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\title{Additive manufacturing of magnesium alloys: Characterization and post-processing }


\author{Shambhu Kumar Manjhi a, Prithivirajan Sekar ${ }^{b}$, Srikanth Bontha ${ }^{a}$, A.S.S. Balan ${ }^{a, *}$\\
a Department of Mechanical Engineering, National Institute of Technology Karnataka, Surathkal, Mangalore 575025, India\\
${ }^{\mathrm{b}}$ Department of Mechanical Engineering, Indian Institute Technology Madras, Chennai 600036, India}
\date{}


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Review



\section*{A R T I C L E I N F O}
\section*{Article history:}
Received 16 November 2022

Received in revised form

1 May 2023

Accepted 19 June 2023

Available online 22 June 2023

\section*{Keywords:}
Additive manufacturing

Direct energy deposition

Wire arc additive manufacturing

Magnesium alloy

Post-processing

Laser powder bed fusion

\begin{abstract}
A B S T R A C T Magnesium and its alloys remain perilous in the framework of light weighting and advanced devices structure such as rockets and satellites. However, the utilization of Magnesium $(\mathrm{Mg})$ is increasing every year, revealing growing demands in manufacturing industries. Manufacturing of $\mathrm{Mg}$ components is challenging because of their HCP crystal structure and limited ductility. In this context, additive manufacturing (AM) provides the flexibility to manufacture complex shape components with excellent dimensional stability. It also provides a new possibility for utilizing novel component structures that increase the applications for $\mathrm{Mg}$ alloy. This review herein pursues to holistically explore the additive manufacturing of $\mathrm{Mg}$ alloy with a synopsis of processes used and microstructure, mechanical properties, corrosion behaviour and postprocessing of AMed Mg alloy. The challenges and future scope of AMed Mg alloys are critically explored.\\
(C) 2023 The Authors. Publishing services by Elsevier B.V. on behalf of KeAi Communications Co. Ltd. This is an open access article under the CC BY-NC-ND license (\href{http://creativecommons.org/licenses/by-nc-nd/}{http://creativecommons.org/licenses/by-nc-nd/}
\end{abstract}

\section*{1. Introduction}
Magnesium is the lightest structure among all engineering materials. $\mathrm{Mg}$ has the lowest density of $1.74 \mathrm{~g} / \mathrm{cm}^{3}$ compared to the density of Aluminium (Al), Titanium (Ti) and steel, which are 2.71, 4.5 and $7.8 \mathrm{~g} / \mathrm{cm}^{3}$, respectively. The low density, high specific strength, and biodegradable nature make them an attractive material for manufacturing nosewheel doors, flap cover skin, oil tanks, floorings, wingtips, ducts, and seats, fuselage parts in aerospace industries [1], Front end structure, transfer case, engine cradle, incremental panel, steering wheel cores, cam covers, seat back, third-row seat frame in automotive industries [2] and cardiovascular stents, MAGNEZIX screw, micro clip for laryngeal microsurgery, biodegradable orthopaedic implant, wound-closing devices in biomedical sectors [3]. However, the application of Mg and Mg alloy is still limited due to their low ductility, corrosion resistance and flammability [4]. In addition, the manufacturing of $\mathrm{Mg}$ is challenging owing to its heat-sensitive nature. Currently, $>95 \%$ of $\mathrm{Mg}$ alloy products are manufactured using high die-pressure casting
\footnotetext{\begin{itemize}
  \item Corresponding author
\end{itemize}

E-mail address: \href{mailto:balan@nitk.edu.in}{balan@nitk.edu.in} (A.S.S. Balan).

Peer review under responsibility of Editorial Board of International Journal of Lightweight Materials and Manufacture.
}

[5]; however, it is unable to manufacture complex shapes such as porous structures and large components like cryogenic containers for aerospace applications [6], propellers for marine applications [7] with desired properties. In contrast, the application of wrought $\mathrm{Mg}$ alloy is less due to insufficient ductility, around $6 \%$ of elongation [8]. The processability using cold rolled, hot rolled, and the forging of $\mathrm{Mg}$ alloy is challenging owing to its HCP crystal structure, which is responsible for insufficient formability [9]. Apart from these challenges, the low corrosion resistance of $\mathrm{Mg}$ alloy makes their limited application. An additive manufacturing process can be used for manufactured $\mathrm{Mg}$ components to overcome these challenges. Further heat treatment and surface modification can be used as a post-process to achieve desired properties.

Additive manufacturing (AM) is a material fabricating technology with immense potential to deposit complex geometry parts, layer by layer, with less human intervention and high material efficiency [10]. In the last two decades, AM technique is becoming the most popular and preferable manufacturing process to be employed in the automobile [11], aerospace [12], architectural [13], military [13], medical [14] and electronics industries [15]. Indeed, $\mathrm{AM}$ has numerous benefits, such as handling materials and achieving near-net shape parts that reduce the need for tooling and re-fixturing. AM technique produces a whole component with lowers manufacturing costs and material waste. As a result, a\\
significantly lower buy-to-fly (BTF) ratio is attained in AM than in other conventional manufacturing techniques involving material removal [16]. However, limited research has been carried out in the field of $\mathrm{AM}$ of $\mathrm{Mg}$ alloy to date owing reactive nature of $\mathrm{Mg}$, which raises the possibility of safety concerns. In addition, other challenges, such as oxidation and evaporation, occurs during $\mathrm{Mg}$ deposition [17]. Several researchers [18-24] recently successfully deposited WE43Mg alloy using the laser powder bed fusion (LPBF) process. They also demonstrated risk control during the deposition of $\mathrm{Mg}$ powder. The Risk-controlling factors include storage of the $\mathrm{Mg}$ powder to maintain quality; personal training is required to control the process, cleaning the powder hopper properly and controlling the reactive gases during the deposition of $\mathrm{Mg}$. In addition, apart from safety concerns, another concern is the development of high-quality LPBFed Mg. Besides LPBF, various AM techniques have also been discovered for manufacturing Mg alloy, including wire arc additive manufacturing (WAAM), friction stir processing additive manufacturing (FSP-AM) and the sintering process. Among them, the WAAM process is highly preferable for the deposition of Mg alloy. Although FSP-AM is still under consideration in the community of $A M$, however, the technique is accepted because it follows AM strategy in a general sense. Recently structure and properties of additive manufacturing components were reviewed by Debroy et al. [18] and Zeng et al. [19]. Debroy et al. reviewed the process and physics, and Zeng et al. studied the microstructure of LPBFed of Mg alloy. Some review papers were published on additive manufacturing of Mg alloy [25-27]. However, they did not symmetrically explore the microstructure, mechanical properties, corrosion performance, and WAAMed and LPBFed Mg alloy post-processing. Derekar et al. [23] extensively reviewed the challenges in controlling the quality and part accuracy without considering the microstructure and mechanical properties of metals fabricated using WAAM. They studied the WAAM of metals, including Ti, Al, steel and nickel-based alloys. Wu et al. [24] have formulated a methodology to achieve high-quality WAAM parts. Indeed, they emphasized process selection, feedstock optimization and post-processing treatments necessary to manufacture structurally sound WAAM components. In addition, some of the reviews on WAAM were focused on specific materials, including aluminium [28], stainless steel [29] and titanium [25]. Several consolidated studies have been published regarding the additive manufacturing of magnesium alloys [7-15]. The composition-processing-microstructural properties relationship in AM is not symmetrically established yet. A significant cause for this is that results of microstructural properties relationships for AMed $\mathrm{Mg}$ alloys have revealed some discrepancies in various reports. Based on the review papers available in the open literature, it is found that few of the reviews were dedicated to AM of magnesium due to the minimal number of studies carried out. The additive manufacturing of magnesium is extensively reviewed to bridge this research gap in this work, with specific attention paid to WAAM and LPBF processes. Moreover, microstructure, mechanical properties, corrosion behaviour and post-processing of $\mathrm{Mg}$ alloys deposited using WAAM and LPBF are also summarized. In addition, challenges and future scopes in the WAAMed and LPBFed Mg are also discussed. The overview of this review paper is demonstrated in Fig. 1.

\section*{2. Publication trend on additive manufacturing of $\mathbf{M g}$ alloys}
Based on a search from three primary journal databases, including Scopus, Web of Science (WOS) and google scholar, 125 articles are dedicated to laser powder bed fusion (LPBF), and 28 publications are precisely matched to magnesium and magnesium alloy deposited using LPBF process. Similarly, 74 articles are dedicated to WAAM, out of which 14 papers are entirely compared to WAAMed Magnesium alloy. The comparison of publication trends year-wise data related to LPBF and WAAMed Magnesium alloy are vividly shown in Fig. 2.

Fig. 2 (a) shows the number of publications on LPBF and WAAMed Mg alloy has rapidly increased since 2020. This increasing trend of publications signifies the importance of research in AMed magnesium alloy. In particular, the comparison of worldwide publication percentages on AMed Mg alloys related to LPBF and WAAM of magnesium alloy is represented in Fig. 2 (b). This figure shows that Asia and Europe are contributing to AM of Mg alloys to a greater extent.

\section*{3. Wire arc additive manufacturing (WAAM) process}
Wire arc additive manufacturing (WAAM) is a direct energy deposition (DED) AM process consisting of a wire feeder system and an electric arc heating source, depicted in Fig. 3. During the WAAM process, filler metal wire is heated through an electric arc and deposited as a bead on the substrate with the help of industrial robots and gantries. The WAAM system is implemented into conventional welding set-up, reducing the machine's cost. The comparison between commonly used traditional welding-based WAAM processes is listed in Table 1.

The first WAAM process was patented dates back to 1920 [27]. This manufacturing method proves suitable over other AM processes due to less material waste, low machine set-up cost, low production cost, higher production rate, $100 \%$ wire material utilization and superior fusion of layers within parts of the components [31].

For instance, Guo et al. reported that the deposition rate of WAAM is around $10 \mathrm{~kg} / \mathrm{h}$ with high deposition efficiency, which ensures the capability to deposit large components [32]. Similarly, Zhang et al. [33] claimed that deposition time and post-milling time could be decreased by $40-60 \%$ and $15-20 \%$ compared with conventional manufacturing. However, the production of large parts with fewer intricacies with minimized duration and components obtained using the WAAM process is in disparity with those fabricated in conventional ways. However, the surface roughness of components is subordinate to ones manufactured using traditional techniques [34]. However, challenges such as poor surface finish, tensile residual stress, deformation due to high heat input and less dimensional accuracy results in stair-stepping are significant issues in stabilizing the WAAM process [35]. In contrast to conventional WAAM, the cold metal transfer (CMT)-WAAM process is highly suitable for reducing the above defects due to controlled heat input. The CMT machine detects a short circuit which sends a signal that retracts the filler material and gives the weld time to cool before each drop is placed. Therefore, it is named cold metal transfer welding. Because of the cold metal transfer mechanism, welds are smoother and more robust than hotter welds [36]. In addition, CMT allows the material transfer to occur with a relatively lesser current flow. Therefore, this technique is preferable during the deposition of heat-sensitive materials such as Magnesium ( $\mathrm{Mg}$ ) and Aluminium (Al).

\subsection*{3.1. WAAM set-up and implementation}
The complete process flow of WAAM is depicted in Fig. 4. The software system mainly consists of three independent methods: prototype modelling, layer slicing and tool path planning. For instance, various contour fitting methods for single bead [37,38] and multi-bead overlap models $[39,40]$ were studied. The tool path strategy is also an important parameter influencing the deposition quality and efficiency. For example, raster [28] and Zig-Zag [29] tool

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-03}
\end{center}

Fig. 1. An overview of this review paper.

path strategies are mostly adapted when depositing simple geometrical structures. In Zig-Zag tool path planning, relatively fewer start and stop points lead to higher deposition efficiency, but some defects are observed at arc off or on conditions. However, in the case of complex geometry deposition like curves and circles, contour [41] and spiral [42] tool path strategies are more significant. In addition, both tool paths exhibit a high deposition rate, but the contour tool path created a closed curve due to high start and stop points that lead to the uneven surface and protrusion of the deposited surface. After tool path planning, materials deposition is carried out. For material deposition, selecting a WAAM set-up is essential to achieve sound deposition quality.

Generally, three types of WAAM processes listed in Table 1 are commonly used. The symmetric view of the WAAM process is vividly shown in Fig. 5. These three types of WAAM process are unfavourable to deposited batter-quality magnesium alloy due to high heat input, which leads to higher oxidation, spattering, stress, distortion and cracks. Therefore, reducing the heat input with sufficient current to fabricate Mg alloy is crucial. Many authors used the cold metal transfer-based WAAM process to address these challenges while depositing the Mg alloy [32,33] because the heat input is 33\% lower than conventional WAAM processes with less spattering. Schierl et al. [44] reported that the droplet-cutting mode occurred without the aid of electromagnetic force.\\
Therefore, the spatter is significantly less in the CMT-WAAM process. The low heat input of the process is due to a stable short circuit with a low current. This low heat input prevents the oxidation and development of residual stress. The comparison of deposition quality of materials using conventional WAAM and CMT-WAAM is shown in Fig. 6 (a, b, c \& d). These deposited beads are evidence of high spattering in GMAW, humping defect or irregular bead in GTAW and high oxidation in PAW-based WAAM deposition. However, in CMT-WAAM deposition, no such type of defects are observed. Therefore, the above results ensure that the CMT-WAAM process is highly suitable and recommended for the deposition of $\mathrm{Mg}$ alloy.

