prolongs the cooling time of each deposited layer, providing favorable conditions for the diffusion transition from the $\beta$ phase to the $\alpha$ phase.

\subsection*{3.2. Influence of phase transformation on mechanical properties}
Table 2 shows the mechanical properties of Ti alloys fabricated by 3 AM techniques. L-PBF-fabricated Ti alloys typically exhibit relatively higher hardness and strength, but lower ductility compared to Ti alloys fabricated by EB-PBF and DED. This can be ascribed to the formation of more non-equilibrium phases in L-PBF-fabricated Ti alloys. According to Attar et al. [130], compared with $\alpha$ phase, there is a higher density of dislocations, increased amount of closely spaced interfaces and separation of adjacent laths and plates in the martensite $\alpha^{\prime}$ phase. In addition, the martensitic $\alpha^{\prime}$ phase also can result in the lattice structure deformation in Ti alloys [59]. Therefore, the dislocation slip can be effectively hindered by the martensitic structure during deformation,

Table 2

Mechanical and tensile properties of various Ti alloys fabricated by different AM techniques.

\begin{center}
\begin{tabular}{|c|c|c|c|c|c|c|c|}
\hline
Material & Method & $\sigma_{0.2}(\mathrm{MPa})$ & $\sigma_{\text {UTS }}(\mathrm{MPa})$ & $H_{\mathrm{v}}$ & $E(\mathrm{GPa})$ & $\varepsilon(\%)$ & Refs. \\
\hline
CP-Ti & L-PBF & $555 \pm 3$ & $757 \pm 12.5$ & 261 & - & $19.5 \pm 1.8$ & [130] \\
\hline
CP-Ti & L-PBF & $620 \pm 20$ & $703 \pm 16$ & $213 \pm 10$ & $112 \pm 3$ & $5.2 \pm 0.3$ & [95] \\
\hline
CP-Ti & EB-PBF & $377 \pm 10$ & $475 \pm 15$ & - & - & $28.5 \pm 0.5$ & [96] \\
\hline
CP-Ti & DED & $518 \pm 5$ & $640 \pm 6$ & 241 & $126 \pm 3$ & 29 & [97] \\
\hline
Ti-6Al-4V & L-PBF & $1056 \pm 64$ & $1166 \pm 107$ & $383 \pm 11$ & $132 \pm 16$ & $6.1 \pm 2.6$ & [95] \\
\hline
Ti-6Al-4V & L-PBF & 1036 & 1086 & - & - & 10.4 & [104] \\
\hline
Ti-6Al-4V & L-PBF & 1110 & 1267 & 409 & 109 & 7.3 & $[100]$ \\
\hline
Ti-6Al-4V & L-PBF & $1143 \pm 30$ & $1219 \pm 20$ & - & - & $4.9 \pm 0.6$ & [101] \\
\hline
Ti-6Al-4V & L-PBF & $1150 \pm 67$ & $1246 \pm 134$ & - & - & $1.4 \pm 0.5$ & $[72]$ \\
\hline
Ti-6Al-4V & L-PBF & $962 \pm 47$ & $1166 \pm 25$ & - & - & $1.7 \pm 0.3$ & [102] \\
\hline
Ti-6Al-4V & L-PBF & 736 & 1051 & 360 & 110 & 11.9 & [103] \\
\hline
Ti-6Al-4V & L-PBF & $960 \pm 12$ & $1042 \pm 16$ & - & - & $9.3 \pm 1.04$ & [134] \\
\hline
Ti-6Al-4V & L-PBF & 1234 & 1286 & - & - & 5.22 & [135] \\
\hline
Ti-6Al-4V & EB-PBF & 645 & 778 & - & - & 12.2 & [104] \\
\hline
Ti-6Al-4V & EB-PBF & $869 \pm 7.2$ & $928 \pm 9.8$ & 311 & - & $9.9 \pm 1.7$ & [101] \\
\hline
Ti-6Al-4V & EB-PBF & 1001 & 1073 & - & - & 11 & [105] \\
\hline
Ti-6Al-4V & EB-PBF & $890 \pm 15$ & $930 \pm 22$ & 347 & 125 & 11 & [106] \\
\hline
Ti-6Al-4V & EB-PBF & $846 \pm 7$ & $976 \pm 11$ & - & - & $15 \pm 2$ & $[72]$ \\
\hline
Ti-6Al-4V & EB-PBF & $830 \pm 5$ & $910 \pm 10$ & 328 & $118 \pm 5$ & - & [136] \\
\hline
Ti-6Al-4V & EB-PBF & 735 & 775 & 377 & 93 & 2.3 & [137] \\
\hline
Ti-6Al-4V & EB-PBF & $13.2-16.3$ & $944.5-964.5$ & 330 & - & 14.6 & [138] \\
\hline
Ti-6Al-4V & DED (low power) & 1005 & 1103 & - & - & 4 & [105] \\
\hline
Ti-6Al-4V & DED (high power) & 990 & 1042 & - & - & 7 & [105] \\
\hline
Ti-6Al-4V & DED & $1020 \pm 13$ & $1100 \pm 9$ & 395 & 125 & 8 & [106] \\
\hline
Ti-6Al-4V & DED & $976 \pm 24$ & $1099 \pm 2$ & $360 \pm 10$ & - & $4.9 \pm 0.1$ & [139] \\
\hline
Ti-6Al-4V & DED & $916 \pm 26$ & $1032 \pm 31$ & - & $113 \pm 5$ & $19 \pm 4$ & [140] \\
\hline
Ti-6Al-4V & DED & $950 \pm 2$ & $1025 \pm 2$ & - & - & $5 \pm 1$ & [141] \\
\hline
Ti-6Al-4V & DED & $945 \pm 13$ & $1041 \pm 12$ & - & - & $14.5 \pm 1.2$ & [142] \\
\hline
Ti-2448 & L-PBF & $563 \pm 38$ & $665 \pm 18$ & $220 \pm 6$ & $53 \pm 1$ & $13.8 \pm 4.1$ & [120] \\
\hline
Ti-2448 & L-PBF & $490 \pm 16$ & $700 \pm 6$ & $219 \pm 8$ & $49 \pm 1$ & $22 \pm 1$ & [143] \\
\hline
Ti-2448 & EB-PBF & - & - & 255.1 & - & - & [121] \\
\hline
Ti-25Nb-3Zr-3Mo-2Sn & L-PBF & $592 \pm 21$ & $716 \pm 14$ & - & - & $37 \pm 5$ & [115] \\
\hline
Ti-35Nb-7Zr-5Ta & L-PBF & $569.5-668.5$ & $575.1-675.0$ & - & $76.7-85.26$ & $26.2-31.7$ & [125] \\
\hline
Ti-35Nb-7Zr-5Ta & L-PBF & $816 \pm 26$ & $830 \pm 26$ & - & $66.5 \pm 1.5$ & $16.5 \pm 1.8$ & [116] \\
\hline
Ti-35Nb-7Zr-5Ta & L-PBF & 309 & 631 & $205 \pm 10$ & 81 & $14-15$ & $[124]$ \\
\hline
Ti-35Nb-7Zr-5Ta & DED & 813.6 & 834.2 & - & 55 & 19 & [127] \\
\hline
Ti-35Nb-2Ta-3Zr & L-PBF & $\sim 430$ & 552 & - & - & 21 & [117] \\
\hline
$\mathrm{Ti}-35 \mathrm{Nb}$ & L-PBF & $648 \pm 13$ & $803 \pm 33$ & $274 \pm 7$ & $84 \pm 2$ & $3.9 \pm 1.1$ & $[38]$ \\
\hline
$\mathrm{Ti}-35 \mathrm{Nb}$ & L-PBF & $636 \pm 80$ & $750 \pm 51$ & $241 \pm 14$ & $85 \pm 2$ & $2.2 \pm 1.4$ & $[39]$ \\
\hline
$\mathrm{Ti}-35 \mathrm{Nb}$ & L-PBF & $485 \pm 28$ & $645 \pm 9$ & $174 \pm 7$ & $72 \pm 1$ & $23.5 \pm 2.2$ & $[39]$ \\
\hline
\end{tabular}
\end{center}

$\sigma_{0.2}$ - tensile yield strength; UTS - ultimate tensile strength; $H_{\mathrm{V}}$ - Vickers hardness; $E$ - elastic modulus and $\varepsilon$ - tensile elongation.

which results in dislocation strengthening and improved strength of Ti alloys [131-133]. However, the delayed dislocation movement can negatively impact the ductility of the material. Zhao et al. [99] have reported that the martensitic $\alpha^{\prime}$ phase has high strength and low ductility, which is associated with Ti alloys containing the $\alpha^{\prime}$ phase. Therefore, it is reasonable that L-PBF-fabricated Ti alloys present relatively higher strength and hardness, but lower ductility in comparison to those produced by EB-PBF and DED methods.

\section*{4. Grain size and morphology in AM-fabricated Ti alloys}
\subsection*{4.1. Understanding grain size}
In addition to variations in phase distinctions, the grain size in Ti alloys fabricated by different AM techniques also varies. In general, grain size is mainly influenced by growth time (i.e., solidification rate) [21], where longer growth time leads to the formation of coarser grain, while shorter growth time typically results in fine grains. However, accurately predicting the growth time of grains in Ti alloys is a complex task due to several reasons. Firstly, different types of grains in Ti alloys nucleate and grow within different temperature ranges. For example, $\beta$ grains usually form when the temperature is below the melting point and above the allotropic transition temperature, so the growth time of $\beta$ grains in Ti alloys depends on the dwell time between the allotropic transition temperature and the melting point [144]. In contrast, $\alpha$ grains form when the temperature is below the allotropic transition tempera- ture, whereby the growth time of $\alpha$ grains is mainly impacted by the cooling rate [107]. Moreover, the final grain size in Ti alloys varies due to differences in the manufacturing processes employed by different AM techniques. Currently, the literature on Ti-6Al-4V alloy fabricated by AM techniques is most abundant, which provides adequate evidence for comparing the grain sizes of Ti alloys. Accordingly, the following sections will present the Ti-6Al-4V alloy as an example to discuss size distinctions in grain sizes among $\mathrm{Ti}$ alloys fabricated by different AM techniques.

\subsection*{4.1.1. Prior $\beta$ grain size}
As discussed above, when the temperature drops below the allotropic transition temperature, $\beta$ grains will partially decompose into $\alpha / \alpha^{\prime}$ grains in Ti-6Al-4V alloys, based on different producing conditions. Thus, the $\beta$ grain is commonly referred to as the "prior $\beta$ grain". Usually, the Ti-6Al-4V alloy fabricated by the DED process tends to exhibit larger prior $\beta$ grains compared to those produced by L-PBF and EB-PBF techniques. Fig. 6(a) presents electron backscatter diffraction (EBSD) images of the horizontal layer (scan plane) in the L-PBF-fabricated Ti-6Al-4V alloy, showcasing a considerably homogeneous microstructure characterized by a basketweave and needle-like morphology of $\alpha / \alpha^{\prime}$ phases, with an average width of less than $10 \mu \mathrm{m}$. These phases undergo a phase transformation from the prior $\beta$ grains. The corresponding reconstructed EBSD microstructure of the prior $\beta$ grains is depicted in Fig. 6(b), revealing extensive variation in the size of $\beta$ grains with different orientations, averaging around $40-150 \mu \mathrm{m}$. Simonelli et al.\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-09(1)}\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-09}

Fig. 6. EBSD images for horizontal layers (scan plane) of Ti-6Al-4V alloys fabricated using different methods: (a) by L-PBF and (b) corresponding reconstructed parent prior $\beta$ grains (Reproduced with permission from Ref. [147]. Copyright (2023), Elsevier), and (c) by EB-PBF, small circles in the IPF at the lower right corner depict the parent $\beta$ grains in the scan plane, with the labeled regions (1)-(7) representing these parent $\beta$ grains (Reproduced with permission from Ref. [148]. Copyright (2020), Elsevier); (d) three dimensional optical images of DED-processed Ti-6Al-4V showing prior $\beta$ grains and $\alpha$ laths (Reproduced with permission from Ref. [149]. Copyright (2021), Elsevier), and (e) results of the $\beta$ grain width and its evolution trend with different laser powers (Reproduced with permission from Ref. [150]. Copyright (2021), Elsevier).

[145] have also reported a prior $\beta$ grain thickness of $103 \mu \mathrm{m}$ in L-PBF-fabricated Ti-6Al-4V alloys. Fig. 6(c) displays the EBSD orientation map of the scan plane for the Ti-6Al-4V alloy produced using EB-PBF. The microstructure reveals basket-weave $\alpha$ laths/plates distributed through the microstructure. The labeled regions (1)-(7) represent the prior parent $\beta$ grains, which exhibit size variations ranging from approximately $50 \mu \mathrm{m}$ to $200 \mu \mathrm{m}$. Similarly, Lancaster et al. [146] have found the thickness size of prior $\beta$ grains to be around $246 \mu \mathrm{m}$ in EB-PBF-fabricated Ti-6Al-4V alloys. Fig. 6(d) presents the three-dimensional (3D) optical images of the Ti-6Al$4 \mathrm{~V}$ counterpart fabricated using DED, where the images reveal the presence of columnar prior $\beta$ grains along the build direction, with an average length ranging from $1.1 \pm 0.4 \mathrm{~mm}$ to $1.7 \pm 0.8 \mathrm{~mm}$. When viewed from the top of the scan plane, equiaxed $\beta$ grains of similar size can be observed, with an average size ranging from $280 \mu \mathrm{m}$ to $510 \mu \mathrm{m}$, with increased laser heat input density from $141 \mathrm{~J} / \mathrm{mm}$ to $283 \mathrm{~J} / \mathrm{mm}$. At higher magnification, the images showcase the presence of needle-like $\alpha$-laths inside the prior $\beta$ grains, with widths ranging from $0.7 \mu \mathrm{m}$ to $1.7 \mu \mathrm{m}$. Fig. 6(e) provides additional insights into the effect of laser power on the width of $\beta$ grains. The results indicate that as the laser power increases, the width of the $\beta$ grains also increases. More specifically, the width rises from $140 \mu \mathrm{m}$ at $700 \mathrm{~W}$ to $160 \mu \mathrm{m}$ at $850 \mathrm{~W}$, and then jumps to $200 \mu \mathrm{m}$ at $1000 \mathrm{~W}$. Moreover, this variation in $\beta$ grain width\\
Table 3

Prior $\beta$ grain size of Ti-6Al-4V alloys fabricated by various AM technologies.

\begin{center}
\begin{tabular}{lllll}
\hline
Material & Method & Grain types & Grain width $($ thickness) $(\mu \mathrm{m})$ & Refs. \\
\hline
Ti-6Al-4V & L-PBF & prior $\beta$ & 87 & $[107]$ \\
Ti-6Al-4V & L-PBF & prior $\beta$ & $40-150$ & $[147]$ \\
Ti-6Al-4V & L-PBF & prior $\beta$ & 103 & $[145]$ \\
Ti-6Al-4V & EB-PBF & prior $\beta$ & 93 & $[107]$ \\
Ti-6Al-4V & EB-PBF & prior $\beta$ & $50-200$ & $[148]$ \\
Ti-6Al-4V & EB-PBF & prior $\beta$ & 246 & $[146]$ \\
Ti-6Al-4V & DED & prior $\beta$ & 202 & $[107]$ \\
Ti-6Al-4V & DED & prior $\beta$ & $280-510$ & $[149]$ \\
Ti-6Al-4V & DED & prior $\beta$ & $140-200$ & $[150]$ \\
Ti-6Al-4V & DED & prior $\beta$ & $200-400$ & $[151]$ \\
Ti-6Al-4V & DED & prior $\beta$ & $300 \pm 60$ & $[152]$ \\
\end{tabular}
\end{center}

has a slight impact on the volume fraction of the basket-weave $\alpha$ phase.

Table 3 provides a summary of the prior $\beta$ grain sizes in Ti-6Al$4 \mathrm{~V}$ alloys produced through different AM methods. As indicated in the table, DED-manufactured Ti-6Al-4V alloys typically exhibit larger prior $\beta$ grains, while those produced via L-PBF have smaller prior $\beta$ grains. EB-PBF-fabricated Ti-6Al-4V alloys fall in between these two categories.\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-10}

Fig. 7. Thermal history of a single layer of Ti-6Al-4V during the production process of (a) L-PBF (Reproduced with permission from Ref. [153]. Copyright (2020), John Wiley and Sons), (b) EB-PBF (Reproduced with permission from Ref. [154]. Copyright (2021), Elsevier), and (c) DED (Reproduced with permission from Ref. [155]. Copyright (2021), Elsevier) techniques.