The CMT process is a modified version of the MIG welding process, based on short-circuiting transfer developed by Fronius Pvt Ltd Austria in 2004, and the schematic views of the CMT-WAAM mechanism can be seen shown in Fig. 7 (a), which consists of three electrical signal cycles: peak current, background current and short-circuit phase. During the peak current phase, a constant voltage corresponding to a high current pulse owing to the arc ignition leads to heating the electrode wire to form a molten droplet. In the background current phase, the current and voltage drastically dropped for a few milliseconds, preventing the globular transfer of molten droplets on the tip of the wire. This phase remains till a short circuit occurs. During the short circuit phase, the\\
\includegraphics[max width=\textwidth, center]{2024_04_13_7e5fbe1213c12de51fd6g-04}

Fig. 2. Publication trends of AMed Mg Alloy (a) comparison showing publication trends of LPBF and WAAM of Mg alloys year-wise (b) comparison of publication \% worldwide on PBF and WAAM of Mg alloy.

arc voltage drops to almost zero, and immediately, a signal is provided to return the wire that leads to molten droplet cutting and transfer to the substrate. These phases of the electrical cycles and droplet transfer mechanism are shown in Fig. 7 (b \& c).

\subsection*{3.2. Oxidation prevention}
Magnesium is highly reactive with oxygen at elevated temperatures. Therefore, high oxidation occurs on the fabricated component during deposition due to continuous heat input [50]. Consequently, oxidation is one of the significant causes that hinder the adaption of the WAAM process for the deposition of magnesium components. One of the simplest methods widely accepted to prevent oxidation is to perform WAAM operation in a closed chamber with inert gas purging, as shown in Fig. 8 (a \& b). During WAAM, this chamber is filled with an inert gas to prevent oxidation. However, the installation of WAAM set up in a closed chamber limits the axial movement of a robotic arm, thereby reducing the size of the component to be fabricated. In addition, the shielding gas inside the closed chamber must be replaced frequently, influencing the deposition efficiency. To summarise, the shielded chamber WAAM set-up limits the fabrication capabilities of large and complex shaped components [51]. To overcome this limitation, shielding gas devices in WAAM set-up are gradually developing. Nowadays, a tracing shielding gas device is attached to the welding torch, which helps to form a volume region filled with shielding gas. This type of WAAM set-up can deposit relatively better significant parts with the prevention of oxidation.

The design of tracing inert gas devices is generally compact and enables uniform distribution of inert gas flowing over the deposition region. The tracing shielding gas attachment in the WAAM setup is usually custom-made to fit in a specific welding torch for a particular deposition. For instance, Fig. 9 depicts the tracing shielding gas set-up in which a ceramic nozzle is mounted on a welding torch. However, the extent of inert gas convergence provided at the deposition zone is not always adequate.

Moreover, The Welding Institute (TWI) developed a tracing shielding gas on GTAW when applying it to trace shielding gas on a root pass groove [56]. The groove geometry, such as width and depth, determines whether shielding gas flood at weld pass is appropriate and whether sufficient shielding gas protection away from the welding torch is provided. The GMAW and automated GTAW requires relatively more significant tracing shielding during high-speed deposition. Resistance glass is employed instead of metal for shielding for better visibility. Hence, any successful tracing shielding gas design that requires experience and proven commercial production should also be available for circumferential fillet and strain in welding. The three typical tracing shielding gas devices with a complete configuration that can be used in the deposition of Mg alloy by WAAM are depicted in Fig. 9 (a, b \& c). These designs are generally preferred for large shielding volumes. Conversely, the design shown in Fig. 9 (c) employs a shielding device with a hollow cylindrical configuration. Moreover, this design uses a relatively smaller gas shielding area, increasing the operating freedom during deposition.

\subsection*{3.3. Magnesium alloy system deposited by WAAM}
Limited research has been carried out in the wire arc additive manufacturing of $\mathrm{Mg}$ alloy since this new technique is at the early stage of commercialization; however, in present days, the attention of researchers has moved towards the deposition of $\mathrm{Mg}$ alloy using modern manufacturing processes, such as additive manufacturing and hybrid Additive manufacturing. This is because the demand for $\mathrm{Mg}$ and $\mathrm{Mg}$ alloys is significantly increasing in the aerospace, automotive and biomedical industries.

The literature survey revealed that most research was performed on the $\mathrm{AZ}$ series of $\mathrm{Mg}$ alloys. Apart from the $\mathrm{AZ}$ series, only AEX11 Mg was deposited by the WAAM process to date [57]. Therefore, immense scope exists in the deposition of different grades of $\mathrm{Mg}$ alloy and in realizing the benefits and challenges of extending the application of AM Mg.

\subsection*{3.3.1. AZ Series of $M g$ alloy deposited using WAAM}
Recently, Cao et al. [58] deposited AZ31 Mg deposited by the ultrasonic pulse frequency (UPF)-based WAAM process. They achieved equiaxed grain in the bottom, middle and top sections. During UPF-WAAM, the UPF arc induced significant vibration in the molten pool, influencing heat dissipation and fluid flow. This phenomenon enhanced the \% of isotropic mechanical properties. Interestingly, Fang et al. [59] studied the microstructure and mechanical properties of GTA-WAAMed AZ31 Mg alloy. They reported deposited thin walls composed of completely fine equiaxed grain with approximately $0.1 \%$ porosity without heat treatment. However, due to rapid cooling, primarily microstructure revealed $\alpha-\mathrm{Mg}$ phase with negligible Al-Mn phase precipitation. The UTS and \% EL

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-05}
\end{center}

Fig. 3. Classification of the additive manufacturing process based on the heat source and material feeding strategy [30].

Table 1

Different types of WAAM processes [24].

\begin{center}
\begin{tabular}{|c|c|c|}
\hline
Welding Machine setup & Source of Energy & Feature \\
\hline
GMAW & Gas Metal-Arc & Consumable electrode, Deposition rate $3-4 \mathrm{~kg} / \mathrm{h}$. Poor arc stability, spattering \\
\hline
CMT & Tandem Gas Metal-Arc & \begin{tabular}{l}
Reciprocating consumable wire electrode, Deposition rate $2-3 \mathrm{~kg} / \mathrm{h}$. Low heat input, zero spattering, high process \\
tolerance \\
\end{tabular} \\
\hline
GTAW & Gas Tungsten-Arc & Non-consumable electrode, Separate wire feeder, rotation of torch and electrode, Deposition rate 1-2 kg/h. \\
\hline
PAW & Plasma & The mechanism is the same as GTAW, only differs based on the energy source \\
\hline
Tandem/twin wire & MIG/MAG/CMT & Consumable wire has a very high deposition rate $(15-20 \mathrm{~kg} / \mathrm{h})$ \\
\hline
\end{tabular}
\end{center}

in the build direction were slightly lower compared to the travel direction. This is because of several pores found in fractured build direction surfaces, which are responsible for poor mechanical properties. Zhang et al. [60] investigated the effect of solution annealing on microstructure and corrosion performance of WAAMed AZ91Mg thin wall. They observed WAAMed sample consists of an equiaxed $\mathrm{Mg}$ matrix, $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ and $\mathrm{Al}_{8} \mathrm{Mn}_{5}$ secondary phase with an average grain size of $7.9 \mu \mathrm{m}$. During the annealing process, the secondary phases of $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ completely dissolved; only the $\mathrm{Al}_{8} \mathrm{Mn}_{5}$ phase with the increased grain size of $62.4 \mu \mathrm{m}$ remained, as depicted in Fig. 10.

This is because $\mathrm{Al}-\mathrm{Mn}$ particle sizes were larger than the size of $\mathrm{Mg}-\mathrm{Al}$. Therefore, the comparatively small particles dissolved completely. Moreover, the corrosion resistance of as-deposited samples was higher than heat-treated samples because of the high amount of secondary phase particles presented in asdeposited samples. Gao et al. [61] investigated the microstructure evolution and mechanical properties of the WAAMed AZ80M Mg thin wall. They reported the microstructure of the bottom, middle and top sections were slightly different. The top zone consists of many equiaxed and dendritic grains, similar to arc-welded microstructure. In contrast, the middle and bottom sections comprised equiaxed grains with dendrite segregation and columnar dendritic grains because of large temperature gradients between the melt pool and substrate, leading to epitaxial growth of columnar grain along the fusion line.

\subsection*{3.4. Microstructure and mechanical properties}
The mechanical properties of WAAMed components are higher than the cast and comparable to wrought Mg alloys. Therefore, the contribution of WAAM is rapidly increasing, especially in largescale manufacturing industries. Generally, the mechanical properties of WAAMed thin walls are influenced by welding parameters such as voltage, current and travel speed. In addition, the mechanical properties of WAAMed components are also affected by microstructure, crystallographic orientation and mesostructured. For instance, Yang et al. [62] investigated the microstructural and mechanical properties of the bottom, middle and top sections of WAAMed AZ31 Mg alloy. They reported the bottom, middle and top sections of deposited thin walls composed of vertical columnar, directionally changed, and equiaxed dendrites microstructure owing to different temperature history at each layer, as shown in Fig. 11(a-d). Moreover, the middle and bottom sections were composed of columnar dendrites without secondary dendrite arms, as seen in Fig. 11(c-f). The reason behind this phenomenon is temperature gradient of the molten pool gradually decreased due to heat accumulation. Thus, the transition of secondary dendrite

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-06}
\end{center}

Fig. 4. Flow of wire and arc additive manufacturing process [43].

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-06(1)}
\end{center}

Fig. 5. Symmetric representation of the WAAM process [45].

arms to primary dendrite occurred. The average primary dendrite arm spacing of the bottom, middle, and top regions are $17 \mu \mathrm{m}$, $29 \mu \mathrm{m}$, and $39 \mu \mathrm{m}$, respectively [62]. The microstructure and mechanical properties of various grades of $\mathrm{Mg}$ alloy fabricated using WAAM are listed in Table 2. This table shows that some authors reported that the mechanical properties in the travel direction are higher than that of the build direction. This is because, during the deposition of the next layer, the upper portion of the previous layer is remelted and mixed with melted-fed wire. This led to the formation new melt pool followed by rapid solidification.

The space between the newly deposited layer and the previous layer is not only an interface but also a solid-liquid boundary. Therefore, the distribution of grains is complex in this zone due to temperature differences, which leads to a mixture of coarse and fine grains with micropores and microcracks. Because of micropores and microcracks, failure occurs at the interface, resulting in

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-07}
\end{center}

Fig. 6. Comparison of deposition quality of different WAAM processes (a) GMAW [46], (b) GTAW [47], (c) PAW-WAAM [48], (d) CMT-WAAM.

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-07(1)}
\end{center}

Fig. 7. (a) Schematic diagram of CMT-WAAM (b) droplet formation and cutting image [49] (c) electric phase cycle [49].\\
\includegraphics[max width=\textwidth, center]{2024_04_13_7e5fbe1213c12de51fd6g-08}

Fig. 8. GTAW-based WAAM set up in a closed chamber (a) physical Map (b) schematic view [52].

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-08(1)}
\end{center}

Fig. 9. Tracing shielding gas deceive (a) design 1 [53] (b) design 2 [54] (c) design 3 [55].

anisotropic mechanical properties of the WAAMed component. Therefore, a deep study is essential to avoid the anisotropic mechanical properties of deposition.

The influence of pulse frequency on microstructure and mechanical properties of AZ31 Mg alloy fabricated using the GTAWbased WAAM process was studied by Guo et al. [32] and reported that no pores were observed in the microstructure, as shown in Fig. 12(a-f). They also noticed a substantial microstructure difference in grain size with increased pulse frequency obtained. The grain size first decreased with increasing pulse frequency and then increased marginally. The microstructure of the AZ31 sample deposited at pulse frequencies $10 \mathrm{~Hz}$ and $5 \mathrm{~Hz}$ is relatively fine and uniform with a grain size of $21 \mu \mathrm{m}$, vividly shown in Fig. 12 (c \& d). The coarse grain size of $39 \mu \mathrm{m}$ is obtained at pulse frequencies $500 \mathrm{~Hz}, 100 \mathrm{~Hz}$, and $1 \mathrm{~Hz}$. as shown in Fig. 12 (a, b \& f), respectively [71]. Generally, in the CMT arc welding process, the wire starts to melt at the peak current phase, and the molten droplet is still attached to the wire tip owing to surface tension till the base current occurs [72]. After that, when the base current turn to the peak current phase, the adhered droplet on the tip of the wire drops on the last deposited layer. Consequently, droplets have more time to spread over the layers at the low pulse frequency and become larger melt pools, resulting in fewer coarse grains than those at the high frequency. From Fig. 12, it is noted that pulse frequency appreciably influenced the microstructure, especially grain size. Mainly, two factors of pulse frequency contribute more to the change in grain size. The first factor is stirring the melt pool by pulse current. The pulse current stirs the melt pool, producing a high cooling rate that refines the grains. In addition, when increased the pulse current, the plasma momentum with electromagnetic force rises automatically. The plasma momentum generates pressure and shearing force on the surface of the melt pool. Hence, the electromagnetic force pushed the fluid inside the melt pool rapidly inwards and then downwards or down to the axis [73], resulting in rises in weld pool oscillations leading to dendrite fragmentation and producing more heterogeneous nucleation, consequently refining the grains [74]. Secondly, the pulse current directly increased the cooling rate, resulting in grain refinement. The peak current is mainly used as heat input to melt the metal wire, while the base current is much smaller than the peak current and is used to maintain the arc. During the base current phase, the heat input suddenly decreased; therefore, the cooling rate of the weld pool increased, leading to heterogeneous nucleation resulting in grain refinement.

Guo et al. [65] deposited AZ80M Mg alloy deposited using GTAW-based WAAM. This study conducted three different heat treatment processes, $\mathrm{T} 4, \mathrm{~T} 5$ and $\mathrm{T} 6$, on the deposited specimen. The as-deposited specimen was composed of an equiaxed dendrite, and secondary phase particles were distributed along the interdendritic zone to form a network structure. T4 heat treatment dissolved the eutectic phase and eliminated the inhomogeneity of the interlaminar structure. In contrast, T5 inherited the characteristics of dendritic structure and discontinuous secondary phase distribution along the eutectic phase. However, T6 heat treatment formed a continuous phase in the microstructure, effectively enhancing grain uniformity. In addition, the as-build AZ80 Mg alloy specimen

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-09}
\end{center}

Fig. 10. SEM + EDS mapping of WAAMed AZ91 samples (a) As deposited (b) heat treated [60].

exhibited mechanical anisotropy properties and travel direction, displaying significantly better mechanical properties than those observed in the build direction. From the above results, the mechanical properties of AZ80 Mg alloy are remarkably more excellent than AZ31 Mg. This could be attributed to the additional Wt \% of Al in AZ80 Mg alloy [75]. Wang et al. [63] varied the parameters of CMT-WAAM during single tracks deposition of AZ31 Mg alloy. They obtained three different characteristics of tracks with the obvious change in their corresponding weld geometries. Among them, the CMT characteristic that exhibited relatively better wettability corresponding to a contact angle of $114.1^{\circ}$ was chosen for the deposition of the thin wall. Optical micrographs revealed fine columnar dendrites microstructure growing along the build direction in the bottom, middle and top layers. Moreover, equiaxed grains were observed in the heat-affected zone between each layer in travel or welding direction with some pores, as shown in Fig. 13(a-c). EBSD analysis of cross section revealed the difference in grain morphology, texture intensity, and Schmid factor in both build and travel direction, resulting in anisotropic mechanical properties as depicted in Fig. 13 (a1 \& c1). Interestingly, TD and BD tensile test specimens failed at the HAZ zone between interlayers [76]. Gussev et al. [77] reported that the hard grains with a Schmid factor $\leq 0.35$ were hard to slip, while soft grains with a Schmid factor $\geq 0.4$ were easy to slip. According to the statistics of the Schmid factor, the fraction of soft grains was much higher than that of the fraction of hard grains, which indicated that CMT-WAAMed AZ31 Mg components have better ductility than cast and wrought $\mathrm{Mg}$ alloy. Moreover, the average fraction of hard grains at layer boundaries and at the interlayer of HAZ are $30.8 \%$ and $42 \%$, respectively. From this figure, it is inferred that the density of hard grains at the interlayer of HAZ is superior to deposited layers, which indicates that the interlayer of HAZ exhibited comparatively less plastic strain accumulation, resulting in lower plastic properties than deposited layers. In addition, Somekawa et al. [78] reported that the low angle grains boundaries (LAGBs) $\leq 15^{\circ}$ exhibited little effect on dislocation motion and can reduce the intergranular crack initiation, while high angle grain boundaries (HAGBs) $>15^{\circ}$ lead to dislocation pile-up, resulting in high plastic deformation resistance. Therefore, according to the statistics chart of misorientation angles, given in Fig. $13\left(a_{1}-d_{1}\right)$, all regions show more LAGBs $\left(\leq 15^{\circ}\right)$ microstructure, which indicates that plastic deformation in CMTWAAMed AZ31 Mg alloy occurred easily. Form the above observation, it is noted that wire arc additively manufactured $\mathrm{Mg}$ alloys exhibited equiaxed grain and excellent mechanical properties with higher elongation \%.

Gneiger et al. [57] fabricated rare-earth-containing Mg alloys using custom-made AEX11 Mg wires. The mechanical properties were significantly lower than the WAAMed AZ61A in both travel and build directions. Interestingly, after T6 heat treatment, the mechanical properties of AEX11 are relatively greater than WAAMed- AZ61A because of dissolution and change in the shape of the $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ secondary phase particles due to precipitate Han et al. [79] fabricated AZ91 Mg alloy using a CMT-based WAAM technique that exhibited $2 \%$ anisotropy in UTS and $\%$ elongation. Guo et al. [65] monitored thermal cycles during the deposition of AZ80 Mg alloy using a thermocouple placed at the bottom of the substrate. They observed that the peak temperature reached a maximum of $450^{\circ} \mathrm{C}$ during the first $500 \mathrm{~s}$. Consequently, the temperatures gradually dropped below $300{ }^{\circ} \mathrm{C}$ by the end of $3500 \mathrm{~s}$ due to an increased distance between the deposited layers and thermocouples.\\
\includegraphics[max width=\textwidth, center]{2024_04_13_7e5fbe1213c12de51fd6g-10}

Fig. 11. SEM micrographs of the longitudinal section (X-Z plane) of the AZ31 thin wall (a) low and (b) high magnification micrographs of the top region, (c) low and (d) high magnification micrographs of the middle region, (e) low and (f) high magnification micrographs of the bottom region [62].

WAAMed AZ80 Mg evinced columnar to equiaxed grains transition (CET), resulting in mechanical anisotropy. However, the mechanical properties of WAAM AZ80 Mg were comparable to wrought AZ80 Mg [80]. Takagi et al. [67] build three AZ31B specimens using MIGWAAM by maintaining constant current and voltage with varying wire feed rates of 400,600 , and $800 \mathrm{~mm} / \mathrm{min}$. Among these parameters, samples with minor defects were obtained at $100 \mathrm{~A}, 10 \mathrm{~V}$ and $800 \mathrm{~mm} / \mathrm{min}$. Indeed, the porosity ratio \% of 0.00025 during WAAM of AZ31 B was negligible compared to $50 \%$ and $48 \%$ exhibited by die-casting [81] and SLM techniques [82].

The Grain size plays a crucial role in determining the mechanical properties of materials, such as strength, hardness, ductility, fatigue and creep resistance. It is well established that according to Hall and Petch's equation, a relatively larger grain size adversely affects the tensile strength of $\mathrm{Mg}$ [61]. Various techniques such as heat treatment, serve plastic deformation, and addition of foreign particles used as a nano-reinforcement technique in magnesium matrix is employed to achieve grain refinement. LI et al. [69] achieved fine grain of WAAMed AZ31 Mg compared to cast Mg. This is because, during the WAAM process, the metallic wire melted locally by electric arc generation due to the occurrence of a short circuit. As a result, the heat input is controlled, and due to the rapid cooling of beads, highly homogeneous nucleation results in grain refinements.

EBSD analysis is shown in Fig. 14 conforms such a type of assertion. In addition, EBSD micrographs revealed that the grains of\\
WAAMed AZ31 Mg were significantly refined. The grain size achieved during WAAMed $\mathrm{Mg}$ is comparable to the grain refinement obtained by severe plastic deformation. Indeed, the mechanical properties of WAAMed Mg samples were relatively higher when compared to their cast $\mathrm{Mg}$ counterparts [83].

Moreover, non-textured equiaxial crystallographic orientation is also revealed with grain refinement in WAAMed magnesium alloy. This is related to the transient solidification of the molten pool because of allowing the end of the wire to contact the substrate plate. The grain refinement of various grades of $\mathrm{Mg}$ alloy is listed in Table 2.

\subsection*{3.5. Corrosion performance}
Mg exhibits relatively lower corrosion resistance, which remains a hindrance in extending the application of $\mathrm{Mg}$ alloys. Indeed, understanding of corrosion mechanism of AMed Mg alloy is of prime importance to widen application. The corrosion performance of $\mathrm{Mg}$ alloy depends on grains size, formation of secondary phase particles and texture of materials. Therefore, Li et al. [69] manufactured AZ31 Mg alloy using CMT-WAAM and studied the effect of grains size and orientation on corrosion performance in $0.5 \mathrm{wt} \% \mathrm{NaCl}$ solution. They reported that WAAMed AZ31 Mg specimens' corrosion resistance was slightly higher or comparable to cast specimens. The formation of $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ phases is higher in the cast; in contrast, WAAMed specimens are composed of negligible $\mathrm{Mg}_{17} \mathrm{Al}_{12}$

Table 2

Microstructure and UTS, YS and \%EL of WAAMed Mg system.