In the case of AM techniques, which employ a unique layerby-layer production process, the thermal history becomes significantly more complex. This is ascribed to the transfer and accumulation of heat within the deposited layers from newly added layers. Consequently, this phenomenon can significantly affect dwell time within the temperature range required for the nucleation and growth of $\beta$ grains. Fig. 7 shows the thermal history of a single layer of Ti-6Al-4V during different AM techniques. As depicted in Fig. 7(c), DED provides an extended dwell time in the temperature range, conducive to $\beta$ grain growth. On the one hand, the temperature peak reaches the $\beta$ grain growth temperature zone over five times within a single layer. On the other hand, as shown in Fig. 7(a) and (b), the temperature peaks in L-PBF and EB-PBF only reach the $\beta$-grain growth temperature region about three times. The discrepancy can be attributed to the higher penetration depth of DED, which has been discussed in detail in Section 3.1. Therefore, the longer dwell time achieved by DED technology allows for adequate growth time, resulting in larger sizes of prior $\beta$ grains in Ti-6Al-4V alloy compared to counterparts fabricated by L-PBF and EB-PBF.\\
Table 4

$\alpha / \alpha^{\prime}$ lath thickness of Ti-6Al-4V alloys fabricated by different AM technologies.

\begin{center}
\begin{tabular}{lllll}
\hline
Material & Method & Grain types & Grain width (thickness) $(\mu \mathrm{m})$ & Refs. \\
\hline
Ti-6Al-4V & L-PBF & $\alpha / \alpha^{\prime}$ & $0.62-1.32$ & $[160]$ \\
Ti-6Al-4V & L-PBF & $\alpha / \alpha^{\prime}$ & 0.5 & $[161]$ \\
Ti-6Al-4V & L-PBF & $\alpha / \alpha^{\prime}$ & $0.37 \pm 0.10$ & $[156]$ \\
Ti-6Al-4V & L-PBF & $\alpha / \alpha^{\prime}$ & $0.2-1$ & $[72]$ \\
Ti-6Al-4V & L-PBF & $\alpha / \alpha^{\prime}$ & $0.2-1.2$ & $[158]$ \\
Ti-6Al-4V & EB-PBF & $\alpha / \alpha^{\prime}$ & $3.61 \pm 0.92$ & $[156]$ \\
Ti-6Al-4V & EB-PBF & $\alpha / \alpha^{\prime}$ & 1.7 & $[107]$ \\
Ti-6Al-4V & EB-PBF & $\alpha / \alpha^{\prime}$ & $1-1.2$ & $[162]$ \\
Ti-6Al-4V & EB-PBF & $\alpha / \alpha^{\prime}$ & 1.9 & $[163]$ \\
Ti-6Al-4V & EB-PBF & $\alpha / \alpha^{\prime}$ & 1.4 & $[99]$ \\
Ti-6Al-4V & EB-PBF & $\alpha / \alpha^{\prime}$ & $2.09-2.14$ & $[164]$ \\
Ti-6Al-4V & DED & $\alpha / \alpha^{\prime}$ & $0.75-0.8$ & $[165]$ \\
Ti-6Al-4V & DED & $\alpha / \alpha^{\prime}$ & $0.7-1.7$ & $[149]$ \\
Ti-6Al-4V & DED & $\alpha / \alpha^{\prime}$ & 1.81 & $[166]$ \\
Ti-6Al-4V & DED & $\alpha / \alpha^{\prime}$ & 0.7 & $[158]$ \\
Ti-6Al-4V & DED & $\alpha / \alpha^{\prime}$ & 0.8 &  \\
\hline
\end{tabular}
\end{center}

\subsection*{4.1.2. $\alpha / \alpha$ ' lath thickness}
The purpose of this section is to offer a macroscopic understanding of grain size variations in Ti alloys produced by different AM methods by comparing the sizes of $\alpha$ and $\alpha^{\prime}$ grains, collectively referred to as $\alpha / \alpha^{\prime}$ grains. Typically, the unique principles inherent to each AM technique can lead to various sizes of $\alpha / \alpha^{\prime}$ grains in Ti-6Al-4V alloys fabricated by different AM methods. Some researchers have conducted a comparison of the grain size and morphology between Ti-6Al-4V processed through L-PBF and EB-PBF $[156,157]$. Fig. 8(a) presents the EBSD inverse pole figure (IPF) of L-PBF fabricated Ti-6Al-4V, revealing a basketweave needleshaped $\alpha$ ' morphology with a width (thickness) of $0.37 \pm 0.10 \mu \mathrm{m}$ (Fig. 8(c)). In contrast, Fig. 8(b) demonstrates that the EB-PBF fabricated counterpart exhibited a Widmanstätten $\alpha$ morphology with an average thickness of $3.61 \pm 0.92 \mu \mathrm{m}$ (Fig. 8(c)). Fig. 8(d) and (e) shows the $\alpha$ lath displayed an average size of $1.81 \mu \mathrm{m}$ in DED-fabricated Ti-6Al-4V. Lee et al. [158] further conducted a comparison of the needle-like $\alpha / \alpha^{\prime}$ lath between Ti-6Al-4V processed through L-PBF and DED, where in L-PBF, the $\alpha^{\prime}$ laths exhibited a width ranging from $0.2 \mu \mathrm{m}$ to $1.2 \mu \mathrm{m}$ and a length ranging from $2 \mu \mathrm{m}$ to $12 \mu \mathrm{m}$. Contrasted to this, the DED counterpart displayed $\alpha$ laths with a width of $0.8 \mu \mathrm{m}$ and a length of $20 \mu \mathrm{m}$.

Table 4 lists the $\alpha / \alpha^{\prime}$ lath thicknesses reported in recent studies for Ti-6Al-4V alloys fabricated using 3 AM techniques. Based on the results of these studies, it is apparent that the $\alpha / \alpha^{\prime}$ grains generally attain a larger size in EB-PBF-fabricated Ti-6Al-4V alloys than those counterparts fabricated by L-PBF and DED. In addition, the size of $\alpha / \alpha^{\prime}$ grains of DED-fabricated Ti-6Al-4V alloys is slightly bigger than that of L-PBF-fabricated Ti-6Al-4V alloys.

\subsection*{4.1.3. Influence of grain size on mechanical properties}
There is also a significant connection between grain size and the mechanical properties of Ti alloys fabricated by AM techniques. As evident in Table 2, Ti-6Al-4V alloy samples fabricated by LPBF and DED present relatively higher YS and ultimate tensile strength (UTS) compared to those fabricated by EB-PBF. In addition, the highest hardness is observed in Ti-6Al-4V parts fabricated by L-PBF, whereas EB-PBF-fabricated forms present relatively lower hardness when compared to counterparts fabricated by the other 2 AM methods. The improved strength and hardness in Ti alloys fabricated by L-PBF and DED can be attributed to their finer $\alpha / \alpha^{\prime}$ grain size. According to Whang [167], both the strength and hardness of metallic materials generally increase with decreasing grain size, with the maximum strengthening effect observed when grains are refined to 20-30 nm (through the Hall-Petch relationship). The relation between grain size and mechanical properties can be repre-

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-11}
\end{center}

Fig. 8. EBSD IPF images of Ti-6Al-4V (Ti-64) along the build direction for (a) L-PBF, (b) EB-PBF, and (c) corresponding lath size distribution (Reproduced with permission from Ref. [156]. Copyright (2022), Elsevier); (d) DED and (e) corresponding lath size distribution histograms (Reproduced with permission from Ref. [159]. Copyright (2022), Elsevier).

sented by a Hall-Petch equation [168]:


\begin{equation*}
H=H_{0}+k_{\mathrm{H}} \delta_{\alpha_{\text {lath }}}^{-\frac{1}{2}} \tag{2}
\end{equation*}


where $H$ is the hardness or YS, $\delta_{\alpha \text { lath }}$ is the $\alpha$ lath thickness, $H_{0}$ and $k_{\mathrm{H}}$ are constants. It can be seen from Eq. (2) that as grain size decreases, both the YS and hardness of the material are enhanced. This relationship has been supported by the findings of Xiao et al. [169], examining the relationship between grain size and YS of Ti$6 \mathrm{Al}-4 \mathrm{~V}$ using EB-PBF and L-PBF, and concluding that most of the data align with the Hall-Petch theory. Moreover, Galarraga et al. [73] have evaluated the influence of $\alpha$ lath thickness on the mechanical properties of Ti-6Al-4V alloys, showing that increasing the lath thickness of the $\alpha$ grain leads to a slight reduction in UTS, while the YS decreases significantly. Furthermore, the hardness of the Ti-6Al-4V alloy was decreased with an increase in $\alpha$ lath thickness. The relationship between $\alpha$ lath thickness and mechanical properties of Ti-6Al-4V is shown in Fig. 9. As can be seen from Fig. 9(a), the influence of the average $\alpha$ lath width on the tensile YS of L-PBF fabricated Ti-6Al-4V is depicted. The curve represent- ing the Hall-Petch relationship demonstrates a good fit between the width and strength (red symbols), consistent with other findings in the literature. Fig. 9(b) and (c) displays the Hall-Petch relationship between the average width of $\alpha$ lath and the average width of prior $\beta$ grain on the Vickers hardness of EB-PBF fabricated Ti-6Al-4V alloys. Fig. 9(b) demonstrates a significant association between the increase in the average width of $\alpha$ lath and corresponding decrease in Vickers hardness. The fitted curve coefficient $\left(R^{2}=0.91\right)$ applies to both the fabricated rectangular plate (RP) and round bar (RB), indicating a robust relationship between the mean $\alpha$ lath width and hardness in a single fitting line. On the one hand, the significant correlation observed between the average width of $\alpha$ lath and hardness highlights the potential of the average width of $\alpha$ lath as a promising microstructural factor for fine-tuning the mechanical properties of EB-PBF-built Ti-6Al-4V. On the other hand, Fig. 9(c) examines the effect of the average width of prior $\beta$ grain on the hardness of the fabricated RP and RB Ti-6Al-4V samples. The data reveal two separate fitting curves for the fabricated RB and RP samples, with an $R^{2}$ value of $\sim 0.86$.\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-12}

Fig. 9. (a) Influence of $\alpha$ lath thickness (width) on the YS of L-PBF Ti-6Al-4V (Reproduced with permission from Ref. [170]. Copyright (2019), Elsevier), influence of $\alpha$ lath width (b) and prior $\beta$ grain width (c) on the hardness of EB-PBF Ti-6Al-4V (Reproduced with permission from Ref. [171]. Copyright (2019), Elsevier), and (d) effect of $\alpha$ lath size on YS of DED Ti-6Al-4V (Reproduced with permission from Ref. [150]. Copyright (2021), Elsevier).

This $R^{2}$ value is noticeably lower than the $R^{2}$ value of 0.91 for the mean $\alpha$ lath width. Consequently, the results suggest that prior $\beta$ grain width cannot be considered as an important microstructural indicator, and that controlling the $\alpha$ lath width is the primary factor for determining the mechanical properties of EB-PBF-built Ti-6Al-4V. Fig. 9(d) presents a comparison of the influence of $\alpha$ lath width on the tensile YS of DED-fabricated Ti-6Al-4V. The data points were fitted to a Hall-Petch relationship, whereby the results demonstrate that the variation in $\alpha$ lath width correlates well with the YS, with an $R^{2}=0.97937$. In contrast, the fitting of the HallPetch YS curve for $\beta$ grain width (Figure not shown here) only yielded an $R^{2}=0.34831$. These findings suggest that the width of the $\alpha$ lath plays a more important role in affecting the mechanical properties of DED-fabricated Ti-6Al-4V compared to the width of the $\beta$ grains. Moreover, the fitting results in Fig. 9 affirm the reliability of the Hall-Petch relationship in predicting YS and hardness, which can be directly applied to Ti-6Al-4V alloys produced using $\mathrm{L}-\mathrm{PBF}, \mathrm{EB}-\mathrm{PBF}$, and DED, despite the variations in grain sizes and morphologies resulting from these different methods.

This mechanism can be explained by dislocation theory. In crystalline materials, smaller grains create a larger amount of grain boundaries, leading to grain-boundary strengthening. The mechanism of grain-boundary strengthening can be explained with re- gard to two aspects. Firstly, because of the different lattice orientations in neighboring grains, the grain boundaries act as pinning points that prevent the further propagation of dislocations. Secondly, grain boundaries exhibit a high degree of disorder and consist of a large number of dislocations, which also obstruct the continuous movement of dislocations within slip planes. Owing to the limited movement of dislocations, plastic deformation is delayed, thereby enhancing YS and hardness. In addition, fine grains promote a more uniform stress distribution and relieve stress concentration. Therefore, smaller grain size contributes to increased fracture toughness [172], which can result in higher UTS.

\subsection*{4.2. Grain morphology in Ti alloys}
Apart from grain size, the mechanical properties of AMfabricated Ti alloys also can be influenced by grain morphology. The two most common grain morphologies observed in Ti alloys are columnar grains and equiaxed grains, which can occur under different conditions during the production process $[21,173]$. The formation of columnar grains involves the epitaxial growth from parent grains in the substrate across multiple layers and melt pools. In contrast, equiaxed grains form through nucleation and growth within each individual melting pool. Among the mi-\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-13}

Fig. 10. Columnar prior $\beta$ grains formed in Ti-6Al-4V alloys produced by (a) L-PBF (Reproduced with permission from Ref. [40]. Copyright (2016), Elsevier), (b) EB-PBF (Reproduced with permission from Ref. [177]. Copyright (2019), Elsevier), and (c) DED (Reproduced with permission from Ref. [178]. Copyright (2020), Nature Portfolio).

Table 5

Mechanical tensile properties of Ti-6Al-4V alloys fabricated by $3 \mathrm{AM}$ techniques with different building orientations at room temperature.

\begin{center}
\begin{tabular}{|c|c|c|c|c|c|c|c|}
\hline
Material & Method & Condition & Tensile axis orientation & $\sigma_{0.2}(\mathrm{MPa})$ & $\sigma_{\text {UTS }}(\mathrm{MPa})$ & $\varepsilon(\%)$ & Refs. \\
\hline
Ti-6Al-4V & L-PBF & As-built & Horizontal & $1075 \pm 25$ & $1199 \pm 49$ & $7.6 \pm 0.5$ & $[184]$ \\
\hline
Ti-6Al-4V & L-PBF & As-built & Vertical & $967 \pm 10$ & $1117 \pm 3$ & $8.9 \pm 0.4$ & [184] \\
\hline
Ti-6Al-4V & L-PBF & As-built & Horizontal & $1070 \pm 50$ & $1250 \pm 50$ & $5.5 \pm 1$ & [185] \\
\hline
Ti-6Al-4V & L-PBF & As-built & Vertical & $1050 \pm 40$ & $1180 \pm 30$ & $8.5 \pm 1.5$ & [185] \\
\hline
Ti-6Al-4V & EB-PBF & Machined & Horizontal & 1063 & - & 7.1 & [163] \\
\hline
Ti-6Al-4V & EB-PBF & Machined & Vertical & 997 & - & 8.8 & [163] \\
\hline
Ti-6Al-4V & EB-PBF & Machined & Horizontal & $817 \pm 6$ & $916 \pm 13$ & $9.3 \pm 1.6$ & [186] \\
\hline
Ti-6Al-4V & DED & Machined & Horizontal & $960 \pm 26$ & $1063 \pm 20$ & $10.9 \pm 1.4$ & [142] \\
\hline
Ti-6Al-4V & DED & Machined & Vertical & $958 \pm 19$ & $1064 \pm 26$ & $14 \pm 1$ & [142] \\
\hline
Ti-6Al-4V & DED & Machined & Vertical & $941 \pm 4$ & $1027 \pm 1$ & $8.6 \pm 0.1$ & [159] \\
\hline
Ti-6Al-4V & DED & Machined & Horizontal & $1027 \pm 6$ & $1077 \pm 14$ & $2.9 \pm 0.6$ & [187] \\
\hline
Ti-6Al-4V & DED & Machined & Vertical & $1031 \pm 68$ & $1106 \pm 52$ & $6.8 \pm 1.2$ & [187] \\
\hline
\end{tabular}
\end{center}

$\sigma_{0.2}$ - yield strength; UTS - ultimate tensile strength; $\varepsilon$ - elongation.

crostructures of Ti alloys fabricated by most AM techniques, columnar prior $\beta$ grains are more prevalent compared to equiaxed grains. This is ascribed to the complicated thermal history during the AM process, which can result in heat accumulation between newly deposited layers and promote the epitaxial growth of columnar grains. Furthermore, in most AM processes, the chemical components of Ti-6Al-4V between different deposited layers remain nearly identical. This similarity in composition provides essential conditions for grain epitaxial growth because adjacent deposited layers have the same crystal structure. Consequently, the epitaxial growth of metallic grains does not require the nucleation of a new phase [174,175]. As seen in Fig. 10, the columnar prior $\beta$ grains formed in Ti-6Al-4V alloys fabricated by different AM methods. In general, during the solidification process of each layer, columnar grains usually will grow along with the orientation of the temperature gradient $(G)[21,176]$, causing them to extend from the boundary towards the center of the melt pool. Therefore, the growth direction of grains can be significantly influenced by the energy input and the shape of the molten pool that is induced during the process.

The formation of grain boundary $\alpha$ grains (GB- $\alpha$ ) along the prior- $\beta$ grain boundaries in Ti-6Al-4V alloys produced via different $\mathrm{AM}$ processes is a consequence of the diffusion transition from $\beta$ to $\alpha$ phases. As depicted in Fig. 11, elongated GB- $\alpha$ grains are evident along the prior- $\beta$ grain boundaries in Ti-6Al-4V alloys fabricated through L-PBF, EB-PBF, and DED techniques. Notably, the relatively higher cooling rates in DED and L-PBF result in smaller GB- $\alpha$ grain sizes in Ti alloys compared to EB-PBF-fabricated Ti-6Al$4 \mathrm{~V}$ alloys.

The existence of GB- $\alpha$ grains and columnar prior- $\beta$ grains is generally considered as the main contributor to anisotropy observed in the mechanical properties of Ti alloys [182,183]. Thus, Ti alloys fabricated by AM techniques usually exhibit anisotropy in mechanical properties. Table 5 highlights the anisotropy in the mechanical behaviors of AM-built Ti-6Al-4V. Two phenomena are particularly noteworthy here. Firstly, the tensile strength and ductility present a converse trend with respect to the anisotropy property [184]. Secondly, horizontally fabricated Ti-6Al-4V typically exhibit higher strengths than vertically fabricated samples across all 3 AM techniques. Conversely, vertically manufactured parts exhibit better ductility than horizontally manufactured parts for Ti-6Al-4V alloys built by the three studied AM techniques.

Anisotropy in ductility can be ascribed to the difference in loading conditions experienced by GB- $\alpha$ grains and columnar prior $\beta$ grains when subjected to stress in different directions [142]. As shown in Fig. 12(a-c), when a tensile load is applied vertically on the boundaries of prior $\beta$ grains, the short axes of the prior$\beta$ grains and GB- $\alpha$ grains undertake the loads, resulting in the separation of adjacent prior- $\beta$ grains and intergranular fracture. In contrast, as shown in Fig. 12(d), when the applied load is parallel to the boundaries of prior $\beta$ grains, the entire prior- $\beta$ grains and GB- $\alpha$ grains share the load, resulting in an increased effective slip distance between neighboring grains. Therefore, compared to horizontally fabricated samples, vertically fabricated Ti-6Al-4V presents better ductility. In addition to ductility, strength anisotropy can be ascribed to varying numbers of prior- $\beta$ grains and GB- $\alpha$ grains formed in Ti-6Al-4V samples fabricated by AM techniques with different build orientations. Compared to vertically fabricated samples, the horizontally fabricated Ti-6Al-4V parts involve more prior$\beta$ grains. This disparity arises because the column $\beta$ grains usually tend to continuously grow along the build orientation, resulting in elongation of the $\beta$ grains but not a significant change in $\beta$ grains quantity. Thus, according to the Hall-Petch relationship, horizontally fabricated Ti-6Al-4V samples have more prior- $\beta$ grain boundaries which obstruct dislocation slip and enhance the strength of samples [140,188].\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-14}

Fig. 11. GB- $\alpha$ grains in Ti-6Al-4V alloys manufactured by (a, b) L-PBF (Reproduced with permission from Ref. [179]. Copyright (2021), Elsevier), (c, d) EB-PBF (Reproduced with permission from Ref. [180]. Copyright (2020), Elsevier) and (e, f) DED (Reproduced with permission from Ref. [181]. Copyright (2021), Elsevier).

\subsection*{4.3. Columnar to equiaxed transition (CET)}
Apart from columnar grains, researchers have also attempted to realize the columnar to equiaxed transition (CET) in AM-built Ti alloys. One approach to achieve CET is by adjusting printing parameters to change the ratio of temperature gradient $(G)$ and solidification rate $(R)$, denoted as $G / R[21,189]$. A low $G / R$ is favorable for promoting CET [21]. Bontha et al. [190] have explored the influences of processing parameters on the solidification mode of\\
DED-fabricated Ti-6Al-4V alloys by conducting analytical (Rosenthal) and numerical (FEM) modeling. They have indicated that by tuning process variables, such as scan speed and laser power, a graded microstructure along the build deposition depth (height) can be formed, particularly transitioning from columnar to mixed or equiaxed morphology near the surface region with higher laser powers, while keeping the scan speed constant. In a study by Wu et al. [151], these authors found that the length of columnar grains became shorter and was substituted by very large equiaxed grains

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-15(1)}
\end{center}

Fig. 12. (a) Morphology of DED-fabricated Ti-6Al-4V alloys, (b) highlighted rectangular section used for schematic diagram of GB- $\alpha$ and prior- $\beta$ grains subjected to loads in different directions: (c) represent the load orientation perpendicular to the long axes of prior- $\beta$ grains and (d) represent the load orientation parallel to the long axes of prior- $\beta$ grains (Reproduced with permission from Ref. [142]. Copyright (2015), Elsevier).\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-15}

Fig. 13. (a) Equiaxed grains partially form in the DED-fabricated Ti-6Al-4V alloys when adopting the laser power of $1 \mathrm{~kW}$ (Reproduced with permission from Ref. [191]. Copyright (2022), Springer Nature) and (b) equiaxed grains of the L-PBF-built Ti-6Al-4V alloy (Reproduced with permission from Ref. [113]. Copyright (2017), Elsevier).

in the DED-built Ti-6Al-4V alloys near the sample's bottom location when an extremely high laser power was used, while other parameters were kept constant. Similarity, as shown in Fig. 13(a), Marshall et al. [191] also found whilst the high laser power (1 kW) was applied, a partially equiaxed microstructure can occur in the DED-fabricated Ti-6Al-4V alloys because of the reduced $G$. In addition, as seen in Fig. 13(b), Xu et al. [113] demonstrated that equiaxed grains also occurred in L-PBF-built Ti-6Al-4V alloys when adopting a small layer thickness and inter-layer time. This approach promoted heat accumulation, resulting in reduced $G$ and facilitated CET. Although the number of cases is limited, these previous cases have shown that appropriately increasing energy density is an effective strategy to decrease $G$ to achieve CET in Ti alloys manufactured by PBF and DED.

In addition to adjusting $G$ and $R$, providing extra favorable conditions for heterogeneous nucleation is another effective way to realize CET. Zhang et al. [192] have reported that utilizing a combination of high powder feed rate and lower energy density promotes the formation of equiaxed grains in DED-fabricated Ti-6Al-2Sn-2Zr3Mo-1.5Cr-2Nb alloy, as shown in Fig. 14(b). The increased amount of unmelted powder in the molten pool created by the high powder feed rate favors heterogeneous nucleation and promotes CET. In addition, the use of lower laser energy density can also decrease the rate of powder melting, resulting in an increased nucleation

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-15(2)}
\end{center}

Fig. 14. Equiaxed grains formed in the (a) bottom part of EB-PBF-fabricated Ti6Al-4V alloy (Reproduced with permission from Ref. [194]. Copyright (2013), Elsevier), (b) DED-fabricated Ti-6Al-2Sn-2Zr-3Mo-1.5Cr-2Nb (Reproduced with permission from Ref. [192]. Copyright (2016), Elsevier) and (c) Ti-6.5Al-3.5Mo-1.5Zr-0.3Si alloys (Reproduced with permission from Ref. [173]. Copyright (2015), Elsevier).

rate. Similarly, as shown in Fig. 14(c), Wang et al. [173] indicated that equiaxed grains can nearly fully form in DED-fabricated Ti6.5Al-3.5Mo-1.5Zr-0.3Si alloys under conditions of relatively higher powder feed rate and lower energy density. A similar phenomenon was reported in DED-fabricated Ti-25V-15Cr-2Al-0.2C alloys in the research of Ref. [193]. Moreover, Antonysamy et al. [194] noted equiaxed grains forming in the bottom part of EB-PBF-fabricated Ti-6Al-4V alloys (Fig. 14(a)), which could be ascribed to heterogeneous nucleation induced by partly unmelted powder. During the AM production process, the build plate serves as a thermal conductor, often referred to as a "heat sink" [28]. As the build height increases, the efficiency of thermal conductivity from the newly deposited layer to the build plate decreases. Consequently, powder in the upper layers retains more heat energy, enabling better melting. However, this thermal limitation can lead to incomplete powder melting, especially in the lower layers, due to rapid heat dissipation from the build plate. This phenomenon somehow can impact overall print quality and part integrity.

Although CET can be achieved by increasing powder feed rate and decreasing energy density, these changed printing parameters may result in undesired side effects, such as increased porosity, cracks, delamination, and others. Therefore, it is ideal to achieve CET while still utilizing optimized printing parameters. In recent research by Ref. [178], high-intensity ultrasound was verified as an effective method for achieving the CET in Ti-6Al-4V alloy during the fabrication process of DED (Fig. 15(a)). This is due to the induction of acoustic cavitation in the liquid by ultrasonic irradiation, which agitates the metallic liquid and promotes nucleation during the solidification process [195,196]. In addition, Todaro et al. [178] reported that ultrasonic irradiation does not affect the melted rate, as both ultrasound-treated and untreated DEDfabricated Ti-6Al-4V alloys exhibited similar porosity levels. Some researchers have also attempted to achieve CET from alloy constitution modification. It has been observed that larger values of the growth restriction factor $(Q)$ facilitate more nucleation and influence the development of equiaxed grains. According to Kozlov and\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-16}

Fig. 15. (a) OM image of the DED-built Ti-6Al-4V alloy with ultrasound (scale bars, $1 \mathrm{~mm}$ ) (Reproduced with permission from Ref. [178]. Copyright (2020), Nature Portfolio); (b) OM image of the EB-PBF-fabricated Ti6Al4V-7Cu alloy (Reproduced with permission from Ref. [201]. Copyright (2022), Elsevier), and (c) OM micrograph of the DEDfabricated Ti-8.5Cu alloy (Reproduced with permission from Ref. [199]. Copyright (2019), Springer Nature).

Schmid-Fetzer [197], a higher $Q$ not only effectively refines grains, but also affects the formation of equiaxed grains. Based on this principle, some research groups have investigated the addition of alloying elements such as beryllium (Be), silicon (Si) or boron (B) to control the grain microstructure of AM-fabricated Ti alloys [198]. However, the addition of those alloying elements usually results in a decrease in the size of the columnar grains or only achieves partial CET in AM-fabricated Ti alloys. Recent research by Zhang et al. [199] has overcome this challenge. As shown in Fig. 15(c), they found that by using the same printing parameters, the addition of copper $(\mathrm{Cu})$ enables DED-fabricated Ti-8.5Cu to fully consist of equiaxed grains, while DED-fabricated Ti-6Al-4V alloys still comprised of columnar grains. This can be ascribed to the significantly higher $Q(62 \mathrm{~K})$ of the Ti-8.5Cu alloy compared to the $Q$ value of only $8 \mathrm{~K}$ in Ti-6Al-4V alloys. Similarly, Choi et al. [200] employed mixed powders of Ti-6Al-4V and Co-Cr-Mo alloys as feedstock, where they reported that equiaxed grains can fully form in DED-fabricated samples containing $5 \mathrm{wt} \%$ and $10 \mathrm{wt} \%$ Co due to the high $Q$ value of Co. Furthermore, as seen in Fig. 15(b), Mosallanejad et al. [201] also proved that an equiaxed microstructure can be obtained in the EB-PBF-fabricated Ti-6Al-4V-7Cu alloy by using a mixed powder of Ti-6Al-4V and Cu.

Given the existing challenges in achieving CET in Ti alloys, there are only a limited number of studies focused on the me- chanical properties of AM-fabricated Ti alloys containing equiaxed grains. As shown in Fig. 16(a), Xu et al. [113] reported that L-PBFfabricated Ti-6Al-4V alloys with equiaxed grain structures exhibit enhanced strength compared to the mill-annealed Ti-6Al-4V alloy. This improvement is attributed to the finer grain size resulting from the rapid cooling rate during L-PBF. In contrast, when comparing these equiaxed-grain Ti-6Al-4V alloys to typical L-PBFfabricated ones, a slight reduction in strength is observed, but with an increase in ductility. This variance is attributed to the higher energy density utilized in the process, promoting the formation of lamellar $(\alpha+\beta)$ phases within the prior $\beta$ equiaxed grains instead of $\alpha^{\prime}$ martensite, which is less favorable for ductility. Moreover, the unique morphology and the absence of continuous GB- $\alpha$ grains in equiaxed grain structures significantly reduce stress concentrations, effectively enhancing material crack resistance and ductility. In another study by Todaro et al. [178], as shown in Fig. 16(b), both YS and UTS can be improved by $\sim 12 \%$ for DED-fabricated Ti-6Al-4V alloys with ultrasound treatment; however, the ductility of DED-built Ti-6Al-4V alloys remains similar, regardless of ultrasound treatment. Compared with columnar grains, the size of equiaxed grains is smaller in DED-built Ti-6Al-4V alloys with ultrasound treatment. As such, achieving CET through various methods in the AM of Ti alloys can provide more flexibility to adjust mechanical properties.\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-17(1)}

Fig. 16. (a) Engineering stress-strain curves of (a) mill-annealed and L-PBF-built Ti-6Al-4V alloys (Reproduced with permission from Ref. [113]. Copyright (2017), Elsevier) and (b) DED-built Ti-6Al-4V alloys with and without ultrasound treatment (Reproduced with permission from Ref. [178]. Copyright (2020), Nature Portfolio).

\section*{5. Defects analysis and mitigation}
\subsection*{5.1. Porosity}
Porosity is a normal defect in metallic materials fabricated through AM techniques, where its complete elimination remains a challenge [202,203]. Porosity represents the void volume in a material, quantified as the pore volume to the total volume ratio of the material. During AM, the formation of pores can be induced by a series of complex phenomena, which are still subject to ongoing debate regarding their origins [204]. In general, the incomplete closure of a keyhole, entrapped gas and lack of fusion (LOF) are deemed the primary factors for the formation of pores in AMfabricated Ti alloys [66].

The keyhole formation is commonly induced by excessively high-power density during the AM production process. These keyholes, which can be unstable and prone to collapse, trap internal vapor and surrounding gas, leading to the creation of pores in the melt pool [21]. Keyhole-induced pores are usually irregular and depend on the size and shape of keyholes, as shown in Fig. 17(a) and (b). Entrapped gas also can contribute to pores during the AM process. The entrapped gas can originate from two main sources. Firstly, during the powder production stage, the atomization process could cause gas to remain within the powder particles [205]. Subsequently, the gas can be released and retained within the material during the AM process. Tammas-Williams et al. [206] have reported that some small spherical pores observed in EB-PBF-fabricated Ti-6Al-4V originated from the argon gas carried by the powder feedstock during atomization. Secondly, shielded gas used during the AM process can also become entrapped in the melt pool, resulting in pore formation. Generally, pores induced by entrapped gas tend to be nearly spherical and tiny, as shown in Fig. 17(c). Improper printing parameters, such as the excessively high scan speed, low power from heat source or large hatching space, can result in insufficient energy density, which prevents complete melting and results in LOF. When a new layer is deposited over an area with LOF, pores form between the boundaries of the two adjacent layers [207]. LOF-induced pores, which are usually bigger and irregular compared with pores induced by entrapped gas [208], are shown in Fig. 17(d).

In general, the amount, shapes and sizes of pores vary significantly in Ti alloys fabricated by different AM techniques. As shown in Table 6, L-PBF-fabricated Ti alloys exhibit more pores and the highest porosity, while Ti alloys fabricated by EB-PBF and DED tend to have fewer pores and relatively lower porosity. According to Liu\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-17}

Fig. 17. Different types of pores induced by (a) (Reproduced with permission from Ref. [62]. Copyright (2016), Elsevier) and (b) keyhole (Reproduced with permission from Ref. [209]. Copyright (2020), Elsevier), (c) entrapped gas (Reproduced with permission from Ref. [210]. Copyright (2020), Elsevier), and (d) LOF in Ti alloys manufactured by AM techniques (Reproduced with permission from Ref. [211]. Copyright (2015), Elsevier).

et al. [62], L-PBF-fabricated Ti-2448 alloys contain ten times more defects (pores) compared to those fabricated by EB-PBF. On the one hand, the majority of pores that occur in L-PBF-fabricated Ti alloys are often bigger and irregular in shape, with some categorized as pores induced by keyhole and LOF. On the other hand, Ti alloys fabricated by EB-PBF and DED methods typically exhibit smaller, nearly spherical pores, with most pores being gas-induced.