\begin{center}
\begin{tabular}{|c|c|c|c|c|c|c|c|}
\hline
Mg Alloy & \begin{tabular}{l}
Sample \\
Direction \\
\end{tabular} & UTS (Mpa) & $\mathrm{YS}(\mathrm{MPa})$ & $\%$ EL & Grain size & Grain type & Ref. \\
\hline
AZ31 & \begin{tabular}{l}
$\mathrm{TD}$ \\
$\mathrm{BD}$ \\
$\mathrm{BD}$ \\
\end{tabular} & \begin{tabular}{l}
$210.5 \pm 18.2$ \\
$225.7 \pm 12.1$ \\
$210.5 \pm 3.5$ \\
\end{tabular} & \begin{tabular}{l}
$125.9 \pm 5.0$ \\
$85.4 \pm 3.0$ \\
$131.6 \pm 4.2$ \\
\end{tabular} & \begin{tabular}{l}
$17.2 \pm 4.2$ \\
$28.3 \pm 2.0$ \\
$10.55 \pm 1.61$ \\
\end{tabular} &  & \begin{tabular}{l}
The bottom, middle and top sections are \\
composed of columnar dendrites, \\
direction-changed columnar dendrites \\
and equiaxed dendrites \\
\end{tabular} & [63] \\
\hline
AZ61A & \begin{tabular}{l}
$\mathrm{TD}$ \\
$\mathrm{BD}$ \\
$\mathrm{BD}$ \\
\end{tabular} & \begin{tabular}{l}
$264.1 \pm 1.8$ \\
$256.4 \pm 10.1$ \\
$237 \pm 6.3$ \\
\end{tabular} & \begin{tabular}{l}
$104.4 \pm 1.6$ \\
$99.2 \pm 1.7$ \\
$119 \pm 13.4$ \\
\end{tabular} & \begin{tabular}{l}
$15.4 \pm 0.7$ \\
$15.3 \pm 3.5$ \\
$12 \pm 0.7$ \\
\end{tabular} &  & \begin{tabular}{l}
Globular and equiaxed without the \\
formation of elongated grains \\
\end{tabular} & [64] \\
\hline
AZ80M & \begin{tabular}{l}
$\mathrm{TD}$ \\
$\mathrm{BD}$ \\
\end{tabular} & \begin{tabular}{l}
280 \\
230 \\
\end{tabular} &  & \begin{tabular}{l}
$15 \%$ \\
$13 \%$ \\
\end{tabular} & - & equiaxed dendrite & $[65]$ \\
\hline
AZ91 & \begin{tabular}{l}
TD \\
$\mathrm{BD}$ \\
\end{tabular} & \begin{tabular}{l}
$245.2 \pm 1.0$ \\
$250.3 \pm 2.6$ \\
\end{tabular} &  & \begin{tabular}{l}
$16.3 \pm 1.0$ \\
$17.5 \pm 1.6$ \\
\end{tabular} & \begin{tabular}{l}
Grain size above and below the fusion \\
line are $4.8-25.2$ and $20.1-48.5 \mu \mathrm{m}$ \\
\end{tabular} & Equiaxed grain & [66] \\
\hline
AZ80M & \begin{tabular}{l}
$\mathrm{TD}$ \\
$\mathrm{BD}$ \\
\end{tabular} & \begin{tabular}{l}
$308 \pm 6.5$ \\
$237 \pm 6.3$ \\
\end{tabular} & \begin{tabular}{l}
$146 \pm 46.7$ \\
$119 \pm 13.4$ \\
\end{tabular} & \begin{tabular}{l}
$15 \pm 0.5$ \\
$12 \pm 0.7$ \\
\end{tabular} & - & \begin{tabular}{l}
The bottom, middle and Top sections \\
are composed of columnar dendritic, \\
equiaxed with dendrite segregation, \\
and dendritic \\
\end{tabular} & [61] \\
\hline
AZ31 & \begin{tabular}{l}
$\mathrm{TD}$ \\
$\mathrm{BD}$ \\
\end{tabular} &  &  &  &  &  &  \\
\hline
\begin{tabular}{l}
 \\
AEX11 \\
\end{tabular} & \begin{tabular}{l}
$\mathrm{TD}$ \\
$\mathrm{BD}$ \\
\end{tabular} & \begin{tabular}{l}
$275 \& 243$ \\
$272 \& 233$ \\
\end{tabular} &  & \begin{tabular}{l}
$16.8 \& 4.5$ \\
$17.2 \& 5.2$ \\
\end{tabular} &  &  & [57] \\
\hline
AZ31B & BD & 239 &  & 21 & \begin{tabular}{l}
Boundary, Middle and Top sections \\
composed of 18,55 and $80 \mu \mathrm{m}$ grain size \\
\end{tabular} & \begin{tabular}{l}
Fine and coarse dendrite microstructure \\
varies with the wire feed rate \\
\end{tabular} & [67] \\
\hline
AZ31 & \begin{tabular}{l}
$\mathrm{TD}$ \\
$\mathrm{BD}$ \\
\end{tabular} & \begin{tabular}{l}
$222.9 \pm 5.4$ \\
$190.7 \pm 23.3$ \\
\end{tabular} & \begin{tabular}{l}
$109.1 \pm 17.9$ \\
$94.7 \pm 2.0$ \\
\end{tabular} & \begin{tabular}{l}
$20.26 \pm 3.79$ \\
$13.82 \pm 3.98$ \\
\end{tabular} & Avg grain diameter is $24.7 \mu \mathrm{m}$ & \begin{tabular}{l}
Fine equiaxed grains at different \\
regions, bottom, middle and top \\
\end{tabular} & [59] \\
\hline
AZ31 & \begin{tabular}{l}
$\mathrm{TD}$ \\
$\mathrm{BD}$ \\
\end{tabular} & \begin{tabular}{l}
211.2 \\
203.3 \\
\end{tabular} &  & \begin{tabular}{l}
22.27 \\
20.19 \\
\end{tabular} & - & Complete equiaxed microstructure & [58] \\
\hline
AZ91 & - & - & - & - & \begin{tabular}{l}
As WAAMed grain size ranges from \\
$2 \mathrm{~mm}$ to $25 \mathrm{~mm}$. while heat treated \\
$30 \mathrm{~mm}-1050 \mathrm{~mm}$ \\
\end{tabular} & \begin{tabular}{l}
Equiaxed grains are fine in WAAM \\
during the course in HT \\
\end{tabular} & [60] \\
\hline
\begin{tabular}{l}
AZ91D \\
AZ31 \\
\end{tabular} & - & - & - & - & \begin{tabular}{l}
The average grain diameter of $3.35 \mu \mathrm{m}$ \\
Avg grains in TD is $12 \mu \mathrm{m}$ while in BD, \\
$6.9 \mu \mathrm{m}$ \\
\end{tabular} & \begin{tabular}{l}
fine equiaxed grains \\
Equiaxed grain texture \\
\end{tabular} & \begin{tabular}{l}
$[68]$ \\
$[69]$ \\
\end{tabular} \\
\hline
AZ80M & \begin{tabular}{l}
TD \\
$\mathrm{BD}$ \\
\end{tabular} & \begin{tabular}{l}
308.7 \\
237.3 \\
\end{tabular} & \begin{tabular}{l}
146 \\
119 \\
\end{tabular} & \begin{tabular}{l}
15.4 \\
12.2 \\
\end{tabular} &  & Equiaxed and dendritic grains & $[70]$ \\
\hline
\begin{tabular}{l}
AZ31 \\
AZ31 \\
\end{tabular} & \begin{tabular}{l}
- \\
$\mathrm{TD}$ \\
$\mathrm{BD}$ \\
\end{tabular} & \begin{tabular}{l}
260 \\
$151.9 \pm 12.9$ \\
$210.5 \pm 3.5$ \\
\end{tabular} & \begin{tabular}{l}
102 \\
$71.2 \pm 4.5$ \\
$131.6 \pm 4.2$ \\
\end{tabular} & \begin{tabular}{l}
23 \\
$7.54 \pm 1.32$ \\
$10.55 \pm 1.61$ \\
\end{tabular} & \begin{tabular}{l}
Average grain diameter $21 \mu \mathrm{m}$ \\
Decrease the height of vertical \\
dendrites from $1.55 \mathrm{~mm}$ to $0.90 \mathrm{~mm}$ \\
and increase the height of directional \\
change dendrites from $0.65 \mathrm{~mm}$ to \\
$1.50 \mathrm{~mm}$ \\
\end{tabular} & \begin{tabular}{l}
Fine equiaxed \\
The bottom, middle, and top section \\
exhibits vertical columnar dendrites, \\
direction-changed columnar dendrites \\
and columnar to equiaxed transition \\
(CET) \\
\end{tabular} & \begin{tabular}{l}
$[32]$ \\
$[62]$ \\
\end{tabular} \\
\hline
\end{tabular}
\end{center}

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-11}
\end{center}

Fig. 12. Microstructures of the samples deposited by different pulse frequencies (a) $500 \mathrm{~Hz}$, (b) $100 \mathrm{~Hz}$, (c) $10 \mathrm{~Hz}$, (d) $5 \mathrm{~Hz}$, (e) $2 \mathrm{~Hz} \mathrm{and}$ (f) $1 \mathrm{~Hz}$ [32].

phases. However, the grain refinement in as-deposited samples is relatively higher than as cast. Therefore, grain refinement is the most influencing factor for slightly higher corrosion resistance [69].

The electrochemical corrosion behaviour of cast AZ31 Mg substrate and WAAMed wall is illustrated in Fig. 15. Fang et al. [59] investigated the electrochemical performance of rolled and GTAWAAMed AZ31Mg. The corrosion rate was measured to be 13.62 and $3.43 \mathrm{~mm} / \mathrm{yr}$ for substrate and WAAMed AZ31 Mg, respectively. Although the authors pointed out various factors responsible for this significant increase in corrosion resistance of the WAAMed part, the corrosion mechanism needed to be completely understood due to less investigation on the corrosion performance of WAAMed AZ31 Mg alloy. Zhang et al. [60] studied the effect of solution annealing of microstructure and corrosion performance of

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-12}
\end{center}

Fig. 13. Optical micrograph of CMT-WAAMed AZ31Mg (a) OM image of build direction (BD) (b) high magnification image at point A in Fig. (13 (a), (c) OM micrograph of travel direction (TD) (d) high magnification image of point B in Fig. (13 (c), (a1) grain morphologies and statistics of Schmid factor and misorientation angle of layer boundary of BD (b1) EBSD image of HAZ interlayer (c1) EBSD image of layer boundary of TD (d) EBSD micrograph of HAZ interlayer of TD [63].

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-13(1)}
\end{center}

Fig. 14. EBSD image of WAAMed AZ31 magnesium alloy (a) ND-TT direction (b) TD- TT direction [69].

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-13}
\end{center}

Fig. 15. (a) Potentiodynamic polarization curves (b) corresponding surface morphologies of cast AZ31 and WAAM AZ31 immersed in $0.5 \mathrm{wt} \% \mathrm{NaCl}$ solution with 30 min stabilization [69].

WAAMed AZ91 Mg. They found as the build sample exhibited $\mathrm{E}_{\mathrm{corr}}$ and $\mathrm{I}_{\text {corr }}$ of -1.522 and $3.256 \mathrm{~mA} / \mathrm{cm}^{2}$ at $0 \mathrm{~h} .-1.530,2.971 \mathrm{~mA} / \mathrm{cm}^{2}$ at $24 \mathrm{~h}$. In addition, heat-treated samples reveal the $\mathrm{E}_{\text {corr }}$ and $\mathrm{I}_{\text {corr }}$ of -1.552 and $2.751 \mathrm{~mA} / \mathrm{cm}^{2}$ at $0 \mathrm{~h}$, but at $24 \mathrm{~h}$, the $\mathrm{E}_{\text {corr }}$ and $\mathrm{I}_{\text {corr }}$ were -1.546 and $1.447 \mathrm{~mA} / \mathrm{cm}^{2}$, which is slightly decreased. This result illustrates the higher corrosion resistance of the heat-treated sample at $24 \mathrm{~h}$. This is because of grain coarsening, and large particles corresponding to the $\mathrm{Al}-\mathrm{Mg}$ phase was dissolved during heat treatment.

\subsection*{3.6. Post-processing}
In general, post-processing such as thermal (heat treatment, laser melting, laser shock peening) and mechanical (machining, burnishing, shot penning) processes can improve the properties of additive manufactured $\mathrm{Mg}$ alloy; however, minimal literature is available regarding post-processing of AMed $\mathrm{Mg}$ alloy due comparatively expensive additive manufacturing machine, and deposition of Mg alloy is relatively challenging, but since 2020 the number of publication is rapidly increasing, which indicates the importance of research in the field of $\mathrm{AM}$ of $\mathrm{Mg}$ alloy. Generally, in the post-process, defects in the materials resulting from the previous process are eliminated by post-processing and desired properties impossible during manufacturing are achieved. The mechanical properties were enhanced by post-processing due to surface modification techniques such as shock peening, shot peening, coating and burnishing. Hence, post-processing of AMed samples is mandatory to extend their application in automobile, aerospace and biomedical industries. As discussed, AM is an exemplary process for low production costs and high deposition rates. Despite these advantages, achieving desired physical and mechanical properties in AM metallic components takes time and effort.

This is because every layer of AMed wall experiences multiple thermal cycles resulting in simultaneous high heat accumulation and rapid cooling. As a result, two adverse effects are experienced by fabricated AM components (i) development of residual stress and (ii) variation in microstructure along the bottom, middle and top section, as explained elsewhere [84]. As a result of post- processing, the surface of $\mathrm{Mg}$ undergoes plastic deformation, thereby generating refined grains and developing compressive residual stress over the surface of the specimen. Interestingly, the possible defects that arise from AM are also repaired during postprocessing. In summary, post-processing is expected to (i) increase grain refinement and induce compressive residual stress. Existing post-processing techniques, which are widely used on WAAMed Mg alloy, are shown in Fig. 16.

The WAAM process results in anisotropic mechanical properties because of the different temperature gradients with increasing the number of layers. Therefore, the cold rolling process is used immediately after deposition, enhancing material homogeneity and reducing tensile residual stress during the WAAM process. Inter-pass cooling is used to minimize surface oxidation, achieve grain refinement, increase hardness and improve the strength of as-deposited components. Furthermore, inter-pass cooling reduced the dwell time between deposited layers and enhanced production efficiency. Shot peering and (abbreviated) UIT are also used to minimize the local residual stress [86]. In both processes, highenergy media is used to contact the deposited surface and impose compressive residual stress, improving surface properties such as hardness and fatigue life. Typically shot peening produced a shallow depth on the sample surface and induced compressive residual stress.

In contrast, UIT resulted in grain refinement, which enhanced the component's strength. The friction stir process is one of the surface modification techniques used to modify WAAMed parts. In this process, the grain was refined due to dynamic recrystallization that improved the mechanical properties. This process has some limitations, such as not being used for complex shaped components.

For example, Guo et al. evaluated the microstructure and mechanical properties of WAAMed AM80 M after subjecting to T4, T5 and T6 heat treatment [65], as shown in Fig. 17. Due to higher thermal stability, $\mathrm{Ca}$ and $\mathrm{Y}$ particles were not dissolved after T4 heat treatment. Ca is a microelement that can replace the $\mathrm{Mg}$ atom in the $\beta-\mathrm{Mg}_{17} \mathrm{Al}_{12}$ secondary phase, increasing the covalent and noncovelent bound strength of $\mathrm{Mg}-\mathrm{Mg}$ atoms [87]. During the $\mathrm{T} 5$ heat treatment, additional secondary phases were formed adjacent

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-14}
\end{center}

Fig. 16. Various post-processing techniques for enhancing the mechanical properties of WAAMed components [85].\\
\includegraphics[max width=\textwidth, center]{2024_04_13_7e5fbe1213c12de51fd6g-15}

Fig. 17. Microstructure of WAAMed AZ80M Mg alloy with EDS analysis (a \& b) as deposited (c \& d) T4 condition (e \& f) T5 condition and (g \& h) T6 condition [65].

to the inter-dendritic and inter-dendritic secondary phases. In contrast, the secondary phases were dissolved due to T6 heat treatment, and a fine equiaxed grain structure was obtained.