The distinctions in pore characteristics among the 3 AM techniques can be attributed to their different manufacturing processes. The typical melting process of EB-PBF and L-PBF is presented in Fig. 18. According to Liu et al. [62], L-PBF, with its smaller size of laser beam spot, effects of ray reflection and discontinuous scanning mode, can easily form deep key holes, resulting in the creation of more conical keyhole-induced pores. In contrast, EB-PBF, characterized by a larger electron beam spot and continuous scanning track, typically forms a wider melt pool, resulting in the presence of spherical pores in Ti alloys [62]. In addition, the L-PBF process is carried out in a shield gas-filled chamber, which can potentially cause the entrapment of gas in the melt pool [21]. Conversely, EB-PBF employs a vacuum building chamber, effectively

Table 6

Relative porosity and defect size and shape of AM-fabricated Ti alloys.

\begin{center}
\begin{tabular}{|c|c|c|c|c|c|}
\hline
Material & Method & Porosity (\%) & Equivalent diameter $(\mu \mathrm{m})$ & Pore shapes & Refs. \\
\hline
CP-Ti & L-PBF & $0.27-0.25$ & - & - & $[212]$ \\
\hline
CP-Ti & L-PBF & $3.6-0.5$ & - & - & $[130]$ \\
\hline
CP-Ti & L-PBF & $1.1-0.4$ & - & Irregular & $[94]$ \\
\hline
CP-Ti & L-PBF & $1.9 \pm 0.51$ & - & - & $[213]$ \\
\hline
CP-Ti & EB-PBF & 0.3 & - & - & $[96]$ \\
\hline
CP-Ti & DED & $2.5-0.5$ & - & Irregular & $[94]$ \\
\hline
Ti-6Al-4V & L-PBF & 0.01 & $8-63$ & Spherical; flat & $[214]$ \\
\hline
Ti-6Al-4V & L-PBF & 0.29 & $3-108$ & Irregular & [99] \\
\hline
Ti-6Al-4V & L-PBF & 0.5 & - & - & $[215]$ \\
\hline
Ti-6Al-4V & EB-PBF & 0.005 & $3-28$ & Spherical & [99] \\
\hline
Ti-6Al-4V & EB-PBF & 0.0397 & 13.3 & Spherical & $[206]$ \\
\hline
Ti-6Al-4V & EB-PBF & 0.25 & - &  & $[215]$ \\
\hline
Ti-6Al-4V & EB-PBF & 0.072 & 81.7 & Spherical & $[216]$ \\
\hline
Ti-6Al-4 V ELI & EB-PBF & 0.15 & - & Spherical; flat & $[217]$ \\
\hline
Ti-6Al-4V & DED & 0.01 & $26.12-128.98$ & Spherical & $[218]$ \\
\hline
Ti-6Al-4V & DED & 0.1 & $1-3$ & Spherical & $[75]$ \\
\hline
Ti-6Al-4V & DED & 0.00066 & - & Spherical & $[219]$ \\
\hline
Ti-6Al-4V & DED & 0.09 & - & - & $[215]$ \\
\hline
Ti-48Al-2Nb-0.7Cr-0.3Si & EB-PBF & $0.17 \pm 0.24$ & 30 & Spherical & $[220]$ \\
\hline
$\mathrm{Ti}-34 \mathrm{Nb}$ & L-PBF & $0.12 \pm 0.1$ & - & Spherical, irregular & $[221]$ \\
\hline
Ti-35Nb & L-PBF & $\pm 0.6$ & - & \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-18}
 & $[38]$ \\
\hline
Ti-35Nb-7Zr-5Ta & L-PBF & $1-7$ & - & - & $[123]$ \\
\hline
Ti-35Nb-7Zr-5Ta & L-PBF & $1.5-0.5$ & - & Spherical, irregular & $[125]$ \\
\hline
Ti-2448 & L-PBF & 0.37 & $20-140$ & Irregular; conical & $[62]$ \\
\hline
Ti-2448 & EB-PBF & 0.06 & $20-120$ & Spherical & $[62]$ \\
\hline
TI-6Al-2Sn-4Zr-2Mo & EB-PBF & $0.6-0.8$ & - & Spherical & $[222]$ \\
\hline
\end{tabular}
\end{center}

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-18(1)}
\end{center}

\section*{Electron beam melting process}
Fig. 18. Simplified diagram of the melting process of L-PBF and EB-PBF (Reproduced with permission from Ref. [62]. Copyright (2016), Elsevier).

avoiding shield gas-induced pores in Ti alloys. Further, Gaytan et al. [223] have noted that the pores in EB-PBF-fabricated Ti alloys are primarily caused by gas contained in the powder feedstock and the vaporization of low-boiling elements [223]. Moreover, due to the relatively lower energy density and the higher likelihood of LOF during the L-PBF process, elongated and flat pores are more commonly observed in L-PBF-fabricated Ti alloys, resulting in a higher pore content compared to EB-PBF-fabricated Ti alloys. In contrast to PBF systems, the DED process directly adds metallic powder through a nozzle into the active melt pool without a pre-deposited powder bed. This can result in insufficient release of gas from the bottom of the powder bed, leading to the formation of gas-induced pores. In addition, the size of the laser beam spot in DED is similar to that of EB-PBF and bigger than that of L-PBF [72], resulting in fewer keyhole-induced pores in Ti alloys fabricated by DED. Moreover, due to heat accumulation and higher energy density in the DED process, LOF occurs less frequently in DED-fabricated Ti alloys. Therefore, DED-fabricated Ti alloys generally exhibit relatively lower porosity compared to Ti alloys fabricated by PBF.

\subsection*{5.2. Surface roughness}
Surface roughness is also a common defect in Ti alloys manufactured by AM techniques. This is generally related to two factors:

Table 7

Surface roughness of Ti-6Al-4V alloys manufactured by different AM techniques.

\begin{center}
\begin{tabular}{lllll}
\hline
Method & L-PBF & EB-PBF & DED & Refs. \\
\hline
Actual surface roughness (ASR, $\mu \mathrm{m})$ & 8.485 & $28.803-45.7$ & $7.867-63.9$ & $[105,225,226]$ \\
Edge surface roughness (ESR, $\mu \mathrm{m})$ & $5-40$ & $25-131$ & $0.24-13.3$ & $[227-230]$ \\
\hline
\end{tabular}
\end{center}

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-19}
\end{center}

Fig. 19. Surface morphology of Ti-6Al-4V manufactured by (a) L-PBF and (b) EB-PBF (Reproduced with permission from Ref. [231]. Copyright (2018), MDPI).

the actual surface roughness (ASR) and the edge surface roughness (ESR). The ASR presents the surface roughness of the flat layer and can be calculated using Eq. (3) [224]:


\begin{equation*}
R_{\mathrm{as}}=\frac{1}{N} \sum_{i=1}^{N}\left|f_{\mathrm{n}}\right| \tag{3}
\end{equation*}


where $R_{\mathrm{as}}$ is the ASR, $f_{n}$ is the height of a peak or the depth of a valley in one measured position, and $N$ represents the number of measured positions. The ASR can vary among Ti alloys manufactured by different AM techniques. A comparison in Table 7 reveals that the Ti-6Al-4V alloy fabricated by both L-PBF and DED methods can achieve relatively lower ASR in comparison with the counterparts manufactured by EB-PBF.

Two reasons can explain this phenomenon. Firstly, the size of powder feedstock used in L-PBF is smaller than the powder adopted in EB-PBF. As reflected in Fig. 19, many larger unmelted metallic powders adhere to the surface of Ti-6Al-4V fabricated by the EB-PBF, while the size of powders is much smaller on the surface of L-PBF-fabricated Ti-6Al-4V alloy. In addition, the DED technique has a special feeding system that allows unused metallic powders to be blown away, resulting in fewer powder particles adhering to the surface of the fabricated products. Therefore, Ti alloys fabricated by L-PBF and DED usually can achieve lower ASR than those of EB-PBF-fabricated Ti alloys.

In addition to ASR, the ESR also can effectively influence the roughness of products fabricated by AM techniques. The "stair step effect" is the primary factor contributing to ESR, as shown in Fig. 20(a). This effect occurs when a thicker deposited layer and more inclined build direction are present, leading to an increase in the severity of the "stair step effect". The average roughness $\left(R_{\mathrm{a}}\right)$ caused by "stair step effect" can be calculated using Eq. (4) [232]:


\begin{equation*}
R_{\mathrm{a}}=1000 t_{\mathrm{l}} \sin \left(\frac{90-\theta}{4}\right) \tan (90-\theta) \tag{4}
\end{equation*}


where $t_{1}$ is the layer thickness and $\theta$ represents the build angle. Table 7 demonstrates that both L-PBF and DED techniques yield Ti alloys with lower ESR compared to the EB-PBF technique. Fig. 20(b) shows that the thicker deposited layer $(70 \mu \mathrm{m})$ during the EB-PBF contributes to higher ESR in fabricated Ti-2448 alloys compared to deposited layers $(50 \mu \mathrm{m})$ in the L-PBF production process. This disparity in layer thickness leads to the higher ESR observed in EBPBF-fabricated Ti-2448 alloys. Furthermore, in the case of DED, although the deposited layer is also thick, the inter-layer fusion is more pronounced because of the higher energy density and heat accumulation effect during the DED process. Therefore, Ti alloys fabricated by both L-PBF and DED present higher surface accuracy than those fabricated by EB-PBF.

\subsection*{5.3. Residual stress}
Residual stress (RS) is stress that exists in solid material when products are under the conditions of thermal equilibrium and not subjected to external stress [233]. The RS in metals is usually originated from the large-scale temperature variation experienced during the production process. According to DebRoy et al. [28], when a metallic material suffers substantial temperature changes during production, the thermal strain of the material can exceed the limitation of its elastic strain, leading to the conversion of strain from elastic to plastic and resulting in stress accumulation. The relationship between the thermal strain of metals and temperature change can be expressed as Eq. (5) [28]:

$\varepsilon_{\mathrm{T}}=\alpha\left(T-T_{0}\right)$

where $\varepsilon_{\mathrm{T}}$ is the thermal strain of metals, $\alpha$ is the coefficient of thermal expansion (CTE) of material, $T$ and $T_{0}$ are the local temperature and defined initial temperature, respectively. The CTE of Ti alloys is usually above $1 \times 10^{-5} \mathrm{~K}^{-1}$, which means even a temperature change of only a few hundred $\mathrm{K}$ can cause stress accumulation [28]. Therefore, due to the large thermal gradient $(G)$ in the melt pool, RS can more easily occur in Ti alloys manufactured by AM techniques compared to those produced with traditional methods [234]. The L-PBF method, characterized by large G and rapid cooling rates, tends to result in relatively higher levels of RS [235]. XRD patterns are commonly used to analyze the magnitude and type of RS in material. A right shift in the diffraction peak represents the compressive RS occurring in the fabricated material, while a left shift indicates tensile RS. Moreover, the larger the displacement distance (shift) of the diffraction peak, the larger the $\mathrm{RS}$ remains near the material's tested surface region. As shown in the inset of Fig. 21, compared to EB-PBF-fabricated Ti-6Al-4V alloys, the diffraction peaks of $\alpha / \alpha^{\prime}$ phase in L-PBF-fabricated Ti-6Al$4 \mathrm{~V}$ exhibit a greater right shift, which indicates the presence of higher compressive RS near the region. In addition, Sharma et al. [236] have also reported a high RS of $295 \pm 12.5 \mathrm{MPa}$ in L-PBFfabricated Ti-6Al-4V alloy. Szost et al. [237] have also indicated that Ti-6Al-4V fabricated by DED has a maximum tensile RS of $280 \mathrm{MPa}$ and a maximum compressive RS of $250 \mathrm{MPa}$. Contrastingly, compared to Ti alloys fabricated by L-PBF and DED methods, the RS is very small and can often be ignored in Ti alloys fabricated by the EB-PBF $[105,238,239]$. This is because the production process of EB-PBF typically involves prolonged exposure to high build platform temperatures $\left(\sim 650-750{ }^{\circ} \mathrm{C}\right)$ [73], which is similar to the annealing process. Therefore, the process of producing EB-PBF can effectively mitigate RS in final fabricated Ti alloys.

The distribution of RS can vary in Ti alloys manufactured by AM techniques. The RS distribution of Ti-6Al-4V fabricated by L$\mathrm{PBF}$ is shown in Fig. 22, where it can be seen that high tensile RS occurs near the surface edge surrounding area, which gradually transitions into compressive RS towards the central area of surface layer. Furthermore, Ahmad et al. [240] have shown that the maximum tensile RS ranges from 805 to $837 \mathrm{MPa}$ in the area that is

\begin{center}
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(a)

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\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-20}
\end{center}

EBM sample

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\end{center}

SLM sample

\begin{center}
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\end{center}

EBM melting process

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-20(4)}
\end{center}

(b)

Fig. 20. (a) Schematic diagram of the "stair step effect" (Reproduced with permission from Ref. [232]. Copyright (2015), Springer Nature) and (b) surface morphologies and melting schematic diagram of Ti-2448 alloys manufactured by L-PBF and EB-PBF (Reproduced with permission from Ref. [62]. Copyright (2016), Elsevier).

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-20(5)}
\end{center}

Fig. 21. XRD pattern of Ti-6Al-4V alloys fabricated by L-PBF and EB-PBF (Reproduced with permission from Ref. [104]. Copyright (2021), Elsevier).

near the surface edge surrounding area, while the maximum compressive RS ranges from 434 to $459 \mathrm{MPa}$ in the area close to the center of surface layer in the Ti-6Al-4V alloy manufactured by L-\\
PBF. Moreover, Yakout et al. [241] have reported tensile RS in the range of 445-528 MPa at a depth of $0.4 \mathrm{~mm}$ below the top surface, which decreases to around 20-264 MPa at a depth of $1 \mathrm{~mm}$ below the top surface.

\subsection*{5.4. Inhomogeneity and chemical composition alteration}
The design types and volume of constituent elements are crucial in determining the properties of the final fabricated products. Even minor changes in the chemical composition can have a substantial influence on the performance of products. However, it is currently challenging to completely avoid chemical composition changes in final products fabricated by AM techniques. Therefore, it is necessary to gain a deep understanding of the mechanism behind the chemical composition changes in Ti alloys during different AM techniques to improve the stability of the properties of final products. Interstitial element pickup and alloying elements evaporation are commonly two main sources for chemical composition changes. The interstitial elements pickup typically leads to an increase in the total element content in AM-fabricated Ti alloys. Conversely, the evaporation or loss of alloying elements results in a decrease in their concentration in the alloys. Accordingly, understanding and controlling these factors are crucial to guaran-\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-20(6)}

Fig. 22. Stress distribution at the surface layer (scan plane) of Ti-6Al-4V alloys manufactured by L-PBF: (a) contour stress maps and (b) stress distribution along the diagonal (Reproduced with permission from Ref. [240]. Copyright (2018), Elsevier).\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-21}

Fig. 23. (a) Relationship between the interstitial element concentration in Ti-6Al-4V alloys manufactured by L-PBF and the 0 concentration of building chamber (Reproduced with permission from Ref. [242]. Copyright (2020), Elsevier) and (b) trend of O concentration and hardness of L-PBF-built Ti-6Al-4V alloys with the increase of laser beam exposure time from increasing amount of laser passes (Reproduced with permission from Ref. [243]. Copyright (2019), Elsevier).

tee the stability of desired chemical constitution and properties of AM-fabricated Ti alloys.

\subsection*{5.4.1. Interstitial element pickup}
During the AM process, additional impurities can be introduced into the fabricated Ti alloys from the surrounding environment. Researchers have investigated the effects of oxygen $(0)$ concentration and laser parameters on chemical composition changes in Ti alloys fabricated using the L-PBF. Dietrich et al. [242] have investigated chemical composition changes in L-PBF-fabricated Ti-6Al-4V alloys under different $O$ concentrations in the building chamber. As shown in Fig. 23(a), increasing the 0 concentration in the build chamber results in an increase in the $\mathrm{O}$ and nitrogen $(\mathrm{N})$ content in the Ti-6Al-4V alloy compared to the powder raw material. The 0 content increased by up to $\sim 41 \%$ (from $\sim 0.13 \mathrm{wt} . \%$ to $\sim 0.18 \mathrm{wt} . \%$ ), and the $\mathrm{N}$ content increased by $\sim 222 \%$ (from $\sim 0.0062 \mathrm{wt} \%$ to $\sim 0.02 \mathrm{wt} . \%$ ). Additionally, the content of extra interstitial elements can also change in AM-fabricated Ti alloys, even at a constant gas concentration in the building chamber. Velasco-Castro et al. [243] have applied L-PBF to produce Ti-6Al-4V samples with different numbers $(1,3,5,9$ beam passes) of repeated passes of the beam on each layer, whereby there was more oxygen in LPBF-built Ti-6Al-4V samples with the greater number of beams passes in each layer. This is because, as shown in Fig. 23(b), more beams passing increases the total energy input, which results in a longer melting time during the production process. As a result, the molten pool will repeatedly be exposed to the building chamber, promoting the absorption of interstitial elements. Similarly, $\mathrm{Na}$ et al. [244] have reported $\mathrm{O}$ and $\mathrm{N}$ elements to increase in L-PBFbuilt CP-Ti alloys with increased laser power. However, Wang et al. [39] have stated that the L-PBF-built Ti-35Nb alloy ( 0.12 wt.\% 0) using prealloyed powder exhibited a very low $\mathrm{O}$ pick-up of approximately $0.02 \mathrm{wt} . \%$ compared to the initial $0.10 \mathrm{wt} \% \mathrm{O}$ in the prealloyed powder under optimized processing parameters. Conversely, the Ti-35Nb produced using a powder mixture showed a higher $\mathrm{O}$ content of $0.30 \mathrm{wt} . \%$ as a result of incorporating elemen-

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-22}
\end{center}

(a)

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-22(2)}
\end{center}

(b)

Fig. 24. (a) Ti-6Al-4V alloy manufactured by DED (the color of the O-contaminated area is darker) (Reproduced with permission from Ref. [142]. Copyright (2015), Elsevier) and (b) concentration of $\mathrm{N}$ and $\mathrm{O}$ elements in the powder feedstock and Ti-6Al-4V samples built by DED with different processing conditions: ArS (Argon Shielding, general condition), SE-HB (Sealed Environment, Hot Base), SE-CB (Sealed Environment, Cold Base) (Reproduced with permission from Ref. [245]. Copyright (2020), MDPI).

tal Ti powder (with 0.10 wt.\% O) and Nb powder (with 0.32 wt.\% 0). Additionally, both built prealloyed and powder mixture Ti-35Nb showed very low $\mathrm{N}$ content, with values of $0.024 \mathrm{wt} \% \mathrm{~N}$ and 0.030 wt.\% N, respectively.

With regard to Ti alloys fabricated using the DED process, studies have also reported the occurrence of additional interstitial elements, particularly O. Carroll et al. [142] have found that the 0 content of DED-fabricated Ti-6Al-4V alloy was about 0.2046 wt.\%, approximately $0.0316 \mathrm{wt}$.\% higher than the 0 content of the initial powder. In addition, due to water-cooling system leaking, they reported that in areas affected by leaked water vapor, the $O$ content was measured to be 0.2170 wt.\%. As shown in Fig. 24(a), the intrusion of $\mathrm{O}$ caused oxidation and darkening of the lower parts of the fabricated products. Moreover, Carrozza et al. [245] have also found both $\mathrm{O}$ and $\mathrm{N}$ contents to increase in DED-fabricated Ti-6Al$4 \mathrm{~V}$ alloys, with the $\mathrm{O}$ content showing a greater increase compared to the $\mathrm{N}$ content, as indicated in Fig. 24(b).

Although L-PBF and DED processes typically take place in a building chamber with an inert gas atmosphere provided by feeding systems, the presence of mixed interstitial elements in fabricated Ti alloys can still be observed. Firstly, it is challenging for industrial inert gas to completely eliminate $\mathrm{O}$ during its production process, whereby a portion of the $\mathrm{O}$ can enter the chamber along with the inert gas via the gas supply systems. Additionally, the sensitivity of $\mathrm{O}$ sensors varies, whereby the detection limit for $\mathrm{O}$ content is typically around 1000 ppm [243]. Moreover, apart from 0 , the inert gas itself is also a source of interstitial elements. According to Kornilov et al. [246], both $\mathrm{O}$ and $\mathrm{N}$ atoms have small enough sizes (van der Waals radii of 152 and $155 \mathrm{pm}$, severally) to fit into the voids of the Ti crystal structure. Therefore, it is challenging to fully avoid the additional interstitial elements occurring in Ti alloys fabricated by the L-PBF and DED methods. Whilst it is widely known that the production process of the EB-PBF system takes place in a vacuum chamber, $\mathrm{O}$ intrusion can still occur in $\mathrm{Ti}$ alloys fabricated by the EB-PBF method. Formanoir et al. [163] have reported $\mathrm{O}$ contamination in EB-PBF-fabricated Ti-6Al-4V, noting that some water moisture from the air can freeze on the walls of the building chamber during the evacuation of air to achieve a vacuum environment.

\subsection*{5.4.2. Evaporation of alloying elements}
During the AM process, the high temperature of melting and deposition processes can result in the evaporation of certain alloying elements, leading to the alloy composition deviation between the feedstock and produced parts [247,248]. This phenomenon occurs because different alloying elements have varying melting

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-22(1)}
\end{center}

Fig. 25. Relationship between the saturated vapor pressure and temperature for different elements (Reproduced with permission from Ref. [252]. Copyright (2020), MDPI).

points, with some elements being more volatile and susceptible to selective vaporization during AM processes [249]. In general, metallic elements with lower melting points, such as aluminium (Al), are more prone to evaporation and loss during the hightemperature melting process. Keaveney et al. [250] analyzed the condensate from the L-PBF building chamber after the production of Ti-6Al-4V alloy and indicated a higher concentration of Al compared to vanadium $(\mathrm{V})$ in the collected condensate. In addition, the amount of the alloying elements loss varies among different AM techniques. In general, compared to L-PBF, EB-PBF-fabricated Ti alloys experience a more significant loss of alloying elements. Brice et al. [251] have reported a decrease of $\sim 0.9 \mathrm{wt} . \% \mathrm{Al}$ element in Ti-6Al-4V alloy manufactured by the EB-PBF method, While Gaytan et al. [223] have shown a $\sim 0.6-1.0$ wt.\% reduction of the $\mathrm{Al}$ element in EB-PBF-fabricated Ti-Al-4V alloys. In contrast, the reduction of the Al element was $\sim 0.55$ wt.\% in L-PBF-built Ti-6Al-4V [99].

As illustrated in Fig. 25, the evaporation of alloying elements is affected by temperature, with higher temperatures leading to increased loss. A higher saturated vapor pressure can result in a larger loss of alloying element [252]. Moreover, the evaporation of alloying elements usually occurs on the molten pool surface during the AM process, whereby the ratio of the surface area to vol-

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-23}
\end{center}

Fig. 26. (a) EBSD maps for DED-fabricated Ti6Al4V-Mo alloys with different Ti6Al4V and Mo mixed ratio (wt.\%) at four interfaces: (b) measurements taken pre and post $\beta$-phase reconstruction at the $100 \% \mathrm{Ti6Al4V} / 75 \% \mathrm{Ti6Al4V}-25 \%$ Mo interface, while (c-e) $\beta$-phase with Mo increasing at three interfaces (Reproduced with permission from Ref. [257]. Copyright (2017), Elsevier); (f) and (g) BSEM graph and EDS elemental maps for Ti-35Nb fabricated via L-PBF adopting mixed and prealloyed powder (Reproduced with permission from Ref. [39]. Copyright (2022), Elsevier).

ume can greatly influence the level of alloying loss [28,253]. According to [62], the molten pool in EB-PBF is wider and shallower (width: $\sim 280 \pm 23 \mu \mathrm{m}$; depth: $\sim 152 \pm 15 \mu \mathrm{m}$ ) than the molten pool in L-PBF (width: $\sim 146 \pm 17 \mu \mathrm{m}$; depth: $\sim 172 \pm 21 \mu \mathrm{m}$ ). This indicates that the surface area-to-volume ratio of the molten pool is smaller during the production process of L-PBF. According to Klassen et al. [247], the vacuum environment of the EB-PBF building chamber can further promote the evaporation of alloying elements.

\subsection*{5.4.3. Inhomogeneity challenges}
Inhomogeneity challenges in the microstructure of AM-built Ti alloys have also been investigated. For instance, Wang et al. [39] have contrasted the melting behavior of the Ti-35Nb alloy manufactured by L-PBF, adopting a powder mixture and prealloyed powder. The backscattered SEM (BSEM) image and energy dispersive X-ray spectroscopy (EDS) mapping in Fig. 26(f) show that the Ti-35Nb sample produced from a powder mixture exhibited a heterogeneous microstructure with undissolved $\mathrm{Nb}$ particles of vary-\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-24}

Fig. 27. Influence of surface roughness (a) and porosity (b) on the fatigue properties of Ti-6Al-4V samples manufactured by L-PBF and EB-PBF (Reproduced with permission from Ref. [238]. Copyright (2018), Elsevier).

ing sizes. In contrast, the Ti-35Nb specimen from prealloyed powder (Fig. 26(g)) displayed a uniform and homogeneous microstructure with complete melting of Nb. Previous research has also highlighted the challenge of inhomogeneity caused by unmelted $\mathrm{Nb}$ particles in Ti alloys produced via L-PBF, resulting in both microstructural and chemical inhomogeneity [37,38,254,255]. Similarly, addressing the issue of inhomogeneity through parameter optimization in EB-PBF poses a significant challenge. The utilization of EB-PBF with blended elemental powders in the production of a Ti-10at.\%Nb alloy has revealed the formation of a heterogeneous microstructure [256]. Additionally, DED offers more advantages in producing functionally graded Ti alloys, including the ability to make in-situ adjustments to powder or mixture. However, it also poses challenges related to inhomogeneity, both in terms of chemical composition and microstructure. When examining DEDfabricated Ti-6Al-4V alloys with the addition of varying percentages of Mo content, two significant sources of inhomogeneity become apparent. Firstly, chemical inhomogeneity exists within each functionally graded layer due to insufficient energy to fully melt the high melting point of Mo (Fig. 26(a)). Secondly, microstructural inhomogeneity becomes more pronounced, particularly near the interfaces between gradient transitions along the build direction. This is evident in Fig. 26(b-e), where varying Mo contents result in different grain morphologies. For example, $100 \mathrm{wt} . \%$ Ti-6Al-4V forms coarse columnar $\beta$-grains, while the addition of $25 \mathrm{wt} \%$ Mo reduces the size of columnar $\beta$-grains without a specific crystallographic orientation. Further increasing the Mo content to $50 \mathrm{wt} . \%$ leads to the formation of fine equiaxed $\beta$-grains, and a Mo content of 75 wt.\% further decreases the equiaxed grain size to approximately 5-60 $\mu \mathrm{m}$. Thus, the challenge remains to optimize the melting process for powder mixtures, taking advantage of their cost-effectiveness and fast production while achieving better melting and homogeneity in the microstructure. Thus, continuous and future efforts will be necessary to enhance process control and mitigate challenges associated with interstitial element pickup, alloying element evaporation and inhomogeneity.

\subsection*{5.5. Mechanical properties affected by defects}
Understanding the impact of both macro and micro defects on the mechanical performance of AM-fabricated Ti alloys is crucial for optimizing their performance. As noted above, in AM-fabricated Ti alloys various defects can arise, including macro defects such as surface roughness, porosity and residual stress (RS), as well as micro defects resulting from changes in chemical composition and inhomogeneity issues. These defects, regardless of their scale, have the potential to impact the material's structural integrity and overall mechanical behavior.\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-25}

Fig. 28. Relationship between fatigue crack growth ( $\mathrm{d} a / \mathrm{d} n$ ) and stress intensity factor range (dK) in Ti-6Al-4V fabricated by (a) L-PBF with different conditions and (b) EB-PBF (Reproduced with permission from Ref. [262]. Copyright (2019), Elsevier).\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-25(1)}

Fig. 29. (a) Trend of hardness with increasing interstitial elements ( $O$ and N) in CP-Ti fabricated by L-PBF (Reproduced with permission from Ref. [244]. Copyright (2018), Elsevier) and (b) relationship between mechanical properties and the quantity of 0 element in L-PBF built Ti-6Al-4V (Reproduced with permission from Ref. [242]. Copyright (2020), Elsevier).

\subsection*{5.5.1. Mechanical properties related to macro defects}
It has been reported that macro defects can greatly affect the fatigue property of AM-fabricated Ti alloys [258]. Moreover, variations in production processes among different AM techniques can result in distinct fatigue properties of Ti alloys. Table 8 provides the fatigue properties of Ti alloys manufactured by different AM techniques, indicating that Ti-6Al-4V alloys manufactured by L-PBF and DED methods exhibit higher fatigue strength compared to EBPBF-fabricated samples. In contrast, the EB-PBF-built Ti-6Al-4V alloy presents higher fatigue toughness than those of Ti-6Al-4V alloys manufactured by L-PBF and DED.

The fatigue strength of Ti alloys manufactured by different AM techniques can be attributed to two key aspects related to crack initiation behavior. Firstly, Ti-6Al-4V alloys manufactured by L-PBF

Table 8

Fatigue properties of Ti-6Al-4V alloys fabricated by different AM techniques.

\begin{center}
\begin{tabular}{llllll}
\hline
Producing methods & Condition & $R$ & $\Delta \sigma_{\mathrm{w}}$ & $\Delta K_{\text {th }}$ & Refs. \\
\hline
L-PBF & As-built & 0.1 & 550 & 1.7 & $[101,259]$ \\
L-PBF & As-built & 0.1 & $220 \pm 24$ & - & $[231]$ \\
EB-PBF & As-built & 0.1 & $200-250$ & 3.8 & $[162,239]$ \\
EB-PBF & As-built & 0.1 & $115 \pm 13$ & - & $[231]$ \\
DED & As-built, machined & 0.1 & $482 *$ & $2.8-3.5$ & $[105,112]$ \\
\hline
\end{tabular}
\end{center}

$R$ - fatigue stress ratio; $\Delta \sigma_{\mathrm{w}}$ - threshold stress; $\Delta K_{\mathrm{th}}$ - crack propagation threshold; “*"- at the condition of $1 \times 10^{6}$ cycles.

\begin{center}
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(c)

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(d)

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\end{center}

Fig. 30. (a) Schematic diagram of interstitial $O$ elements inside the octahedral and tetrahedral lattices of Ti-35Ta alloy with a bcc lattice structure and (b) the change trend of lattice constant with the $O$ increase (Reproduced with permission from Ref. [272]. Copyright (2021), Elsevier); (c) HAADF-STEM graph of a dislocation impeded by the interstitial $\mathrm{O}$ array (top-left inset presents the geographic phase of the strain for the selected area; the $\mathrm{O}$ interstitial array is showed by iDPC-STEM image at the bottom-right inset) and (d) enlarged HAADF-STEM image illustrates the dislocation movement is impeded by a large amount of interstitial O element (Reproduced with permission from Ref. [265]. Copyright (2023), Nature Portfolio).

and DED methods contain the presence of martensite $\alpha$ ' phases and finer $\alpha / \alpha^{\prime}$ grains. During fatigue tests, cyclic load leads to the formation of plastic slip localization on the surface of Ti-6Al-4V samples, which serves as the starting point for fatigue crack initiation. The presence of martensite $\alpha^{\prime}$ phases and finer $\alpha / \alpha^{\prime}$ grains can effectively hinder the dislocation slip and delay the formation of the plastic slip localization on the sample surface, resulting in improved fatigue strength of Ti-6Al-4V alloy fabricated by L-PBF and DED methods [101,105]. In addition, surface roughness and porosity are also important factors for the fatigue strength. These defects create local stress concentrations that serve as sites for fatigue crack initiation sites during the fatigue test. Chastand et al. [238] have explored the influences of surface roughness and porosity on fatigue strength. Their findings, as shown in Fig. 27(a) and (b), suggest that a better surface finish (after polishing) can increase fatigue strength by $\sim 100 \%$ for $10^{7}$ cycles compared to asbuilt ones. Additionally, when internal porosity is further reduced through hot isostatic pressing (HIP) and followed by polishing, a significant $\sim 80 \%$ increase in fatigue strength is observed compared to polished-only counterparts. This highlights the substantial impact of both surface roughness and internal porosity on the fatigue strength of Ti alloys produced through AM. Therefore, achieving a high-quality surface finish and minimizing internal porosity are essential through process optimization or post-processing techniques to boost fatigue strength and ensure the reliable performance of AM-fabricated Ti alloys.

Contrasted to fatigue strength, fatigue toughness is almost not influenced at all by surface roughness. This can be attributed to the close relationship between fatigue toughness and fatigue crack growth (FCG) behavior, which is significantly affected by RS. According to Leuders et al. [259], RS is the main factor for the distinctions of FCG in L-PBF-fabricated Ti alloys with and without heat treatment. This is because the effective stress intensity factor ratio can be raised by tensile residual stress in the materials to increase\\
\includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-27}

Fig. 31. SEM images of the tensile fracture surfaces at the cross section for L-PBF fabricated Ti-35Nb: (a, b) specimen using prealloyed powder and (c, d) specimen using an elemental powder mixture (Reproduced with permission from Ref. [39]. Copyright (2022), Elsevier); (e) backscattered SEM for the surface deformation during a compressive test at a strain of $15 \%$ using an elemental powder mixture (Reproduced with permission from Ref. [37]. Copyright (2019), Elsevier), and (f) nanoindentation load-displacement curves for inhomogeneous regions (Reproduced with permission from Ref. [38]. Copyright (2021), Elsevier).

the FCG rate [260]. As reflected in Fig. 28(a), the machined L-PBFfabricated Ti-6Al-4V alloys present similar FCG rates to the as-built counterparts. However, as seen in Fig. 28(b), after the heat treatment, the crack growth resistance of L-PBF-fabricated Ti-6Al-4V alloys has been effectively improved due to the decreased level of RS. In addition, Cain et al. [261] also have indicated that FCG rate can be effectively decreased in the L-PBF-fabricated Ti-6Al-4V alloy after low-temperature stress relief and annealing heat treatments. Similarly, Leuders et al. [259] have reported that the crack propagation threshold of Ti-6Al-4V alloys manufactured by L-PBF can be effectively improved following stress relief through heat treatment. This improvement can be ascribed to the reduction of RS in the heat-treated part, leading to decreased stress accumulation and improved crack growth resistance. Additionally, because of the special producing process, Ti alloys fabricated by the EB-PBF method obtain relatively smaller RS. Therefore, EB-PBF-fabricated Ti alloys exhibit relatively higher fatigue toughness than the Ti alloys manufactured by L-PBF and DED methods. This can be confirmed in Fig. 28(b), where compared with the L-PBF-fabricated Ti-6Al-4V alloys, the EB-PBF-fabricated Ti-6Al-4V alloys present a lower FCG rate.

\subsection*{5.5.2. Mechanical properties related to micro defects}
The influence of chemical composition changes on the mechanical performances of produced parts indicates the importance of ad- dressing this aspect. As reflected in Fig. 23(b) previously, VelascoCastro et al. [243] have reported an improvement in hardness with increasing 0 content in the L-PBF-built Ti-6Al-4V alloy. Moreover, as shown in Fig. 29(a), Na et al. [244] also found that the hardness can be enhanced with the increasing concentration of $\mathrm{O}$ and $\mathrm{N}$ in the CP-Ti alloy fabricated by L-PBF. Apart from hardness, changes in chemical composition also have an influence on properties related to overall plastic deformation in AM-built Ti alloys. In order to exclude the influence of RS and porosity, Dietrich et al. [242] have investigated the mechanical properties of L-PBF-built Ti-6Al-4V alloys after heat treatments (stress relief) and HIP. Fig. 29(b) shows that an increase in $\mathrm{O}$ concentration resulted in the improvement of both YS and UTS of Ti-6Al-4V samples, while the elongation decreased. Furthermore, Carroll et al. [142] have found that DED-built Ti-6Al-4V samples with larger $O$ content exhibited both higher YS and UTS, but their ductility was $1.6 \%$ lower than the samples compared to samples with less $\mathrm{O}$ content. These results agree well with a study by Wang et al. [39], which revealed that the Ti-35Nb alloy fabricated via L-PBF with an O content of $0.30 \mathrm{wt} . \%$ exhibited a higher tensile YS of $636 \pm 80 \mathrm{MPa}$. However, it also exhibited a significantly lower ductility compared to the counterpart with an O of 0.12 wt.\%, which had a YS of $485 \pm 28 \mathrm{MPa}$.