The deposited AZ80M sample exhibited anisotropic tensile properties w.r.t to the deposition and travel directions. In T4 heattreated samples, hardness $(58.78 \mathrm{HV})$ and tensile strength decreased relatively. Interestingly, during T5 and T6 heat treatment conditions, simultaneous increased tensile strength and hardness (78.4 HV) were observed [65]. From the above results, it is found that the mechanical properties of WAAMed AZ80M mg alloy decreased after $\mathrm{T} 4$ heat treatment; in contrast, they increased after T6 [63].

\section*{4. Laser powder bed fusion (LPBF)}
The LPBF is the most popular AM process among all powder bed fusion (PBF)-AM techniques. In LPBF, metal powder is melted using a laser and solidified on the substrate. This process is continued to create a track, and further, depositing $\mathrm{N}$ number of layers forms a component. For the first time, $\mathrm{Ng}$ et al. [88] deposited single tracks of pure $\mathrm{Mg}$ ( $98.4 \%$ pure) using the LPBF process to optimize process parameters in 2010 . They reported that at laser currents 25 and 28 A, no tracks were formed due to insufficient heat to melt the $\mathrm{Mg}$ powder; in contrast, at 34A laser current, they achieved deposited tracks with inconsistent beads. Zhang et al. [89] studied the bead quality and mechanical properties of LPBFed $\mathrm{Mg}-9 \% \mathrm{Al}-\mathrm{Mg}$ alloy and reported poor mechanical properties between laser power $10-20 \mathrm{~W}$ with a scan speed of $0.01-0.16 \mathrm{~m} / \mathrm{s}$. In addition, higher smoke appears with metal remnant between scan speed 30-60W laser power and evaporation of metal occurred at laser pawer $90-100 \mathrm{~W}$. However, the comparatively consistent bead occurred at laser power between 60 and 80W. Jauer et al. [90] deposited AZ91 $\mathrm{Mg}$ alloy using the LPBF process and achieved almost zero porosity. Following this, in 2014, Gieseke et al. [91] fabricated pure\\
magnesium and Mg-0.8 Ca alloy and described the effect of scanning direction on deposition quality and evaporation. This phenomenon is vividly shown in Fig. 18.1),2),3) and 4) represent the heating of top powder particles, vaporizing top powder particles with the shock wave, releasing nano-size particles and track formation with re-oxidation, respectively.

They also recommended that the scanning direction is changed for successive layers considering the excellent quality of the side surface of specimens perpendicular to the scanning direction. In 2015 Gieseke's team collaborated with Magnesium Elecktron, England and deposited WE43 and AZ91 Magnesium alloy using SLM to optimize the process parameters. They reported dendrite grain formation, coarse grain and high metal evaporation at $0.02 \mathrm{~m} / \mathrm{s}$ scan speed with $10 \mathrm{~W}$ laser power, $0.01 \mathrm{~m} / \mathrm{s}$ with $20 \mathrm{~W}$, and $0.08 \mathrm{~m} /$ min with $110 \mathrm{~W}$, respectively [93].

In 2017 Tandon et al. [94] published a review paper on SLMed Magnesium alloy. They reported the contaminants collected on powder surfaces in the form of the oxide layer lead to weaker interlayer bonding. The process parameters, namely laser power and scan time interval (STI), are responsible for oxidation and balling effects. STI denotes delay time in the SLM process, which occurs between the completion of a layer scanning and the melting of the next layer [95]. When the time delay between the scanning of the next layer is appropriated, the deposited components partially cool down, resulting in the minimum formation of oxide and better wettability. These parameters also enhanced the surface integrity by reducing the number of partially melted powder particles stuck on the surface of the deposited components [92]. Pawlak et al. [82] deposited AZ31 Mg alloy using SLM and explained the design of experiments to optimize process parameters. They also reported that the relative density and hardness of SLMed AZ31 Mg samples are relatively higher than those wrought AZ31 Mg due to rapid cooling, resulting in grain refinement and the formation of twins. However, higher material oxidation also occurred owing to excessive heat input. Salehi et al. [96] reported that magnesium powder is highly reactive, with residual oxygen contents present as an impurity in an inert gas-protective atmosphere. In addition, the authors also noted that it is challenging to deposit defect-free magnesium alloy by L-PBF because of the volatile nature of magnesium powder. Salehi et al. [96] and Zhou et al. [97] evaluated the microstructure and mechanical properties of AZ31B and $\mathrm{Mg}-\mathrm{Zn}-\mathrm{Zr}$, deposited using SLM and reported that vaporization of $\mathrm{Mg}$ alloy is inevitable. Due to high energy input and rapid cooling, the fine grain was obtained in the centre of the track, and coarse grains were formed at the boundary of the track $[83,64]$. Another essential parameter to be considered is the preheating of powder that assists in producing flat and regular tracks [98]. In the L-PBF process, major defects are improper powder melting, oxidation and evaporation, and residual stress. Due to Mg's low heat absorption capacity, most of the laser energy is reflected, which leads to improper melting of Mg powder. Indeed, when the laser density increases, the melt pool temperature exceeds $1120^{\circ} \mathrm{C}$, significantly higher than Mg's boiling point. It is well established that the vapour pressure of $\mathrm{Mg}$ alloy is $0.13 \mathrm{KPa}$ at $650{ }^{\circ} \mathrm{C}$. At the melt pool temperature of $1120{ }^{\circ} \mathrm{C}$, vapour pressure increases up to $51 \mathrm{KPa}$, eventually initiating and assisting the evaporation of Mg. Now researchers and industries concentrate on SLM because it produces high-precision components and fabricates highly complex shapes. However, the SLM additive manufacturing process is not recommended for producing large magnesium components due to its low deposition rate and limited build chamber size. In addition, powder preparation of Magnesium alloy is quite challenging, and these powders quickly oxide resulting in combustion and explosion. Therefore, in recent years the attention of researchers has been shifting towards Wire Arc Additive Manufacturing (WAAM) for the deposition of various grades of magnesium alloy.

\subsection*{4.1. Magnesium alloy system deposited using LPBF}
Compared to studies on cast and wrought $\mathrm{Mg}$ alloy manufacturing, few of them have been dedicated to additive manufacturing of $\mathrm{Mg}$ alloy. This is because of the high cost associated with the customized production of atomized pre-alloyed powder, which is hundreds of times higher than the cost of customization for cast and wrought Mg alloys [99]. In present-day pure $\mathrm{Mg}, \mathrm{AZ91}, \mathrm{AZ31}$ and WE43 Mg alloys are predominantly used in various automotive, aerospace and biomedical applications. This is due to their massive market demand, good printability using LPBF and WAAM, and obtained suitable mechanical properties for engineering structures and biomedical implants. AMed Mg alloys are not commercialized yet. However, in the near future, it will be commercialized because many laboratories and research institutes worldwide are studying additive manufacturing of $\mathrm{Mg}$ alloy to understand the challenges during deposition [100].

\subsection*{4.1.1. Pure $m g$}
Ng et al. [88] from Hon Kong Polytechnic University deposited pure Mg in 2010 using the LPBF process with Nd: YAG laser. This study used various scan speeds and laser power to deposit single tracks to optimize process parameters. The irregular and coarse

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-16}
\end{center}

Fig. 18. Evaporation during deposition of Mg alloy using LPBF [92].\\
powder size resulted in improper tracks. Interestingly soundtracks were successfully deposited with atomized fine and spherical shape powder under appropriate conditions. The grain size of LPBFed pure Mg is around 2-5 $\mu \mathrm{m}$ [101]. Previously Morishige et al. [102] achieved an $8-10 \mu \mathrm{m}$ grain size of pure Mg, processed by SPD. These studies confirmed that LPBF has the potential for grain refinement, a key advantage over the conventional manufacturing process. The study of pure $\mathrm{Mg}$ deposited using LPBF also reported the hardness value lies between 60 and $89 \mathrm{HV}(0.59-0.87 \mathrm{MPa})$ [101]. In addition, high oxidation and cracks also occurred along the grain boundaries due development of tensile residual stress and distortion. The first thin and thick wall of pure $\mathrm{Mg}$ was deposited with powder sizes 26 and $43 \mu \mathrm{m}$ using LPBF by Hu et al. [95] from Chongqing University. They reported insufficient energy density could not be used to print thin, thick walls of $\mathrm{Mg}$ alloy.

On the other hand, the high energy density resulted in severe evaporation of pure Mg. Apart from LPBF, one more Additive manufacturing technique, the DLD AM process, is used to deposit pure Mg with irregular and coarse powder particles as feedstock materials, resulting in high porosity, cracks and poor surface finish $[84,85]$.

\subsection*{4.1.2. Magnesium-Aluminum ( $\mathrm{Mg}-\mathrm{Al}$ )}
AZ31 $\mathrm{Mg}$ alloy is the most useable commercial composition in the form of cast and wrought among all $\mathrm{Mg}-\mathrm{Al}-$ series $\mathrm{Mg}$ alloys. However, in the field of LPBF, a few studies have been carried out on AZ31Mg alloy. Most studies on additive manufacturing of AZ31 Mg alloy are based on the wire arc additive manufacturing process. In fact, AZ91 Mg is mainly deposited using LPBF because Aluminium (Al) provides solid solution strengthening. In addition, the secondary phase $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ improves printability and also plays a vital role in grain refinement by superheating [103]. However, Pawlak et al. deposited AZ31 Mg alloy using the LPBF process and reported a very low porosity level of $>0.5 \%$ [104]. Hence in the studies of AZ61 and AZ91 processed using LPBF, high density was also achieved, demonstrating the printability and acceptability of the $\mathrm{Mg}-\mathrm{Al}$ system. The results of $\mathrm{AZ61}$ [105] and AZ91 represent the fine and equiaxed grain size of around 1-3 $\mu \mathrm{m}$ with random texture, vividly shown in Fig. 19 (a). The secondary phase $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ is uniformly distributed along the grain boundary, interconnecting with each other, which is depicted in Fig. 19 (b) [87,88]. Fig. 19 (c) shows the elongated grain boundary along the build direction. Moreover, the $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ intermetallic phase is distributed along the grain boundary, and some of the high density of sphericalshaped $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ with $100-300 \mathrm{~nm}$ secondary phase nanoparticle is found inside the grain, as vividly depicted in Fig. 19 (d). The section summarises that the microstructure of the $\mathrm{Mg}-\mathrm{Al}$ system deposited using LPBF is tunable by adjusting the process parameters.

\subsection*{4.1.3. Magnesium-rare earth ( $M g-R E)$ based alloy system}
Among all the Mg-RE alloy systems, WE43 Mg alloy is most extensively deposited by the LPBF process for biomedical implant applications. As mentioned above, it has outstanding printability and a large processing window to achieve high density than AZ91Mg alloy. Moreover, Al is a neurotoxic element and is not recommended for biomedical application of fear of Alzheimer's disease. Therefore, WE43 Mg alloy is biocompatible and used in scaffolds, compressive plates and screws. Zumdick et al. [107] studied the microstructure of LPBFed WE43 Mg alloy and observed refined equiaxed grain, as shown in Fig. 20 (a); however, some of the abnormal grain growth occurred due to huge temperature difference between the bottom and top layer of deposited parts. The size of refined equiaxed grains was 1-3 $\mu \mathrm{m}$ with random orientation. Interestingly Bar et al. studied the LPBFed WE43 in 2019 and found irregular firm and large basal textured of size $20.4 \pm 6.3 \mu \mathrm{m}$, clearly shown in Fig. 20 (b) [108]. From this figure, equiaxed grain of $4.7 \pm 0.4 \mu \mathrm{m}$ exhibiting random texture was

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-17}
\end{center}

Fig. 19. Microstructure of LPBFed AZ91 Mg alloy (a) fine equiaxed grain [106] (b) manganese rich precipitate [106] (c) EBSD Map, grain orientation [19] (d) spherical Mg ${ }_{17} \mathrm{Al}_{12}$ nanoparticle [19].