The strengthening mechanism of interstitial elements in Ti alloys is similar to the solid solution strengthening. According to Kornilov et al. [246], $\mathrm{O}$ and $\mathrm{N}$ elements can easily occupy voids\\
within the Ti crystal structure and form solutes in the alloys. Fig. 30(a) shows the interstitial 0 element occupying octahedral and tetrahedral holes in Ti-35Ta alloy with a bcc lattice structure, where the increasing $\mathrm{O}$ content results in an increased lattice constant in the Ti-35Ta (Fig. 30(b)). The presence of interstitial solute atoms leads to lattice distortion, which impedes dislocations slip and low temperature twinning $[263,264]$. This phenomenon has been observed in research conducted by Ref. [265], where integrated differential phase contrast (iDPC)-scanning transmission electron microscopy (STEM) and high-angle annular darkfield (HAADF)-STEM technologies were used. Fig. 30(c) and (d) show that the dislocation movement was impeded in the area containing a large amount of the 0 interstitial array. Apart from the dislocation movement, according to Zaefferer [266], deformation twinning also can be totally inhibited in CP-Ti alloy with 2000 ppm of $O$ content. Therefore, both the hardness and strength can be improved by the increased interstitial elements in Ti alloys. However, this effect can come at the expense of reduced ductility. During the deformation, Ti alloys undergo a combination of twinning and slip mechanisms, which significantly influence the ductility of the material [267-271]. Therefore, with the increasing 0 content, AM-fabricated Ti alloys tend to present the relatively lower ductility.

Fig. 31 provides valuable insights into the inhomogeneity issues affecting the mechanical properties of Ti-35Nb alloy manufactured by L-PBF, thus allowing for an investigation into the impact of inhomogeneity on the mechanical properties of Ti alloys. Specifically, Fig. 31(a) shows the tensile fracture surfaces of L-PBFfabricated Ti-35Nb alloy with a homogeneous chemical distribution. The presence of elongated necking (Fig. 31(a)) and homogeneous ductile fine dimples (Fig. 31(b)) confirms higher tensile ductility $(23.5 \% \pm 2.2 \%)$. As described earlier, the inhomogeneity resulting from undissolved $\mathrm{Nb}$ particles in Ti-Nb alloys has a substantial impact on the mechanical properties. Further examination, as depicted in Fig. 31(c) and (d), reveals the presence of microcracks and unmelted powders. The fracture surface is dominated by transgranular fracture and smooth cleavage facets, leading to a low ductility of $2.2 \% \pm 1.4 \%$. In Fig. 31(e), the surface compressive deformation provides additional evidence of an inhomogeneous $\beta$ phase microstructure. The weaker interface between the $\beta$ matrix and undissolved $\mathrm{Nb}$ particles hinders shear band propagation near the interfaces. The nanoindentation load-displacement curves presented in Fig. 31(f) exhibit distinct response behaviors within the inhomogeneous microstructure of Ti-Nb $\beta$ phase, undissolved $\mathrm{Nb}$ particle, and their interface. The Ti-Nb $\beta$ phase displays the smallest penetration depth (displacement), indicating higher hardness and strength, whereas $\mathrm{Nb}$ particle exhibits the lowest hardness. This further confirms the weak bonding of the interfaces, which contributes to crack initiation and premature tensile and compressive failure. Improving homogeneity has the potential to significantly enhance ductility. These findings in Fig. 31 highlight the detrimental effects of inhomogeneity and emphasize the importance of addressing this issue in order to optimize the mechanical performance of Ti alloys.

\section*{6. Conclusions}
Additive manufacturing (AM) provides a complementary, rather than a replacement, approach to traditional manufacturing processes due to its enhanced flexibility in fabricating shape-complex parts and solving machining challenges, resulting in reduced lead times for custom designs. This article has systematically examined phase transformation, grain size and morphology, as well as defects, and discusses their impacts on the mechanical properties of Ti alloys manufactured by three commonly used powder-type AM techniques: laser powder bed fusion (L-PBF), electron beam powder bed fusion (EB-PBF) and directed energy deposition (DED). The differences in production processes lead to distinct microstructures and defects in Ti alloys fabricated by these AM techniques. The formation of non-equilibrium phases $\left(\alpha^{\prime}, \alpha^{\prime \prime}, \omega\right)$ usually can impede the dislocation movement, which can delay the deformation of Ti alloys. Therefore, the L-PBF-fabricated Ti alloys generally exhibit relatively higher strength and hardness compared with counterparts fabricated by DED and EB-PBF methods. However, nonequilibrium phases are unfavorable for the elastic modulus and ductility of Ti alloys.

Grain size and morphology also impact the mechanical properties of Ti alloys. Smaller grain size (more grain boundaries) can prevent further propagation of dislocations and decrease stress concentration, resulting in higher strength and hardness in Ti alloys fabricated by L-PBF and DED methods compared to EB-PBFfabricated counterparts. Additionally, the columnar to equiaxed transition (CET) can be potentially achieved in AM-built Ti alloys by optimizing and adjusting printing parameters to modify the ratio of $G / R$ (thermal gradient over solidification rate) or by promoting heterogeneous nucleation via the addition of refining elements. Compared to the columnar grains, the AM-built Ti alloys consisting of equiaxed grain can present both higher strength and ductility because of decreased stress concentration and grain size.

Despite the numerous benefits of AM, it remains challenging to produce Ti alloys without defects at the current stage. AMfabricated $\mathrm{Ti}$ alloys typically contain macro defects (e.g., porosity, surface roughness, residual stress) and micro defects (e.g., chemical composition changes and inhomogeneity issues). These macro defects, such as porosity and surface roughness, can obviously impact the fatigue property and overall performance of AM-fabricated Ti alloys by inducing fatigue crack initiation, while residual stress (RS) promotes crack growth and significantly influences fatigue toughness. In addition to defects, the introduction of oxygen $(\mathrm{O})$ and nitrogen $(\mathrm{N})$ atoms during the AM production process can alter the chemical composition of Ti alloys and result in lattice distortion, which impedes the dislocations slip and twinning, ultimately leading to improved strength but decreased ductility. Moreover, chemical inhomogeneity in Ti alloys can have a negative impact on their mechanical properties, as undissolved elements can induce microstructural inhomogeneity, alter the phase composition, and weaken bonding between the matrix and undissolved particles, ultimately leading to reduced ductility and strength.

This article provides a comprehensive understanding of the mechanical properties of $\mathrm{Ti}$ alloys manufactured using various $\mathrm{AM}$ techniques, covering both macro and micro perspectives. Although the mechanical properties of AM-fabricated Ti alloys are influenced by several factors, understanding the relationship between these properties and the AM techniques used remains a significant challenge. As such, this review highlights the need for further research to optimize processing parameters and microstructure design to achieve a more desirable balance of mechanical properties. In addition, efforts to develop novel materials and alloys tailored for AM processes will unlock new possibilities for creating advanced components with superior properties. To fully unlock the potential of emerging multi-laser AM in improving production efficiency and utilize the capabilities of multi-material AM to create high-quality multifunctional Ti alloys for diverse industrial applications, it is important to prioritize and conduct more research in those areas. Researchers also need to improve their understanding of the fundamental mechanisms governing the relationships between powder characteristics, processing conditions, underlying principles of AM techniques, microstructure evolution, and resultant mechanical properties.

\section*{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.

\section*{CRediT authorship contribution statement}
H.Y. Ma: Writing - original draft, Writing - review \& editing. J.C. Wang: Writing - original draft, Writing - review \& editing. P. Qin: Writing - review \& editing. Y.J. Liu: Writing - review \& editing. L.Y. Chen: Writing - review \& editing. L.Q. Wang: Writing review \& editing. L.C. Zhang: Conceptualization, Writing - original draft, Writing - review \& editing, Supervision.

\section*{Acknowledgements}
The authors would like to acknowledge the financial support provided by the industrial grant (No. G1006320). J.C Wang is grateful for the support of the Forrest Research Foundation PhD scholarship. The authors would like to thank the Australian Government Research Training Program Scholarship. The authors also acknowledge the facilities, and the scientific and technical assistance of the Australian Microscopy \& Microanalysis Research Facility at the centre for Microscopy, characterisation \& Analysis, The University of Western Australia, a facility funded by the University, State and Commonwealth Governments.

\section*{References}
[1] A. Sargeant, T. Goswami, Mater. Des. 27 (2006) 287-307.

[2] P. Majumdar, S.B. Singh, M. Chakraborty, J. Mech. Behav. Biomed. Mater. 4 (2011) 1132-1144

[3] Y.J. Liu, H.L. Wang, S.J. Li, S.G. Wang, W.J. Wang, W.T. Hou, Y.L. Hao, R. Yang, L.C. Zhang, Acta Mater. 126 (2017) 58-66

[4] M. Qian, W. Xu, M. Brandt, H.P. Tang, MRS Bull. 41 (2016) 775-784.

[5] X. Wang, P. Qin, L.Y. Chen, H. Sun, L.C. Zhang, Corros. Sci. 221 (2023) 111336.

[6] S.W. Ge, P. Hu, J. Deng, S.L. Li, H.R. Xing, J.Y. Han, X.J. Hua, L. Wang, J.Z. Yang, B. Jin, W. Zhang, K.S. Wang, Tungsten (2023), doi:10.1007/ s42864-023-00228-y.

[7] S.F. Jawed, Y.J. Liu, J.C. Wang, C.D. Rabadia, L.Q. Wang, Y.H. Li, X.H. Zhang, L.C. Zhang, J. Mech. Behav. Biomed. Mater. 110 (2020) 103867

[8] C.D. Rabadia, Y.J. Liu, C.H. Zhao, J.C. Wang, S.F. Jawed, L.Q. Wang, L.Y. Chen, H. Sun, L.C. Zhang, Mater. Sci. Eng. A 766 (2019) 138340.

[9] R. Blackman, N. Barghi, C. Tran, J. Prosthet. Dent. 65 (1991) 309-315.

[10] M. Taira, J.B. Moser, E.H. Greener, Dent. Mater. 5 (1989) 45-50.

[11] L.C. Zhang, J. Das, H.B. Lu, C. Duhamel, M. Calin, J. Eckert, Scr. Mater. 57 (2007) 101-104.

[12] M. Calin, L.C. Zhang, J. Eckert, Scr. Mater. 57 (2007) 1101-1104.

[13] L.C. Zhang, H.B. Lu, C. Mickel, J. Eckert, Appl. Phys. Lett. 91 (2007) 051906.

[14] J. Raval, A. Kazi, O. Randolph, X. Guo, R. Zvanut, C. Lee, B. Tai, J. Manuf. Process. 94 (2023) 539-549

[15] P.J. Kwok, S.M. Oppenheimer, D.C. Dunand, Adv. Eng. Mater. 10 (2008) 820-825.

[16] G.E. Ryan, A.S. Pandit, D.P. Apatsidis, Biomaterials 29 (2008) 3625-3635

[17] C.Q. Ning, Y. Zhou, Biomaterials 23 (2002) 2909-2915.

[18] A. Carman, L.C. Zhang, O.M. Ivasishin, D.G. Savvakin, M.V. Matviychuk, E.V. Pereloma, Mater. Sci. Eng. A 528 (2011) 1686-1693.

[19] Y.J. Chen, B. Feng, Y.P. Zhu, J. Weng, J.X. Wang, X. Lu, Mater. Lett. 63 (2009) 2659-2661.

[20] F.L. Yu, C. Kong, H.L. Yu, Tungsten 5 (2023) 522-530.

[21] J.C. Wang, R. Zhu, Y.J. Liu, L.C. Zhang, Adv. Powder Mater. 2 (2023) 100137.

[22] B. Baufeld, O.V.d. Biest, R. Gault, Mater. Des. 31 (2010) S106-S111.

[23] J. Karlsson, A. Snis, H. Engqvist, J. Lausmaa, J. Mater, Process. Technol. 213 (2013) 2109-2118.

[24] D.D. Gu, W. Meiners, K. Wissenbach, R. Poprawe, Int. Mater. Rev. 57 (2012) 133-164.

[25] J. Kranz, D. Herzog, C. Emmelmann, J. Laser Appl. 27 (2015) S14001.

[26] Y.J. Liu, J.S. Zhang, X.C. Liu, X. Wu, J.C. Wang, Y.S. Zhang, L.Q. Wang, L.C. Zhang, Mater. Sci. Eng. A 822 (2021) 141696.

[27] F.G. Caballero, Encyclopedia of Materials: Metals and Alloys, Elsevier, 2021.

[28] T. DebRoy, H.L. Wei, J.S. Zuback, T. Mukherjee, J.W. Elmer, J.O. Milewski, A.M. Beese, A.Wilson Heid, A. De, W. Zhang, Prog. Mater. Sci. 92 (2018) $112-224$.

[29] T.D. Ngo, A. Kashani, G. Imbalzano, K.T.Q. Nguyen, D. Hui, Compos. B-Eng. 143 (2018) 172-196.

[30] J.O. Milewski, Additive Manufacturing of Metals, Springer, 2017.\\
[31] D. Herzog, V. Seyda, E. Wycisk, C. Emmelmann, Acta Mater. 117 (2016) 371-392.

[32] K. Wei, F. Li, G. Huang, M. Liu, J. Deng, C. He, X. Zeng, Mater. Sci. Eng. A 802 (2021) 140644.

[33] S. Li, J. Yang, Z. Wang, Powder Technol. 384 (2021) 100-111.

[34] D. Wang, L. Liu, G. Deng, C. Deng, Y. Bai, Y. Yang, W. Wu, J. Chen, Y. Liu, Y. Wang, X. Lin, C. Han, Virtual Phys. Prototyp. 17 (2022) 329-365.

[35] A. Nazir, O. Gokcekaya, K. Md Masum Billah, O. Ertugrul, J. Jiang, J. Sun, S. Hussain, Mater. Des. 226 (2023) 111661

[36] C. Wei, L. Li, Virtual Phys. Prototyp. 16 (2021) 347-371.

[37] J.C. Wang, Y.J. Liu, P. Qin, S.X. Liang, T.B. Sercombe, L.C. Zhang, Mater. Sci. Eng. A 760 (2019) 214-224.

[38] J. Wang, Y. Liu, C.D. Rabadia, S.X. Liang, T.B. Sercombe, L.C. Zhang, J. Mater. Sci. Technol. 61 (2021) 221-233.

[39] J.C. Wang, Y.J. Liu, S.X. Liang, Y.S. Zhang, L.Q. Wang, T.B. Sercombe, L.C. Zhang, J. Mater. Sci. Technol. 105 (2022) 1-16

[40] N. Dai, L.C. Zhang, J. Zhang, Q. Chen, M. Wu, Corros. Sci. 102 (2016) 484489.

[41] W.Q. Lu, Y.J. Liu, X. Wu, X.C. Liu, J.C. Wang, Surf. Coat. Technol. 470 (2023) 129849.

[42] P. Qin, Y. Chen, Y. Liu, J. Zhang, L. Chen, Y. Li, X. Zhang, C. Cao, H. Sun, L.C. Zhang, ACS Biomater. Sci. Eng. 5 (2019) 1141-1149.

[43] L.C. Zhang, H. Attar, Adv. Eng. Mater. 18 (2016) 463-475.

[44] L.C. Zhang, L.Y. Chen, Adv. Eng. Mater. 21 (2019) 1801215

[45] L.C. Zhang, Y. Liu, S. Li, Y. Hao, Adv. Eng. Mater. 20 (2018) 1700842.

[46] L.C. Zhang, L.Y. Chen, S. Zhou, Z. Luo, J. Alloy. Compd. 936 (2023) 168099.

[47] L.Y. Chen, S.X. Liang, Y. Liu, L.C. Zhang, Mater. Sci. Eng. R-Rep. 146 (2021) 100648.

[48] T. Ron, A. Shirizly, E. Aghion, Materials (Basel) 16 (2023) 2454.

[49] S.K. Everton, M. Hirsch, P.I. Stavroulakis, R.K. Leach, A.T. Clare, Mater. Des. 95 (2016) 431-445

[50] W.S.W. Harun, M.S.I.N. Kamariah, N. Muhamad, S.A.C. Ghani, F. Ahmad, Z. Mohamed, Powder Technol. 327 (2018) 128-151.

[51] K.L. Li, J.H. Chen, C.C. Zhao, Z.J. Shen, W. Liu, Tungsten 3 (2021) 218-233.

[52] J. Parthasarathy, B. Starly, S. Raman, A. Christensen, J. Mech. Behav. Biomed. Mater. 3 (2010) 249-259.

[53] Y.J. Liu, Z. Liu, Y. Jiang, G.W. Wang, Y. Yang, L.C. Zhang, J. Alloy. Compd. 735 (2018) 1414-1421

[54] T.M. Mower, M.J. Long, Mater. Sci. Eng. A 651 (2016) 198-213.

[55] J.P.M. Pragana, R.F.V. Sampaio, I.M.F. Bragança, C.M.A. Silva, P.A.F. Martins, Adv. Ind. Manuf. Eng. 2 (2021) 100032.

[56] W.E. Frazier, J. Mater. Eng. Perform. 23 (2014) 1917-1928.

[57] B. Vandenbroucke, J.P. Kruth, Rapid Prototyp. J. 13 (2007) 196-203.

[58] L. Thijs, F. Verhaeghe, T. Craeghs, J.V. Humbeeck, J.P. Kruth, Acta Mater. 58 (2010) 3303-3312.

[59] J.H. Tan, W.L.E. Wong, K.W. Dalgarno, Addit. Manuf. 18 (2017) 228-255.

[60] L.Y. Chen, H.Y. Zhang, C. Zheng, H.Y. Yang, P. Qin, C. Zhao, S. Lu, S.X. Liang, L. Chai, L.C. Zhang, Mater. Des. 208 (2021) 109907.