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-18}
\end{center}

Fig. 20. Grain orientation Map of LPBFed WE43 Mg alloy (a) fine, equiaxed, and randomly orientated [111] (b) fine, equiaxed, and randomly orientated grains in the last melt pool [108] (c) large, irregular-shape, and basal-orientated grains [17] (d-e) EDXS maps acquired from the same material [17].

observed only at the centre of the last melt pool. Columnar grains were found only between the equiaxed and large irregular grains. During the successive solidification of the melt pool, equiaxed grain nucleate in partially cooled metal liquid ahead of the columnar grain region and start to transition from columnar to equiaxed [109]. From Fig. 20 (b), it is observed that the columnar grain with strong basal texture elongated along the build direction in region II and irregular grains with strong basal texture in region III. The same observation is also reported by Esmaily et al. [17]. In addition, they also noted that the deposition of WE43 on a wide range of energy density $\left(120-300 \mathrm{~J} / \mathrm{mm}^{3}\right)$ and strong-basal texture occurred for all, as shown in Fig. 20 (c). Fig. 20 (d \& e) shows the EDXS mapping at two magnifications on obtained microstructure. The coarse flake contained Zr, Nd and Y, vividly demonstrated in Fig. 20 (d). While the finer flakes oxides contained $\mathrm{Zr}$ and $\mathrm{Y}$ without $\mathrm{Nd}$, as shown in Fig. 20 (e). Despite the high density of RE oxides, the large and basal-orientated grains are still dominant in the LPBFed WE43. In other studies, a research team from Shanghai Jiaotong University customized the WE43 Mg alloy powder composition and developed an Mg-Gd alloy system [110,111]. After that, they deposited MgGd-based alloy using LPBF and observed equiaxed grain size around 1-2 $\mu \mathrm{m}$ with random orientations. They also concluded minimal oxide, less porosity and a high relative density of around 99.95\% were obtained [112]. Liao et al. [110] deposited Mg-Gd Rare earth Mg alloy (Mg-10Gd-3Y-0.4Zr) alloy using a DLD process with coarse spherical powder size around 100-300 $\mu \mathrm{m}$. They observed a randomly oriented equiaxed average grain size of $19 \mu \mathrm{m}$ with a higher fraction of pores. Interestingly both LPBF and DLD build MgGd-based alloy exhibited superior quality. However, the preferential grain growth of basal-oriented grains depends on the amount of Gd element, which should be $>10 \mathrm{wt} \%$ during deposition of both methods (DLD \& LPBF) [113].

\subsection*{4.1.4. Magnesium-Zinc ( $M g-Z n)$ based alloy system}
$\mathrm{Mg}-\mathrm{Zn}$-based alloys are less explored in $\mathrm{AM}$ than $\mathrm{Mg}-\mathrm{Al}$ and $\mathrm{Mg}$ - RE-alloys. This is attributed to the poor printability of $\mathrm{Mg}-\mathrm{Zn}$ resulting from the significantly lower eutectic temperature of $325^{\circ} \mathrm{C}$ and a large solidification range [114]. Wei et al. reported that less porosity was achieved at a $\mathrm{Zn}$ concentration of $\leq 1 \mathrm{wt} \%$ and high porosity at $\geq 12 \mathrm{wt} \%$; nevertheless, when $\mathrm{Zn}$ concentration is in the middle, i.e., $6 \mathrm{wt} \%$ (the Zn concentration in commercial ZK60 wrought alloy is $6 \mathrm{wt} \%$ ), high porosity and severe cracks on LPBF $\mathrm{Mg}-\mathrm{Zn}$ alloy made it unacceptable for application. The highest relative density of $97 \%$ was achieved in LPBFed ZK60, which is relatively higher than reported so far at $94 \%$ [115]. Therefore, this demonstration confirmed that $\mathrm{Zn}$ should be used as a minor element in $\mathrm{Mg}-\mathrm{Zn}$-based alloy for deposition using the LPBF process. Other than this study, magnesium- Tin ( $\mathrm{Mg}-\mathrm{Sn})$ based alloy [116] and Magnesium-calcium ( $\mathrm{Mg}-\mathrm{Ca}$ ) based alloy [117] are also deposited using LPBF and investigated. The printability of $\mathrm{Mg}-\mathrm{Sn}$ and $\mathrm{Mg}-\mathrm{Ca}$ alloy is relatively better than $\mathrm{Mg}-\mathrm{Zn}$ alloy. They exhibit high eutectic temperatures $\left(510{ }^{\circ} \mathrm{C}\right.$ for $\mathrm{Mg}-\mathrm{Ca}$ and $466{ }^{\circ} \mathrm{C}$ for $\mathrm{Mg} \mathrm{Sn}$ ) and a low solidification range. Both alloys' preliminary results are insufficient to understand the mechanism involved in LPBF thoroughly. Therefore, comprehensive studies are required to realize the microstructure, mechanical properties and corrosion behaviour of LPBFed Mg alloy.

\subsection*{4.2. Microstructure and mechanical properties}
Zhang et al. [89] have selectively melted 56 samples of Mg-9 wt.\% Al alloy by varying laser power and scan speeds using SLM. The process maps revealed that laser power of 10,15 and $20 \mathrm{~W}$ and a scan speed of $0.01,0.02$ and $0.04 \mathrm{~m} / \mathrm{s}$ yielded high-quality specimens. The microstructure and mechanical properties of the\\
samples at four different scanning speeds, viz., 500, 750, 1000 and $1250 \mathrm{~mm} / \mathrm{s}$, were evaluated. The results revealed cellular structures of the $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ secondary phase around primary $\mathrm{Mg}$ columnar grains, as shown in Fig. 21 (a). The addition of SiCnp columnar to equiaxed transition occurred due to heterogenous nucleation and controlled diffusion. As a result, fine equiaxed grains of 1-2 $\mu \mathrm{m}$ were achieved, as shown in Fig. 21 (b c \& d). Interestingly 5\% SiCnp addition exhibited relatively better yield strength, attributed to better bonding strength between SiCnp and AZ91D Mg matrix. In summary, the addition of SiCnp has refined the grain and homogenized the overall microstructure of bulk specimens.

Moreover, with the increased scanning rate from 500 to $1250 \mathrm{~mm} / \mathrm{s}, \mathrm{Mg}$ powders remained unmelted. This is related to sintering rather than melting due to insufficient energy input. As a result of the un-melting of $\mathrm{Mg}$ powders, the mechanical properties deteriorated significantly with increased scanning speed [118].

Similarly, the element vaporization of $\mathrm{Mg}-\mathrm{Zn}-\mathrm{Zr} \mathrm{Mg}$ alloy during SLM was studied by Wei et al. [119] due to element vaporization induced by $\mathrm{Zr}$, the Wt. \% of SLMed parts obtained at scanning speed viz., $300-900 \mathrm{~mm} / \mathrm{s}$ decreased relatively compared to the ZK60 powders. Consequently, the relative density dropped drastically from $94 \%$ for a scanning speed of $300 \mathrm{~mm} / \mathrm{s}$ to $82 \%$ at $900 \mathrm{~mm} / \mathrm{s}$ [119]. $\mathrm{Mg} 0.5 \mathrm{Zn}$ and $\mathrm{Mg} 1 \mathrm{Zn}$ were selectively melted using the Mg and $\mathrm{Zn}$ powder feedstock material. The mechanical properties of LPBFed $\mathrm{Mg}-\mathrm{Zn}$ alloys were comparable to that of as- cast Mg alloys. In addition, visible keyhole pores were observed due to evaporation and incomplete melting of $\mathrm{Mg}$ and $\mathrm{Zn}$ powders [120]. The microstructure and mechanical properties of L-PBF Mg alloys are listed in Table 3.

\subsection*{4.3. Corrosion behaviour of LPBFed Mg alloy}
ZK60 Mg powders were selectively laser melted at four different volume energy densities viz $420,500,600$ and $750 \mathrm{~J} / \mathrm{mm}^{3}$. The samples deposited at $600 \mathrm{~J} / \mathrm{mm}^{3}$ exhibited relatively better relative density and microhardness of $97 \%$ and $89.2 \mathrm{HV}$, respectively. In addition, this condition of ZK60 samples evinced a corrosion rate of $0.006 \mathrm{ml} / \mathrm{cm}^{2} \mathrm{~h}$ during immersion in Hank's solution, which is significantly lower when compared to other SLMed ZK60 samples [123]. The better mechanical properties and corrosion resistance obtained in the SLM ZK60 sample for $600 \mathrm{~J} / \mathrm{mm} 3$ of Ev are attributed to the fine grain microstructure, and homogeneous nucleation of $\mathrm{Mg}$ received resulted from rapid solidification [124]. $\mathrm{Mg} 0.5 \mathrm{Zn}$ and $\mathrm{Mg} 1 \mathrm{Zn}$ samples deposited using the LPBF process exhibited a corrosion rate of $\sim 0.18 \mathrm{ml} / \mathrm{cm}^{2} / \mathrm{h}$. The corrosion resistance of $\mathrm{Mg}-\mathrm{Zn}$ samples was significantly lesser compared to the most commonly employed $\mathrm{Mg}-\mathrm{Y}-\mathrm{RE} \mathrm{Mg}$ alloy system in biomedical applications [125]. Liu et al. [126] manufactured biodegradable WE43 Mg scaffolds comprising diamond-shaped unit cells as building blocks using the LSM technique. The bottom and top\\
\includegraphics[max width=\textwidth, center]{2024_04_13_7e5fbe1213c12de51fd6g-19}

Fig. 21. Typical microstructures of the LPBF-processed specimens (scanning speed $=200 \mathrm{~mm} / \mathrm{s}$ ) [118].

Table 3

Microstructure and mechanical properties of various grades of magnesium alloy powder processed by LPBF.

\begin{center}
\begin{tabular}{|c|c|c|c|c|c|c|}
\hline
Mg Alloy & UTS (MPa) & YS (MPa) & Microstructure & \begin{tabular}{l}
Hardness \\
$(\mathrm{Hv})$ \\
\end{tabular} & \begin{tabular}{l}
Elongation \\
$(\%)$ \\
\end{tabular} & Ref. \\
\hline
AZ31B & $207 \pm 5$ & $183 \pm 3$ & Non-uniform fine $\alpha$-Mg, equiaxed and elongated grain formed with $\gamma \mathrm{MgMg}_{17} \mathrm{Al}_{12}$ eutectic phase & $64 \pm 0.5$ & 7.7 & $[82]$ \\
\hline
AZ91D & $274 \pm 16$ & $237 \pm 17$ & $\alpha-\mathrm{Mg}$ Equiaxed grain formed with $\beta-\mathrm{Mg}_{17} \mathrm{Al}_{12}$ eutectic phase & $85 \pm 0.2$ & 3 & $[121]$ \\
\hline
AZ61 & $239.3 \pm 20$ & $216 \pm 17$ & \begin{tabular}{l}
$\alpha-\mathrm{Mg}$ Equiaxed grain and entirely divorced $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ eutectic phase distributed along the grain \\
boundary \\
\end{tabular} & - & $2-3$ & $[105]$ \\
\hline
$\mathrm{Mg}$ & - & - & Size of the microstructure increase with laser density & $0.59 \mathrm{Gpa}$ & - & $[93]$ \\
\hline
WE43 & 312 & 194 & Elongated Equiaxed fine-grain formed on the melt pool boundaries &  & 14 & $[122]$ \\
\hline
ZK60 &  &  & \begin{tabular}{l}
When increasing the energy density, dendrites to columnar grain formed with $\beta M g_{7} \mathrm{Zn}_{3}$ \\
precipitation \\
\end{tabular} & $70.1-85$ &  & $[92]$ \\
\hline
$\mathrm{Mg}-9 \mathrm{Al}$ & - & - & $\alpha$-Mg equiaxed grain with $\mathrm{Mg}_{17} \mathrm{Al}_{12}, \mathrm{MgO}$ and $\mathrm{Al}_{2} \mathrm{O}_{3}$ precipitation found at the grain boundaries & $66-85$ & - & $[89]$ \\
\hline
\end{tabular}
\end{center}

sections of struts displayed rose-like grains, and cellular grain morphology corroborated moderate and high-temperature gradients. In addition, the $\mathrm{Y}_{23}^{\mathrm{O}}$ phase was uniformly distributed in a bulk Mg matrix within each melt pool, owing to Marangoni convection. Biomechanical properties revealed that compressive yield strength (CYS) decreased chronologically from $22 \mathrm{MPa}$ to $13 \mathrm{MPa}$ when tested after 7- and 24-days immersion [127]. The AMed WE43 samples exhibited level zero and level one cytotoxicity corresponding to 24,48 and $72 \mathrm{~h}$ of immersion time. The corrosion rate (mm/year) of various LPBFed, Mg alloy grades is shown in Fig. 22. The corrosion rate of multiple grades of $\mathrm{Mg}$ alloy obtained by different processes such as coating, burnishing, casting WAAM and SLM process is depicted in Fig. 23. From this figure, it is clear that incorporating coatings and WAAM has reduced the corrosion rate. It can also observe that phosphate-based biocompatible coatings effectively reduce the degradation rate significantly.

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-20}
\end{center}

Fig. 22. The corrosion rate of various grades of LPBFed mg alloy [103-105].

\subsection*{4.4. Post-processing}
Generally, some major concerns, such as the development of residual stress, oxidation, vaporization and under-melting, occurred during the LPBF process. These defects deteriorate the mechanical properties. Benn et al. [128] evaluated the corrosion behaviour and cytocompatibility of printed, etched and machined WE43Mg samples processed by LPBF. The hydrogen evolution rate of these samples after 20 days of immersion was found to be $\sim 140, \sim 68$ and $\sim 85 \mathrm{~mL} / \mathrm{cm}^{2}$, respectively. While the \% relative to toxic control tested by lactate dehydrogenase (LDH) assay of the samples were $\sim 76, \sim 48$ and $\sim 47$, respectively. It is observed that despite the significant increase in corrosion resistance of post-processed samples, their biocompatibility is inferior to that of printed samples. Esmaily et al. [17] evaluated the microstructure and mechanical properties of as-built and HIP-processed WE43 Mg alloy. Star-shaped Mg-rich particles and Y and Zr-rich particles were observed in as-built and HIP conditions. In addition, the $\mathrm{Mg}_{24} \mathrm{Y}_{5}$ phase present in as-build WE43 dissolved due to HIP treatment, along with the appearance of $\mathrm{Mg}_{41} \mathrm{Nd}_{5}$ confirmed by XRD, SEM-EDS and STEM-EDS. Interestingly, HIP treatment significantly reduced the porosity of as-built WE43 Mg. However, the dynamic mechanical behaviour of as-built and HIP-treated WE43 Mg were comparable when tested using Split Hopkinson's pressure bar apparatus. Esmaily et al. subjected SLMed WE43 Mg alloy to postprocessing techniques such as hot isostatic pressing (HIP) and heat treatment (HT). EBSD micrographs revealed firm preferential basal solid texture (0001) for SLM, SLM + HIP and SLM + HIP + HT WE43 Mg samples along the build direction vividly shown in Fig. 24. In addition, due to the higher solidification rate, a common trait of the SLM process, sub-micron cellular structures were observed in WE43 Mg. Hot isostatic pressing of SLMed WE43 Mg alloys successfully reduced the porosity. As a result, the corrosion resistance of SLM + HIP was relatively higher than the SLMed sample. However, the corrosion resistance of SLMed and postprocessed SLMed samples is significantly inferior compared to cast WE43 Mg [129].

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-20(1)}
\end{center}

Fig. 23. Improving the corrosion rate of various grades of Mg alloy obtained by different processes. MEM represents the Minimum Essential Medium solution, SFD represents the simulated body fluid solution, and DMEM represents Dulbecco's modified Eagle Medium solution.

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-21}
\end{center}

Fig. 24. EBSD analysis of the samples in the as-SLMed, SLM + HIP and SLM + HIP + HT conditions (1) at $20 \mathrm{~W}$ with $200 \mathrm{~mm} / \mathrm{s}$ process speed (2) at $80 \mathrm{~W}$ with $400 \mathrm{~mm} / \mathrm{s}$ [process speed (3) $120 \mathrm{~W}$ with $600 \mathrm{~mm} / \mathrm{s}$ process speed (4) $160 \mathrm{~W}$ with $800 \mathrm{~mm} / \mathrm{s}$ process speed (5) $200 \mathrm{~W}$ with $1000 \mathrm{~mm} / \mathrm{s}$ process speed (6) $240 \mathrm{~W}$ with $1200 \mathrm{~mm} / \mathrm{s}$ process speed (7) at $280 \mathrm{~W}$ with $1400 \mathrm{~mm} / \mathrm{s}$ process speed [17].

Zumdick et al. [107] investigated the mechanical properties of additively manufactured, powder extruded (PE) and as-cast WE43 Mg. Lueder bands were observed along the length of AM and PE WE43 Mg samples, eventually resulting in the yield point phenomenon. Indeed, a fine microstructure of about $\sim 1 \mu \mathrm{m}$ was responsible for significantly higher mechanical properties of AM and PE samples' than the as-cast WE43 Mg [130]. Salehi et al. [96] developed green ZK60 compacts using capillary-mediated 3D printing followed by sintering. The sintering temperature of $573{ }^{\circ} \mathrm{C}$ and holding time of $60 \mathrm{~h}$ was optimum to achieve the highest values of ultimate compressive strength (UCS) $174.30 \mathrm{MPa}$, compressive yield strength CYS $39.87 \mathrm{MPa}$, and $33.51 \%$ elongation. Kopp et al. [131] fabricated WE43 Mg using LPBF and postprocessed it with additional heat treatment and PEO coating to improve the biomechanical properties. PEO coating substantially enhanced compressive strength and corrosion resistance compared to the uncoated samples. In contrast, heat treatment resulted only in marginal improvement of the biomechanical properties of LPBF Mg alloy.

\subsection*{4.5. Comparison of WAAMed and LPBFed Mg alloy}
This review article explains the microstructure, mechanical properties, corrosion performance and post-processing of $\mathrm{Mg}$ alloy deposited using wire arc additive manufacturing (WAAM) and laser powder bed fusion (LPBF) processes. In both approaches, the nature and distribution of heat sources are entirely different. In the WAAM process, heat is generated in the form of the arc due to an electrical short circuit; therefore, the distribution of the heat source is unsymmetric and can be predicted by a double ellipsoidal heat source model [132], while in the LPBF heat produced by solid laser (Nd-YAG), the distribution of laser heat is symmetric and can be predicted using Gaussian heat source model [133]. Consequently, the difference in the cooling rate of both processes is significant during deposition, which can influence the type and size of the grains. Therefore, the mechanical properties and corrosion performance of $\mathrm{Mg}$ alloy, deposited by both processes, are not even similar, but whole properties, such as thermal, chemical and mechanical, are different [134]. Therefore, the next section of this\\
article explores the differences between WAAMed and LPBFed Mg specimens in terms of microstructural, mechanical properties and corrosion performance.

\subsection*{4.5.1. Microstructural and mechanical properties and corrosion performance}
The microstructure induced by both processes, WAAM and LPBF, are different owing to the nature of the heat source and cooling rate. The LBPFed Mg alloy mainly exhibited finer grains due to a high cooling rate of approximately $40 \mathrm{~K} / \mu \mathrm{s}$ [135]. However, the grain size increases from the bottom to the top section of the thin wall [136]. This is because heat accumulation due to preheating the previous layer reduced the thermal gradient, which is responsible for the low cooling rate. Consequently, grain size increases towards to deposition direction [137]. Massive grain size variation affects the mechanical properties of LPBFed Mg alloy components. The microstructure depends on process parameters and the environmental condition of the LPBF build chamber [138]. Generally, an AZ series (AZ31, AZ80M AZ91 of Mg alloy, deposited using the LPBF process, exhibited the massive secondary phase particles such as $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ [118] and in most studied Mg alloy WE43 composed of $\mathrm{Mg}_{41} \mathrm{Nd}_{5}, \mathrm{Mg}_{3} \mathrm{Nd}$ particles [108] with globular shape in the considerable amount owing to variation thermal history during deposition, which is comparatively brittle, reduced the ductility of $\mathrm{Mg}$ alloy. After analysis of the microstructure of LPBFed Mg alloy, it is noted that no literature shows the equiaxed microstructure.\\
However, in the WAAMed Mg alloy, maximum literature showed the negligible formation of brittle $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ secondary phase particles with complete equiaxed grains, which can enhance the ductility of deposited components [58]. Therefore, the WAAMed \% EL is significantly higher than LPBF Mg specimens, as shown in Fig. 25.

In the wire arc additive manufacturing field, most AZ series of $\mathrm{Mg}$ alloy are deposited using a cold metal transfer-based WAAM process. This is because of comparatively lower heat input which is suitable for low melting point temperature $\mathrm{Mg}$ alloy. Due to low heat input and comparatively higher thermal gradient, the grains are homogeneous, which is beneficial for mechanical properties. However, from the literature survey, it is noted that only CMTWAAM [69], GTA-WAAM [59] and UPF-WAAMed [58] Mg specimens exhibited equiaxed grain with negligible brittle secondary phase particles. But GTAW, MIG, GMAW based WAAM specimens showed vertical columnar dendrites; the direction changed columnar dendrites, and equiaxed dendrites in sequence; their heights are approximately $0.90 \mathrm{~mm}, 1.50 \mathrm{~mm}$ and $2.20 \mathrm{~mm}$. This is because of high heat input [69].

Moreover, the inhomogeneous microstructure and massive formation of secondary phase particles are caused for lower mechanical properties in LBPFed alloy. However, the smaller grains and secondary phase particles enhanced the corrosion resistance of $\mathrm{Mg}$ alloy. Therefore, the corrosion resistance of LPBFed Mg alloy is higher than that of WAAMed Mg specimen, as seen in Table 4.

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-22}
\end{center}

Fig. 25. Comparison of mechanical properties between (a) WAAMed Mg alloy (b) LPBFed Mg alloy.

Table 4

Summary of corrosion properties of LPBFed and WAAM Mg alloy.

\begin{center}
\begin{tabular}{|c|c|c|c|c|c|c|}
\hline
Mg Alloys & Deposition process & $\mathrm{E}_{\text {corr }}(\mathrm{V})$ & $\mathrm{I}_{\text {corr }}\left(\mathrm{mA} / \mathrm{cm}^{2}\right)$ & Solution & Corrosion rate (Pw, mpy) & Refs. \\
\hline
Pure Mg & LPBF (SS:500 mm/s) & -1.52 & $74.0 \times 10^{-3}$ & HBSS & $2.8(\mathrm{~mm} / \mathrm{y})$ & [139] \\
\hline
Pure $\mathrm{Mg}$ & LPBF (SS:750 mm/s) & -1.53 & $114.1 \times 10^{-3}$ & HBSS & $4.7(\mathrm{~mm} / \mathrm{y})$ & [139] \\
\hline
Pure Mg & LPBF (SS:1000 mm/s) & -1.54 & $162.1 \times 10^{-3}$ & HBSS & $9.6(\mathrm{~mm} / \mathrm{y})$ & [139] \\
\hline
Pure Mg & LPBF (SS: $1250 \mathrm{~mm} / \mathrm{s})$ & -1.53 & $176.6 \times 10^{-3}$ & HBSS & $32.5(\mathrm{~mm} / \mathrm{y})$ & [139] \\
\hline
AZ61 & $\mathrm{SLM}(\mathrm{SS}=22 \mathrm{~mm} / \mathrm{s})$ & $-1.50 \pm 0.02$ & $59 \times 10^{-3}$ & SBF & 1.45 & [140] \\
\hline
AZ61-1.0 RGO/MgO & $\mathrm{SLM}(\mathrm{SS}: 22 \mathrm{~mm} / \mathrm{s})$ & $-1.47 \pm 0.02$ & $199 \times 10^{-3}$ & SBF & 2.19 & [140] \\
\hline
AZ61-2.0 RGO/MgO & $\mathrm{SLM}(\mathrm{SS}: 22 \mathrm{~mm} / \mathrm{s})$ & $-1.48 \pm 0.03$ & $132 \times 10^{-3}$ & SBF & 1.99 & [140] \\
\hline
AZ61-3.0 RGO/MgO & SLM (SS:22 mm/s) & $-1.48 \pm 0.05$ & $42 \times 10^{-3}$ & SBF & 1.05 & [140] \\
\hline
AZ61-4.0 RGO/MgO & SLM (SS:22 mm/s) & $-1.48 \pm 0.02$ & $105 \times 10^{-3}$ & SBF & 1.32 & [140] \\
\hline
AZ61-0.6GO & SLM (SS:15 mm/s) & $-1.57 \pm 0.02$ & $(118 \pm 13) \times 10^{-3}$ & SBF & $2.67 \pm 0.30(\mathrm{~mm} / \mathrm{y})$ & [141] \\
\hline
AZ61 & SLM (SS:15 mm/s) & $-1.54 \pm 0.02$ & $(50 \pm 4) \times 10^{-3}$ & SBF & $1.21 \pm 0.09(\mathrm{~mm} / \mathrm{y})$ & [141] \\
\hline
AZ61-0.2GO & SLM (SS:15 mm/s) & $-1.54 \pm 0.02$ & $(89 \pm 12) \times 10^{-3}$ & SBF & $2.03 \pm 0.27(\mathrm{~mm} / \mathrm{y})$ & [141] \\
\hline
AZ61-0.4GO & SLM (SS:15 mm/s) & $-1.52 \pm 0.03$ & $(212 \pm 16) \times 10^{-3}$ & SBF & $4.84 \pm 0.36(\mathrm{~mm} / \mathrm{y})$ & [141] \\
\hline
AZ91D & $\mathrm{BJ}+\mathrm{TTS}-2\left(680^{\circ} \mathrm{C}\right)$ & -1.535 & $79.8 \times 10^{-3}$ & $3.5 \%$ wt. $\mathrm{NaCl}$ & $120 \pm 12.8(\mathrm{~mm} / \mathrm{y})$ & [142] \\
\hline
AZ91D & $\mathrm{BJ}+$ TTS $-1\left(660^{\circ} \mathrm{C}\right)$ & -1.599 & $149 \times 10^{-3}$ & 3.5\%wt. $\mathrm{NaCl}$ & $172 \pm 17.6(\mathrm{~mm} / \mathrm{y})$ & [142] \\
\hline
ZK30 & SLM (SS:500 mm/min) & -1.57 & 0.10 & SBF & $3.70 \pm 0.10$ & [143] \\
\hline
ZK30-0.3GO & SLM (SS:500 mm/min) & -1.59 & 0.33 & SBF & $10.8 \pm 0.09$ & [143] \\
\hline
ZK30-0.6GO & SLM (SS:500 mm/min) & -1.65 & 0.10 & SBF & $3.38 \pm 0.07$ & [143] \\
\hline
ZK30-0.9GO & SLM (SS:500 mm/min) & -1.51 & 0.49 & SBF & $15.64 \pm 0.13$ & [143] \\
\hline
AZ31 & Hot-rolled & -1.56 & 611.62 & $3.5 \mathrm{wt} \% \mathrm{NaCl}$ & $13.62(\mathrm{~mm} / \mathrm{y})$ & [59] \\
\hline
AZ31 & WAAM-GTA & -1.60 & 154.12 & $3.5 \mathrm{wt} \% \mathrm{NaCl}$ & $3.43(\mathrm{~mm} / \mathrm{y})$ & [59] \\
\hline
AZ91 & WAAM & -1.522 & 3.256 & $0.1 \mathrm{M} \mathrm{NaCl}$ & $0.73(\mathrm{~mm} / \mathrm{y})$ & [60] \\
\hline
AZ91 & WAAM + HT & -1.552 & 2.751 & $0.1 \mathrm{M} \mathrm{NaCl}$ & $0.62(\mathrm{~mm} / \mathrm{y})$ & [60] \\
\hline
AZ91 & WAAM $(24 \mathrm{~h})$ & -1.530 & 2.971 & $0.1 \mathrm{M} \mathrm{NaCl}$ & $0.57(\mathrm{~mm} / \mathrm{y})$ & [60] \\
\hline
AZ91 & WAAM + HT (24h) & -1.546 & 1.447 & $0.1 \mathrm{M} \mathrm{NaCl}$ & $0.42(\mathrm{~mm} / \mathrm{y})$ & [60] \\
\hline
AZ31 & CMT-WAAM & -1.555 & 0.1 & $0.5 \mathrm{wt} \% \mathrm{NaCl}$ & $2.28(\mathrm{~mm} / \mathrm{y})$ & [69] \\
\hline
\end{tabular}
\end{center}