[61] X.Y. Cheng, S.J. Li, L.E. Murr, Z.B. Zhang, Y.L. Hao, R. Yang, F. Medina, R.B. Wicker, J. Mech. Behav. Biomed. Mater. 16 (2012) 153-162.

[62] Y.J. Liu, S.J. Li, H.L. Wang, W.T. Hou, Y.L. Hao, R. Yang, T.B. Sercombe, L.C. Zhang, Acta Mater. 113 (2016) 56-67.

[63] M.T. Mohammed, V.G.S. Semelov, A.V.S. Sotov, Mater. Res. Express. 6 (2019) 122006.

[64] D.S. Shim, G.Y. Baek, J.S. Seo, G.Y. Shin, K.P. Kim, K.Y. Lee, Opt. Laser Technol. 86 (2016) 69-78

[65] W.J. Sames, F.A. List, S. Pannala, R.R. Dehoff, S.S. Babu, Int. Mater. Rev. 61 (2016) 315-360.

[66] N. Sanaei, A. Fatemi, Prog. Mater. Sci. 117 (2021) 100724.

[67] A. Saboori, D. Gallo, S. Biamino, P. Fino, M. Lombardi, Appl. Sci. 7 (2017) 883.

[68] S.H. Pan, G.C. Yao, Y.N. Cui, F.S. Meng, C. Luo, T.Q. Zheng, G. Singh, Tungsten 5 (2023) 1-31.

[69] K.N. Amato, S.M. Gaytan, L.E. Murr, E. Martinez, P.W. Shindo, J. Hernandez, S. Collins, F. Medina, Acta Mater. 60 (2012) 2229-2239.

[70] H.D. Nguyen, A. Pramanik, A.K. Basak, Y. Dong, C. Prakash, S. Debnath, S. Shankar, I.S. Jawahir, S. Dixit, D. Buddhi, J. Mater. Res. Technol. 18 (2022) 4641-4661

[71] Y. Li, D. Gu, Addit. Manuf. 1-4 (2014) 99-109.

[72] B. Wysocki, P. Maj, R. Sitek, J. Buhagiar, K.J. Kurzydłowski, W. Święszkowski, Appl. Sci. 7 (2017) 657.

[73] H. Galarraga, R.J. Warren, D.A. Lados, R.R. Dehoff, M.M. Kirka, P. Nandwana, Mater. Sci. Eng. A 685 (2017) 417-428.

[74] P.A. Kobryn, E.H. Moore, S.L. Semiatin, Scr. Mater. 43 (2000) 299-305.

[75] S. Liu, Y.C. Shin, Mater. Des. 164 (2019) 107552.

[76] S. Ehtemam-Haghighi, Y. Liu, G. Cao, L.C. Zhang, Mater. Sci. Eng. C 60 (2016) 503-510.

[77] S.E. Haghighi, H.B. Lu, G.Y. Jian, G.H. Cao, D. Habibi, L.C. Zhang, Mater. Des. 76 (2015) 47-54.

[78] C. Veiga, J.P. Davim, A.J.R. Loureiro, Rev. Adv. Mater. Sci. 32 (2012) 133-148.

[79] M. Geetha, A.K. Singh, R. Asokamani, A.K. Gogia, Prog. Mater. Sci. 54 (2009) 397-425.

[80] A.T. Sidambe, Materials 7 (2014) 8168-8188 (Basel).

[81] Y. Oshida, Bioscience and Bioengineering of Titanium Materials, Elsevier 2007.

[82] C. Leyens, M. Peters, Titanium and Titanium Alloys: Fundamentals and Applications, Wiley, 2003.

[83] V.A. Joshi, Titanium Alloys: An Atlas of Structures and Fracture Features, CRC Press, 2006.

[84] A.R. McAndrew, P.A. Colegrove, C. Bühr, B.C.D. Flipo, A. Vairis, Prog. Mater. Sci. 92 (2018) 225-257.

[85] J. Cohen, Am. J. Surg. 114 (1967) 31-41.

[86] C.N. Elias, D.J. Fernandes, C.R.S. Resende, J. Roestel, Dent. Mater. 31 (2015) e1-e13.

[87] F. Weng, C. Chen, H. Yu, Mater. Des. 58 (2014) 412-425.

[88] A.K. Srivastava, Mol. Cell. Biochem. 206 182

(2000) 177-

89] R. Azadbakht, T. Almasi, H. Keypour, M. Rezaeivala, Inorg. Chem. Commun. 33 (2013) 63-67.

[90] S.J. Li, L.E. Murr, X.Y. Cheng, Z.B. Zhang, Y.L. Hao, R. Yang, F. Medina, R.B. Wicker, Acta Mater 60 (2012) 793-802.

[91] S.F. Jawed, C.D. Rabadia, Y.J. Liu, L.Q. Wang, Y.H. Li, X.H. Zhang, L.C. Zhang, Mater. Des. 181 (2019) 108064.

[92] L.C. Zhang, L.Y. Chen, L. Wang, Adv. Eng. Mater. 22 (2020) 1901258

[93] M.S. Oh, J.Y. Lee, J.K. Park, Metall. Mater. Trans. A 35 (2004) 3071-3077.

[94] H. Attar, S. Ehtemam-Haghighi, D. Kent, X. Wu, M.S. Dargusch, Mater. Sci. Eng. A 705 (2017) 385-393.

[95] S.L. Sing, W.Y. Yeong, F.E. Wiria, J. Alloy. Compd. 660 (2016) 461-470.

[96] K. Yamanaka, W. Saito, M. Mori, H. Matsumoto, A. Chiba, Addit. Manuf. 8 (2015) 105-109.

[97] Ó. Barro, F. Arias-González, F. Lusquiños, R. Comesaña, J. del Val, A. Riveiro, A. Badaoui, F. Gómez-Baño, J. Pou, Metals 11 (2021) 70.

[98] T. Ahmed, H.J. Rack, Mater. Sci. Eng. A 243 (1998) 206-211.

[99] X. Zhao, S. Li, M. Zhang, Y. Liu, T.B. Sercombe, S. Wang, Y. Hao, R. Yang, L.E. Murr, Mater. Des. 95 (2016) 21-31.

[100] B. Vrancken, L. Thijs, J.P. Kruth, J. Van Humbeeck, J. Alloy. Compd. 541 (2012) $177-185$.

[101] H.K. Rafi, N.V. Karthik, H. Gong, T.L. Starr, B.E. Stucker, J. Mater. Eng. Perform. 22 (2013) 3872-3883.

[102] T. Vilaro, C. Colin, J.D. Bartout, Metall. Mater. Trans. A 42 (2011) 3190-3199.

[103] G. Kasperovich, J. Hausmann, J. Mater. Process. Technol. 220 (2015) 202-214

[104] X.P. Ren, H.Q. Li, H. Guo, F.L. Shen, C.X. Qin, E.T. Zhao, X.Y. Fang, Mater. Sci. Eng. A 817 (2021) 141384.

[105] Y. Zhai, H. Galarraga, D.A. Lados, Eng. Fail. Anal. 69 (2016) 3-14.

[106] M. Rashid, R.L. Narayan, D.L. Zhang, W.Z. Han, J. Mater. Eng. Perform. 31 (2022) 542-551

[107] M. Neikter, P. Åkerfeldt, R. Pederson, M.L. Antti, V. Sandell, Mater. Charact. 143 (2018) 68-75.

[108] S.S. Al-Bermani, M.L. Blackmore, W. Zhang, I. Todd, Metall. Mater. Trans. A 41 (2010) 3422-3434

[109] L.E. Murr, E.V. Esquivel, S.A. Quinones, S.M. Gaytan, M.I. Lopez, E.Y. Martinez, F. Medina, D.H. Hernandez, E. Martinez, J.L. Martinez, S.W. Stafford, D.K. Brown, T. Hoppe, W. Meyers, U. Lindhe, R.B. Wicker, Mater. Charact. 60 (2009) 96-105.

[110] J. Su, H. Xie, C. Tan, Z. Xu, J. Liu, F. Jiang, J. Tang, D. Fu, H. Zhang, J. Teng, Tribol. Int. 175 (2022) 107763

[111] Y.L. Hao, S.J. Li, R. Yang, Rare Met. 35 (2016) 661-671

[112] S.M.J. Razavi, F. Berto, Adv. Eng. Mater. 21 (2019) 1900220.

[113] W. Xu, E.W. Lui, A. Pateras, M. Qian, M. Brandt, Acta Mater. 125 (2017) 390-400.

[114] S.L. Lu, H.P. Tang, Y.P. Ning, N. Liu, D.H. StJohn, M. Qian, Metall. Mater. Trans. A 46 (2015) 3824-3834

[115] Y.J. Liu, Y.S. Zhang, L.C. Zhang, Materialia 6 (2019) 100299.

[116] X. Luo, L.H. Liu, C. Yang, H.Z. Lu, H.W. Ma, Z. Wang, D.D. Li, L.C. Zhang, Y.Y. Li, J. Mater. Sci. Technol. 68 (2021) 112-123.

[117] N. Hafeez, S. Liu, E. Lu, L. Wang, R. Liu, W. lu, L.C. Zhang, J. Alloy. Compd. 790 (2019) 117-126.

[118] N. Hafeez, J. Liu, L. Wang, D. Wei, Y. Tang, W. Lu, L.C. Zhang, Addit. Manuf. 34 (2020) 101264.

[119] N. Hafeez, D. Wei, L. Xie, Y. Tang, J. Liu, H. Kato, W. Lu, L.C. Zhang, L. Wang, Addit. Manuf. 48 (2021) 102376 .

[120] L.C. Zhang, D. Klemm, J. Eckert, Y.L. Hao, T.B. Sercombe, Scr. Mater. 65 (2011) 21-24.

[121] J. Hernandez, S.J. Li, E. Martinez, L.E. Murr, X.M. Pan, K.N. Amato, X.Y. Cheng, F. Yang, C.A. Terrazas, S.M. Gaytan, Y.L. Hao, R. Yang, F. Medina, R.B. Wicker, J. Mater. Sci. Technol. 29 (2013) 1011-1017.

[122] Y. Liu, S. Li, W. Hou, S. Wang, Y. Hao, R. Yang, T.B. Sercombe, L.C. Zhang, J. Mater. Sci. Technol. 32 (2016) 505-508

[123] R.L. Batalha, W.C. Batalha, L. Deng, T. Gustmann, S. Pauly, C.S. Kiminami, P. Gargarella, J. Mater. Res. 35 (2020) 1143-1153.

[124] R. Ummethala, P.S. Karamched, S. Rathinavelu, N. Singh, A. Aggarwal, K. Sun, E. Ivanov, L. Kollo, I. Okulov, J. Eckert, K.G. Prashanth, Materialia 14 (2020) 100941.

[125] N. Nadammal, M. Rajput, S.K. Gupta, E. Ivanov, A.S. Reddy, S. Suwas, K. Chatterjee, ACS Omega 7 (2022) 8506-8517.

[126] X. Luo, C. Yang, R.Y. Li, H. Wang, H.Z. Lu, T. Song, H.W. Ma, D.D. Li, A. Gebert, Y.Y. Li, Biomater. Adv. 133 (2022) 112625

[127] R. Banerjee, S. Nag, S. Samuel, H.L. Fraser, J. Biomed. Mater. Res. A 78 (2006) 298-305.

[128] M.S.K.K.Y. Nartu, D. Flannery, S. Mazumder, S.A. Mantri, S.S. Joshi, A.V. Ayyagari, B. McWilliams, K. Cho, N.B. Dahotre, R. Banerjee, JOM 73 (2021) 1819-1827.\\
[129] D. Gu, Y.C. Hagedorn, W. Meiners, G. Meng, R.J.S. Batista, K. Wissenbach, R. Poprawe, Acta Mater. 60 (2012) 3849-3860.

[130] H. Attar, M. Calin, L.C. Zhang, S. Scudino, J. Eckert, Mater. Sci. Eng. A 593 (2014) 170-177

[131] S. Banerjee, P. Mukhopadhyay, Phase Transformations: Examples From Titanium and Zirconium Alloys, Elsevier, 2010

[132] J. Liu, Y. Jin, X. Fang, C. Chen, Q. Feng, X. Liu, Y. Chen, T. Suo, F. Zhao, T. Huang, H. Wang, X. Wang, Y. Fang, Y. Wei, L. Meng, J. Lu, W. Yang, Sci. Rep. 6 (2016) 35345.

[133] S. Guo, Q. Meng, G. Liao, L. Hu, X. Zhao, Prog. Nat. Sci. 23 (2013) 174-182.

[134] K. Li, J. Yang, Y. Yi, X. Liu, Y. Liu, L.C. Zhang, W. Zhang, W. Li, D. Chen, S. Zhou, Acta Mater. 256 (2023) 119112.

[135] Z. Fan, H. Feng, Results Phys. 10 (2018) 660-664.

[136] L. Facchini, E. Magalini, P. Robotti, A. Molinari, Rapid Prototyp. J. 15 (2009) 171-178.

[137] M. Koike, K. Martinez, L. Guo, G. Chahine, R. Kovacevic, T. Okabe, J. Mater, Process. Technol. 211 (2011) 1400-1408.

[138] X. Tan, Y. Kok, Y.J. Tan, M. Descoins, D. Mangelinck, S.B. Tor, K.F. Leong, C.K. Chua, Acta Mater. 97 (2015) 1-16

[139] J. Yu, M. Rombouts, G. Maes, F. Motmans, Phys. Proc. 39 (2012) 416-424.

[140] J.S. Keist, T.A. Palmer, Mater. Des. 106 (2016) 482-494.

[141] C. Qiu, G.A. Ravi, C. Dance, A. Ranson, S. Dilworth, M.M. Attallah, J. Alloy. Compd. 629 (2015) 351-361.

[142] B.E. Carroll, T.A. Palmer, A.M. Beese, Acta Mater. 87 (2015) 309-320.

[143] M. Hein, N.F. Lopes Dias, S. Pramanik, D. Stangier, K.P. Hoyer, W. Tillmann, M. Schaper, Materials 15 (2022) 3774 (Basel).

[144] N. Stanford, P.S. Bate, Acta Mater. 52 (2004) 5215-5224.

[145] M. Simonelli, Y.Y. Tse, C. Tuck, Metall. Mater. Trans. A 45 (2014) 2863-2872.

[146] R. Lancaster, G. Davies, H. Illsley, S. Jeffs, G. Baxter, Materials 9 (2016) 470 (Basel).

[147] L. Yang, S. Ramachandran, A. Bagasol, Q. Guan, W. Wang, D.J. Browne, D. Dowling, W. Mirihanage, Scr. Mater. 231 (2023) 115430.

[148] J. Ma, J. Tian, M. Yan, Z. Chen, J. Shen, J. Wu, Mater. Charact. 166 (2020) 110440.

[149] S. Nayir, J.S. Keist, T.A. Palmer, Mater. Sci. Eng. A 822 (2021) 141541.

[150] W. Fan, H. Tan, F. Zhang, Z. Feng, Y. Wang, L.C. Zhang, X. Lin, W. Huang, J. Mater. Sci. Technol. 94 (2021) 32-46.

[151] X. Wu, J. Liang, J. Mei, C. Mitchell, P.S. Goodwin, W. Voice, Mater. Des. 25 (2004) 137-144.

[152] A. Carrozza, A. Aversa, F. Mazzucato, E. Bassini, D. Manfredi, S. Biamino, A. Valente, P. Fino, Mater. Charact. 189 (2022) 111958

[153] M.V. Pantawane, Y.H. Ho, S.S. Joshi, N.B. Dahotre, Sci. Rep. 10 (2020) 7579.

[154] A. Gupta, C.J. Bennett, W. Sun, Eng. Fail. Anal. 120 (2021) 105115.

[155] T. Song, T. Dong, S.L. Lu, K. Kondoh, R. Das, M. Brandt, M. Qian, Addit. Manuf. 46 (2021) 102139.

[156] H.Chandra Kaushik, S.I. Shakil, B.S. Amirkhiz, M. Mohammadi, E. Asadi, M. Haghshenas, A. Hadadzadeh, Vacuum 195 (2022) 110690.

[157] A. Hadadzadeh, E. Asadi, S. Imam Shakil, B. Shalchi Amirkhiz, M. Mohammadi, M. Haghshenas, Mater. Lett 301 (2021) 130273.

[158] J. Lee, K. Kim, J. Choi, J.G. Kim, S. Kim, J. Mater. Res. Technol. 23 (2023) 4499-4512.

[159] M. Gwak, S. Kim, D.J. Lee, J.B. Seol, H. Sung, T.H. Nam, S. Kim, J.G. Kim, Mater Sci. Eng. A 836 (2022) 142729.

[160] X. Shi, C. Yan, W. Feng, Y. Zhang, Z. Leng, Opt. Laser Technol. 132 (2020) 106471.

[161] Y. Yang, Y.J. liu, J. Chen, H.L. Wang, Z.Q. Zhang, Y.J. Lu, S.Q. Wu, J.X. Lin, Mater Sci. Eng. A 707 (2017) 548-558,

[162] M. Seifi, A. Salem, D. Satko, J. Shaffer, J.J. Lewandowski, Int. J. Fatigue 94 (2017) 263-287.

[163] C. Formanoir, S. Michotte, O. Rigo, L. Germain, S. Godet, Mater. Sci. Eng. A 652 (2016) 105-119

[164] L. Ednie, R.J. Lancaster, A.A. Antonysamy, F. Zelenka, A. Scarpellini, L. Parimi, R. Maddalena, N.C. Barnard, P. Efthymiadis, Mater. Sci. Eng. A 857 (2022) 144050.

[165] J.S. Keist, S. Nayir, T.A. Palmer, Mater. Sci. Eng. A 787 (2020) 139454.

[166] H.R. Sandgren, Y. Zhai, D.A. Lados, P.A. Shade, J.C. Schuren, M.A. Groeber P. Kenesei, A.G. Gavras, Addit. Manuf. 12 (2016) 132-141.

[167] S.H. Whang, Nanostructured Metals and Alloys, Woodhead Publishing, 2011.

[168] R.E. Reed-Hill, R. Abbaschian, R. Abbaschian, Physical Metallurgy Principles, Van Nostrand, 1973

[169] L. Xiao, W. Song, M. Hu, P. Li, Mater. Sci. Eng. A 764 (2019) 138204.

[170] E. Cerri, E. Ghio, G. Bolelli, Mater. Sci. Eng. A 851 (2022) 143635.

[171] H. Sharma, D. Parfitt, A.K. Syed, D. Wimpenny, E. Muzangaza, G. Baxter, B. Chen, Mater. Sci. Eng. A 744 (2019) 182-194.

[172] R.H.J. Hannink, A.J. Hill, Nanoengi. Metall. Mater. (2006).

[173] T. Wang, Y.Y. Zhu, S.Q. Zhang, H.B. Tang, H.M. Wang, J. Alloy. Compd. 632 (2015) 505-513.

[174] S. Kou, Welding Metallurgy, Wiley, 2003

[175] J.C. Lippold, S.D. Kiser, J.N. DuPont, Welding Metallurgy and Weldability of Nickel-Base Alloys, Wiley, 2011.

[176] S.A. David, J.M. Vitek, Int. Mater. Rev. 34 (1989) 213-245.

[177] M. Liu, S. Liu, W. Chen, C. Chen, Y. Lv, X. Zhang, P. Lei, Y. Lin, K. Zhou, Addit. Manuf. 30 (2019) 100873.

[178] C.J. Todaro, M.A. Easton, D. Qiu, D. Zhang, M.J. Bermingham, E.W. Lui, M. Brandt, D.H. StJohn, M. Qian, Nat. Commun. 11 (2020) 142.

[179] L. Lizzul, M. Sorgato, R. Bertolini, A. Ghiotti, S. Bruschi, Precis. Eng. 67 (2021) 301-310.

[180] C. Dharmendra, A. Hadadzadeh, B.S. Amirkhiz, A. Lloyd, M. Mohammadi, Mater. Sci. Eng. A 771 (2020) 138652.

[181] Q. Chao, S. Mateti, M. Annasamy, M. Imran, J. Joseph, Q. Cai, L.H. Li, P. Cizek, P.D. Hodgson, Y. Chen, D. Fabijanic, W. Xu, Addit. Manuf. 46 (2021) 102173.

[182] C. Liu, Y. Lu, X. Tian, D. Liu, Mater. Sci. Eng. A 661 (2016) 145-151.

[183] Y. Kok, X.P. Tan, P. Wang, M.L.S. Nai, N.H. Loh, E. Liu, S.B. Tor, Mater. Des. 139 (2018) 565-586.

[184] M. Simonelli, Y.Y. Tse, C. Tuck, Mater. Sci. Eng. A 616 (2014) 1-11.

[185] C. Qiu, N.J.E. Adkins, M.M. Attallah, Mater. Sci. Eng. A 578 (2013) 230-239

[186] L. Lan, X. Jin, S. Gao, B. He, Y. Rong, J. Mater. Sci. Technol. 50 (2020) 153-161.

[187] S. Rzepa, Z. Trojanová, J. Džugan, R.Z. Valiev, M. Koukolíková, D. Melzer, M. Brázda, Mater. Charact. 196 (2023) 112622.

[188] P. Åkerfeldt, M.L. Antti, R. Pederson, Mater. Sci. Eng. A 674 (2016) 428-437.

[189] W. Kurz, B. Giovanola, R. Trivedi, Acta Metall. 34 (1986) 823-830.

[190] S. Bontha, N.W. Klingbeil, P.A. Kobryn, H.L. Fraser, Mater. Sci. Eng. A 513-514 (2009) 311-318.

[191] G.J. Marshall, W.J. Young, S.M. Thompson, N. Shamsaei, S.R. Daniewicz, S. Shao, JOM 68 (2016) 778-790.

[192] Q. Zhang, J. Chen, X. Lin, H. Tan, W.D. Huang, J. Mater. Process. Technol. 238 (2016) 202-211.

[193] X. Wu, R. Sharman, J. Mei, W. Voice, Mater. Des. 25 (2004) 103-109.

[194] A.A. Antonysamy, J. Meyer, P.B. Prangnell, Mater. Charact. 84 (2013) 153-168.

[195] G.I. Eskin, D.G. Eskin, Ultrasonic Treatment of Light Alloy Melts, CRC Press, 2014.

[196] O.V. Abramov, Ultrasound in Liquid and Solid Metals, CRC Press, 1994.

[197] A. Kozlov, R. Schmid-Fetzer, IOP Conf. Ser.: Mater. Sci. Eng. 27 (2012) 012001.

[198] D.H. StJohn, S.D. McDonald, M.J. Bermingham, S. Mereddy, A. Prasad, M. Dargusch, Key Eng. Mater. 770 (2018) 155-164.

[199] D. Zhang, D. Qiu, M.A. Gibson, Y. Zheng, H.L. Fraser, D.H. StJohn, M.A. Easton, Nature 576 (2019) 91-95.

[200] G. Choi, W.S. Choi, J. Han, P.P. Choi, Addit. Manuf. 36 (2020) 101467.

[201] M.H. Mosallanejad, B. Niroumand, C. Ghibaudo, S. Biamino, A. Salmi, P. Fino, A. Saboori, Addit. Manuf. 56 (2022) 102878.

[202] R. Morgan, C.J. Sutcliffe, W. O’Neill, J. Mater. Sci. 39 (2004) 1195-1205.

[203] Q. Jia, D. Gu, J. Alloy. Compd. 585 (2014) 713-721.

[204] C. Qiu, S. Yue, N.J.E. Adkins, M. Ward, H. Hassanin, P.D. Lee, P.J. Withers, M.M. Attallah, Mater. Sci. Eng. A 628 (2015) 188-197.

[205] M. Peters, E.G. Brodie, S. Thomas, L. Djumas, M. Brameld, M. Salasi, Z. Quadir, M. Iannuzzi, J. Wang, T. Sercombe, C. Hutchinson, Materialia 32 (2023) 101958.

[206] S. Tammas-Williams, P.J. Withers, I. Todd, P.B. Prangnell, Scr. Mater. 122 (2016) 72-76.

[207] H. Gong, K. Rafi, H. Gu, T. Starr, B. Stucker, Addit. Manuf. 1-4 (2014) 87-98.

[208] Z. Snow, A.R. Nassar, E.W. Reutzel, Addit. Manuf. 36 (2020) 101457.

[209] S. Pal, N. Gubeljak, R. Hudák, G. Lojen, V. Rajtúková, T. Brajlih, I. Drstvenšek, Results Phys. 17 (2020) 103186.

[210] M.A. Buhairi, F.M. Foudzi, F.I. Jamhari, A.B. Sulong, N.A.M. Radzuan, N. Muhamad, I.F. Mohamed, A.H. Azman, W.S.W. Harun, M.S.H. Al-Furjan, Prog. Addit. Manuf. 8 (2023) 265-283.

[211] C. Qiu, C. Panwisawas, M. Ward, H.C. Basoalto, J.W. Brooks, M.M. Attallah, Acta Mater. 96 (2015) 72-79.

[212] B. Wysocki, P. Maj, A. Krawczyńska, K. Rożniatowski, J. Zdunek, K.J. Kurzydłowski, W. Święszkowski, J. Mater. Process. Technol. 241 (2017) 13-23.

[213] J. Jin, S. Zhou, W. Zhang, K. Li, Y. Liu, D. Chen, L.C. Zhang, J. Alloy. Compd. 918 (2022) 165704.

[214] Y.N. Hu, S.C. Wu, P.J. Withers, J. Zhang, H.Y.X. Bao, Y.N. Fu, G.Z. Kang, Mater. Des. 192 (2020) 108708.

[215] O. Tevet, D. Svetlizky, D. Harel, Z. Barkay, D. Geva, N. Eliaz, Materials 15 (2022) 638 (Basel).

[216] S. Tammas-Williams, H. Zhao, F. Léonard, F. Derguti, I. Todd, P.B. Prangnell, Mater. Charact. 102 (2015) 47-61.

[217] H. Galarraga, D.A. Lados, R.R. Dehoff, M.M. Kirka, P. Nandwana, Addit. Manuf. 10 (2016) 47-57.

[218] S. Sahoo, A. Joshi, V.K. Balla, M. Das, S. Roy, Mater. Sci. Eng. A 820 (2021) 141494.

[219] D. Tang, X. He, B. Wu, X. Wang, T. Wang, Y. Li, Theor. Appl. Fract. Mech. 119 (2022) 103322.

[220] G. Baudana, S. Biamino, B. Klöden, A. Kirchner, T. Weißgärber, B. Kieback, M. Pavese, D. Ugues, P. Fino, C. Badini, Intermetallics 73 (2016) 43-49.

[221] S. Huang, R.L. Narayan, J.H.K. Tan, S.L. Sing, W.Y. Yeong, Acta Mater. 204 (2021) 116522.

[222] M. Lopez, C. Pickett, E. Arrieta, L.E. Murr, R.B. Wicker, M. Ahlfors, D. Godfrey, F. Medina, Materials 13 (2020) 2604 (Basel).

[223] S.M. Gaytan, L.E. Murr, F. Medina, E. Martinez, M. Lopez, R.B. Wicker, Mater. Technol. 24 (2009) 180-190.

[224] G. Strano, L. Hao, R.M. Everson, K.E. Evans, J. Mater. Process. Technol. 213 (2013) 589-597.

[225] Q.B. Yan, T. Xue, S.F. Liu, W.L. Wang, Y. Wang, X. Song, X. Yang, W.W. Shang, J. Iron Steel Res. Int. 30 (2023) 165-175.\\
[226] J.T. Tuomi, R.V. Björkstrand, M.L. Pernu, M.V. Salmi, E.I. Huotilainen, J.E. Wolff, P.K. Vallittu, A.A. Mäkitie, J. Mater. Sci. Mater. Med. 28 (2017) $1-7$.

[227] P. Edwards, M. Ramulu, Mater. Sci. Eng. A 598 (2014) 327-337.

[228] K.S. Chan, M. Koike, R.L. Mason, T. Okabe, Metall. Mater. Trans. A 44 (2013) 1010-1022.

[229] M. Hinderdael, M. Strantza, D. De Baere, W. Devesse, I. De Graeve, H. Terryn, P. Guillaume, Materials 10 (2017) 993 (Basel).

[230] R. Molaei, A. Fatemi, N. Phan, Int. J. Fatigue 117 (2018) 352-370.

[231] M. Fousová, D. Vojtěch, K. Doubrava, M. Daniel, C.F. Lin, Materials 11 (2018) 537 (Basel).

[232] S. Rahmati, E. Vahabli, J. Adv. Manuf. Technol. 79 (2015) 823-829.

[233] A. International, ASTM E6-09be1 Standard Terminology Relating to Methods of Mechanical Testing, ASTM International: West Conshohocken, PA, USA, 2009.

[234] L. Qian, J. Mei, J. Liang, X. Wu, Mater. Sci. Technol. 21 (2005) 597-605.

[235] B. Baufeld, E. Brandl, O. van der Biest, J. Mater. Process. Technol. 211 (2011) 1146-1158.

[236] D. Sharma, S.S. Hiremath, N.B. Kenchappa, Fatigue Fract. Eng. Mater. Struct. 46 (2023) 975-986

[237] B.A. Szost, S. Terzi, F. Martina, D. Boisselier, A. Prytuliak, T. Pirling, M. Hofmann, D.J. Jarvis, Mater. Des. 89 (2016) 559-567.

[238] V. Chastand, P. Quaegebeur, W. Maia, E. Charkaluk, Mater. Charact. 143 (2018) $76-81$.

[239] N. Hrabe, T. Gnäupel-Herold, T. Quinn, Int. J. Fatigue 94 (2017) 202-210.

[240] B. Ahmad, S.O. van der Veen, M.E. Fitzpatrick, H. Guo, Addit. Manuf. 22 (2018) 571-582.

[241] M. Yakout, M.A. Elbestawi, S.C. Veldhuis, J. Manuf. Process. 52 (2020) 181-192.

[242] K. Dietrich, J. Diller, S. Dubiez-Le Goff, D. Bauer, P. Forêt, G. Witt, Addit. Manuf. 32 (2020) 100980

[243] M. Velasco-Castro, E. Hernández-Nava, I.A. Figueroa, I. Todd, R. Goodall, Heliyon 5 (2019) e02813.

[244] T.W. Na, W.R. Kim, S.M. Yang, O. Kwon, J.M. Park, G.H. Kim, K.H. Jung, C.W. Lee, H.K. Park, H.G. Kim, Mater. Charact. 143 (2018) 110-117.

[245] A. Carrozza, A. Aversa, F. Mazzucato, M. Lombardi, S. Biamino, A. Valente, P. Fino, Appl. Sci. 10 (2020) 4212.

[246] I.I. Kornilov, V.V. Vavilova, L.E. Fykin, R.P. Ozerov, S.P. Solowiev, V.P. Smirnov, Metall. Mater. Trans. 1 (1970) 2569-2571.

[247] A. Klassen, V.E. Forster, V. Juechter, C. Körner, J. Mater, Process. Technol. 247 (2017) 280-288.

[248] C. Panwisawas, C. Qiu, M.J. Anderson, Y. Sovani, R.P. Turner, M.M. Attallah, J.W. Brooks, H.C. Basoalto, Comput. Mater. Sci. 126 (2017) 479-490.

[249] T. Mukherjee, J.S. Zuback, A. De, T. DebRoy, Sci. Rep. 6 (2016) 19717.

[250] S. Keaveney, A. Shmeliov, V. Nicolosi, D.P. Dowling, Addit. Manuf. 36 (2020) 101514.

[251] C.A. Brice, B.T. Rosenberger, S.N. Sankaran, K.M. Taminger, B. Woods, R. Nasserrafi, Mater. Sci. Forum 618-619 (2009) 155-158.

[252] G. Zhang, J. Chen, M. Zheng, Z. Yan, X. Lu, X. Lin, W. Huang, Metals 10 (2020) 435 (Basel).

[253] T. Debroy, S.A. David, Rev. Mod. Phys. 67 (1995) 85-112.

[254] M. Fischer, D. Joguet, G. Robin, L. Peltier, P. Laheurte, Mater. Sci. Eng. C 62 (2016) 852-859.

[255] Q. Wang, C.J. Han, T. Choma, Q.S. Wei, C.Z. Yan, B. Song, Y.S. Shi, Mater. Des. 126 (2017) 268-277.

[256] M.A. Surmeneva, A. Koptyug, D. Khrapov, Y.F. Ivanov, T. Mishurova, S. Evsevleev, O. Prymak, K. Loza, M. Epple, G. Bruno, R.A. Surmenev, J. Mater. Process. Technol. 282 (2020) 116646.

[257] C. Schneider-Maunoury, L. Weiss, P. Acquier, D. Boisselier, P. Laheurte, Addit. Manuf. 17 (2017) 55-66.

[258] P. Li, D.H. Warner, A. Fatemi, N. Phan, Int. J. Fatigue 85 (2016) 130-143.

[259] S. Leuders, M. Thöne, A. Riemer, T. Niendorf, T. Tröster, H.A. Richard, H.J. Maier, Int. J. Fatigue 48 (2013) 300-307.

[260] G. Servetti, X. Zhang, Eng. Fract. Mech. 76 (2009) 1589-1602.

[261] V. Cain, L. Thijs, J. Van Humbeeck, B. Van Hooreweder, R. Knutsen, Addit. Manuf. 5 (2015) 68-76.

[262] H. Zhang, D. Dong, S. Su, A. Chen, Chin. J. Aeronaut. 32 (2019) 2383-2393.

[263] R. Labusch, Phys. Status Solidi B 41 (1970) 659-669.

[264] H. Conrad, Prog. Mater. Sci. 26 (1981) 123-403.

[265] T. Song, Z. Chen, X. Cui, S. Lu, H. Chen, H. Wang, T. Dong, B. Qin, K.C. Chan, M. Brandt, X. Liao, S.P. Ringer, M. Qian, Nature 618 (2023) 63-68.

[266] S. Zaefferer, Mater. Sci. Eng. A 344 (2003) 20-30.

[267] M.A. Meyers, O. Vöhringer, V.A. Lubarda, Acta Mater. 49 (2001) 4025-4039.

[268] F. Bridier, P. Villechaise, J. Mendez, Acta Mater. 53 (2005) 555-567.

[269] Z.W. Wyatt, W.J. Joost, D. Zhu, S. Ankem, Int. J. Plast. 39 (2012) 119-131.

[270] M. Yan, W. Xu, M.S. Dargusch, H.P. Tang, M. Brandt, M. Qian, Powder Metall. 57 (2014) 251-257.

[271] S.F. Jawed, C.D. Rabadia, Y.J. Liu, L.Q. Wang, Y.H. Li, X.H. Zhang, L.C. Zhang, J. Alloy. Compd. 792 (2019) 684-693.

[272] K. Yokota, A. Bahador, K. Shitara, J. Umeda, K. Kondoh, Mater. Sci. Eng. A 802 (2021) 140677.

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