Fig. 26 shows the comparison of deposited Mg alloy using LBPF and WAAM process and processed LPBFed and WAAMed Mg alloy. This figure shows that the deposited AZ91 Mg alloy specimens using the WAAM process exhibited a colossal amount of secondary phase particles such as $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ and $\mathrm{Al}_{8} \mathrm{Mn}_{5}$; after heat treatment, most of these brittle phases dissolved in the primary phase with uniform distribution. Consequently, the strength and elongation \% improved. In contrast, in SLMed AZ61 specimens, columnar grains with $\alpha-\mathrm{Mg}$ and $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ phases were. In addition, HIP post-process induced grain growth of SLMed AZ61 $\mathrm{Mg}$ alloy at $450{ }^{\circ} \mathrm{C}$ under $103 \mathrm{MPa}$ for $3 \mathrm{~h}$ [144], slightly reducing the yield strength and hardness. However, the dissolution of the $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ phase along the grain boundaries and reduction of pores, as depicted in Fig. 26 (b$\mathrm{b}_{2}$ ), enhanced the total elongation without changing tensile strength.

Ghorbani et al. [146] investigated the dissolution of $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ at $45{ }^{\circ} \mathrm{C}$ based on the JMatPro plot of AZ61 Mg alloy. They reported that the $\mathrm{Mg}_{17} \mathrm{Al}_{12}$ phase completely dissolved at a temperature

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-23}
\end{center}

Fig. 26. Comparison of tensile strength of ( $\left.a-a_{2}\right)$ WAAM Mg alloy (b-b $\left.b_{2}\right)$ LPBF Mg alloy [145].\\
higher than $308^{\circ} \mathrm{C}$ which is identified by the binary phase diagram of $\mathrm{Mg}-\mathrm{Al}$, that exhibited $6 \mathrm{wt} \% \mathrm{Al}$ at the temperature of $450{ }^{\circ} \mathrm{C}$ corresponds to single phase $\alpha-\mathrm{Mg}$, as can be seen in Fig. 27.

In summary, the HIP post-process treatment minimizes the internal defect of additively manufactured components. In the HIP treatment process, inert environments' high pressure and temperature allow materials to deform and collapse the pores. Still, if gas is soluble, it can diffuse out of the pores [147]. Therefore, the extended time high temperature is the cause for microstructure coarsening, which resulted in low mechanical properties and unfavourable microstructure.

\subsection*{4.5.2. Most influencing parameters of LPBF and WAAM}
LPBF and WAAM have various parameters that can cause variations in microstructural, mechanical properties, chemical composition and dimensional accuracy of build components. Consideration of all the process parameters will be difficult; therefore, it is significant to recognize and concentrate on the most influencing parameters such as laser power, scanning speed, hatch distance, laser beam diameter and scan strategy in LPBF- AM and current, wire feed speed, travel speed, and inert gas flow rate in WAAM-AM process. One of the most significant ways to determine the most influencing parameters for the deposition of $\mathrm{Mg}$ alloy is through the design of the experiment (DOF) [104]. For instance, Wei et al. [121] studied the effect of energy input on formability, microstructure and mechanical properties of LPBFed AZ91D Mg alloy. They reported that high laser power drastically reduced porosity. It also illustrates that reducing laser scanning speed at constant laser power produced porous components. Thus, laser power, scanning speed, and hatch distance should be considered while depositing Mg alloy. Similarly, in the WAAM process, Takagi et al. [67] deposited AZ31B Mg alloy using the GMAW-WAAM process and studied the effect of welding current and travel speed on deposited track quality. They reported smooth deposited tracks could not be achieved under welding current $60 \mathrm{~A}$. This is because during deposition using a semiautomatic welding machine such as GMAW, the welding current determines the required length of feed wire. When the welding current decreased, the wire feed rate and the amount of wire supplied per unit length decreased. Therefore, under the current $60 \mathrm{~A}$, the amount of wire supplied was insufficient; hence smooth deposited tracks could not be achieved. Moreover, at a travel speed of $1000 \mathrm{~mm} / \mathrm{min}$ humping defect occurs due to the travel speed exceeding a certain critical point. Hence, current, wire feed speed, travel speed, and inert gas flow rate must be carefully considered during the WAAM process to achieve superior deposited components. The most significant parameters that can affect the quality of various grades of $\mathrm{Mg}$ alloy deposited by the WAAM and LPBF process are listed in Table 5.

\subsection*{4.6. Application}
Demand for Mg and Mg alloys is rapidly increasing due to huge applications in aircraft, automotive, armaments, electronics, sports, construction and marine [149,150]. However, Mg and Mg alloys are extensively employed in biomedical applications owing to their appropriate mechanical properties and biocompatible and biodegradable nature [150,151]. The manufacturing of $\mathrm{Mg}$ alloy is quite challenging. Therefore, additive manufacturing processes are used to manufacture biomaterials such as $\mathrm{Mg}$ and its alloy [151-155]. The additively manufactured $\mathrm{AZ}$ series of $\mathrm{Mg}$ alloys are highly significant in components, including biomedical implants such as stents, cardiovascular stents, hip and knee joint implants and scaffolds [158-164].

Moreover, it is remarkable that to support the stents and other implants devices while healing the bone, the Mg alloy can be utilized in orthopaedic fixation (i.g., bones pins, compressive plates and screws) because of the biodegradable and nontoxic nature of $\mathrm{Mg}$, as can be seen in Fig. 28 [164,165]. Even though Mg is considered for non-bearing applications such as stents due to its low mechanical properties, however, some Mg alloys such as WE43 and AZ80M have the potentials to use for load-bearing applications due to their comparable strength to bone and lower elastic modulus compared to commercial available biomedical implants such as Ti-6Al-4V and SS316 steel [163-167]. From the literature survey of Mg alloy deposited using LPBF and WAAM processes, it is worth noting that both approaches are suitable for depositing $\mathrm{Mg}$

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-24}
\end{center}

Fig. 27. (a) Phase identified using JMatPro (b) phase diagram of Mg-Al [146].

Table 5

Most influencing parameters for WAAM and LPBF deposition of Mg and its alloys.

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-25}
\end{center}

biomaterials and biomedical implants; however, compared to the LPBF process, the WAAM process cannot be manufactured very complex magnesium components. But regarding the size of parts, the WAAM process is more suitable than LPBF. Therefore, both methods can produce excellent $\mathrm{Mg}$ components specifically.

\section*{5. Summary and outlook}
In this review, the design aspect of different gas shielding, deposition, microstructure, mechanical properties, and corrosion behaviour of magnesium alloys fabricated using wire arc additive manufacturing (WAAM) and laser powder bed fusion (LPBF) have been critiqued and summarized. The collected results revealed that WAAM and LPBF are widely used to deposit Mg alloy in additive manufacturing processes. The additive manufacturing process represents a sustainable process for the deposition of $\mathrm{Mg}$ alloy because of high material utilization (especially in WAAM), relatively low production time and less manufacturing cost. In addition, some adverse effects such as oxidation, evaporation, tensile residual stress, distortion, porosity and crack formation are major challenges during the deposition of $\mathrm{Mg}$ alloy using AM process. Therefore, these defects create a grey area for further studies to thoroughly understand the external responsible factors such as man, method, and materials. In this regard, the right strategies like skilled operator (Man), suitable process (technique) and right materials should be selected to eliminate unfavourable conditions, thereby obtaining high-performance, high-quality and reliable products. CMT + pulse-based WAAM with a roller system controls these defects. In CMT + Pulse-based WAAM, low heat input and pulse frequency are responsible for reducing residual stress and applied roller just after deposition of each layer, thereby decreasing pore and other internal defects. Consequently, it reduced the anisotropic behaviour of mechanical properties of deposited components. Similarly, in LPBF, balling effect, residual stress, unmelted powder and porosity are significant concerns; therefore, it is essential to control these defects to achieve a better quality of $\mathrm{Mg}$ parts. Researchers use various techniques, such as optimizing process parameters before deposition and applying an additional set-up for heating and cooling the substrate to minimize residual stress. However, these techniques could not eliminate all the defects; hence, various post-processing techniques such as laser shock peening, shot peening, surface melting and ball-rollerdiamond burnishing were used to eliminate these defects and further enhancement of microstructural, mechanical and corrosion preperformance of deposited Mg alloys. Moreover, the WAAM process is economical and able to manufacture comparatively large\\
\includegraphics[max width=\textwidth, center]{2024_04_13_7e5fbe1213c12de51fd6g-26(2)}\\
\includegraphics[max width=\textwidth, center]{2024_04_13_7e5fbe1213c12de51fd6g-26}\\
\includegraphics[max width=\textwidth, center]{2024_04_13_7e5fbe1213c12de51fd6g-26(1)}

Fig. 28. Mg alloy scaffolds, manufactured using LPBF process (a) extracted from Ref. [166] (b) extracted from Ref. [167] (c) extracted from Ref. [117] (d)) extracted from Ref. [168] (e) extracted from Ref. [169] (f)) extracted from Ref. [131] (g) biomedical components of Mg alloy and their application [158,159].

components; however, dimensional accuracy and highly complex shape deposition are quite challenging. In contrast to WAAM, LPBF processes are highly accurate and can deposit any shape of components. In addition, the one major challenge during the deposition of $\mathrm{Mg}$ alloy is maintaining the inert gas environments owing to the volatile nature of Mg powder. Further, to advancements of WAAM and LPBF process, it must be integrated with computer technology, like decision science, machine learning and process modelling, to achieve high efficiency and real-time monitoring from different places. The machine learning and process modelling approaches can help to optimize the WAAMed and LPBF process and make a better understanding of the process at a low cost. The primary concerns in the AMed component are residual stress development which varies layer by layer. Therefore, it is challenging, timeconsuming, and more expensive to measure each layer experimentally. However, using numerical modelling, these challenges can be addressed easily, consequently saving time and money. Current drawbacks associated with WAAM and LPBF of Mg alloy manufacturing do not seem unbridgeable. However, intensive research is required to realize the potential of WAAMed and LPBFed $\mathrm{Mg}$ alloy in advanced applications such as satellite parts, aerospace and biomedical. Therefore, the future research scopes based on this review are (i) investigating the deposition quality of various additive manufactured Mg alloy numerically and experimentally to

\begin{center}
\includegraphics[max width=\textwidth]{2024_04_13_7e5fbe1213c12de51fd6g-27}
\end{center}

Fig. 29. Future advancement in the deposition of $\mathrm{Mg}$ alloy using laser and arc additive manufacturing process.

optimize and understand the process parameters, (ii) studying the developments of residual stress during deposition, (iii) to investigate the effect of post-processing such as surface melting, shot peening, burnishing and shock peening on mechanical properties and corrosion performance of as-deposited Mg alloy (iv) to investigate the fatigue corrosion behaviour of as-deposited $\mathrm{Mg}$ and post processed Mg alloy. Apart from these, Fig. 29 shows the objectives not being explored to date in the field of additive manufacturing of magnesium alloys.

\section*{Conflicts of interest}
The authors declare that there is no conflicts of interest.

\section*{Acknowledgment}
The Department of Science and Technology (DST), Government of India, Grant No SP/YO/2019/1287(G), financially supported this work.

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