% This LaTeX document needs to be compiled with XeLaTeX. \documentclass[10pt]{article} \usepackage[utf8]{inputenc} \usepackage{ucharclasses} \usepackage{amsmath} \usepackage{amsfonts} \usepackage{amssymb} \usepackage[version=4]{mhchem} \usepackage{stmaryrd} \usepackage{hyperref} \hypersetup{colorlinks=true, linkcolor=blue, filecolor=magenta, urlcolor=cyan,} \urlstyle{same} \usepackage{graphicx} \usepackage[export]{adjustbox} \graphicspath{ {./images/} } \usepackage{multirow} \usepackage{bbold} \usepackage{polyglossia} \usepackage{fontspec} \setmainlanguage{english} \setotherlanguages{hebrew} \newfontfamily\hebrewfont{Noto Serif Hebrew} \newfontfamily\lgcfont{CMU Serif} \setDefaultTransitions{\lgcfont}{} \setTransitionsFor{Hebrew}{\hebrewfont}{\lgcfont} \title{Powder Bed Fusion of nickel-based superalloys: A review } \author{Salomé Sanchez ${ }^{\mathrm{a}}$, Peter Smith ${ }^{\mathrm{a}}$, Zhengkai Xu ${ }^{\mathrm{a}}$, Gabriele Gaspard ${ }^{\mathrm{a}, \mathrm{c}}$, Christopher J. Hyde ${ }^{\mathrm{a}}$,\\ Wessel W. Wits ${ }^{\mathrm{b}}$, Ian A. Ashcroft ${ }^{\mathrm{a}}$, Hao Chen ${ }^{\mathrm{d}}$, Adam T. Clare ${ }^{\mathrm{a}, *}$\\ a Faculty of Engineering, University of Nottingham, Advanced Manufacturing Building, Jubilee Campus, Nottingham, NG7 2RD, United Kingdom\\ ${ }^{\mathrm{b}}$ University of Twente, Drienerlolaan 5, 7522 NB, Enschede, Netherlands\\ ${ }^{\mathrm{c}}$ Department of Mechanical Engineering, EPFL, Route Cantonale, 1015, Lausanne, Switzerland\\ d Department of Mechanical, Materials and Manufacturing Engineering, University of Nottingham Ningbo China, Ningbo 315100, China} \date{} %New command to display footnote whose markers will always be hidden \let\svthefootnote\thefootnote \newcommand\blfootnotetext[1]{% \let\thefootnote\relax\footnote{#1}% \addtocounter{footnote}{-1}% \let\thefootnote\svthefootnote% } %Overriding the \footnotetext command to hide the marker if its value is `0` \let\svfootnotetext\footnotetext \renewcommand\footnotetext[2][?]{% \if\relax#1\relax% \ifnum\value{footnote}=0\blfootnotetext{#2}\else\svfootnotetext{#2}\fi% \else% \if?#1\ifnum\value{footnote}=0\blfootnotetext{#2}\else\svfootnotetext{#2}\fi% \else\svfootnotetext[#1]{#2}\fi% \fi } \begin{document} \maketitle \section*{A R T I C L E I N F O} \section*{Keywords:} Additive manufacturing Powder bed fusion Laser powder bed fusion Electron beam melting Nickel-based superalloys Mechanical properties Microstructure evaluation Tensile Hardness Shear Fatigue Creep Toughness \begin{abstract} A B S T R A C T Powder Bed Fusion (PBF) techniques constitute a family of Additive Manufacturing (AM) processes, which are characterised by high design flexibility and no tooling requirement. This makes PBF techniques attractive to many modern manufacturing sectors (e.g. aerospace, defence, energy and automotive) where some materials, such as Nickel-based superalloys, cannot be easily processed using conventional subtractive techniques. Nickelbased superalloys are crucial materials in modern engineering and underpin the performance of many advanced mechanical systems. Their physical properties (high mechanical integrity at high temperature) make them difficult to process via traditional techniques. Consequently, manufacture of nickel-based superalloys using PBF platforms has attracted significant attention. To permit a wider application, a deep understanding of their mechanical behaviour and relation to process needs to be achieved. The motivation for this paper is to provide a comprehensive review of the mechanical properties of PBF nickel-based superalloys and how process parameters affect these, and to aid practitioners in identifying the shortcomings and the opportunities in this field. Therefore, this paper aims to review research contributions regarding the microstructure and mechanical properties of nickel-based superalloys, manufactured using the two principle PBF techniques: Laser Powder Bed Fusion (LPBF) and Electron Beam Melting (EBM). The 'target' microstructures are introduced alongside the characteristics of those produced by PBF process, followed by an overview of the most used building processes, as well as build quality inspection techniques. A comprehensive evaluation of the mechanical properties, including tensile strength, hardness, shear strength, fatigue resistance, creep resistance and fracture toughness of PBF nickel-based superalloys are analysed. This work concludes with summary tables for data published on these properties serving as a quick reference to scholars. Characteristic process factors influencing functional performance are also discussed and compared throughout for the purpose of identifying research opportunities and directing the research community toward the end goal of achieving part integrity that extends beyond static components only. \end{abstract} Build direction and axis definition are almost uniformly defined as below (as shown in Fig. 1) \section*{1. Introduction} Additive manufacturing (AM) can be defined as "a process of joining materials to make objects from 3D model data, usually layer upon layer, as opposed to subtractive manufacturing methodologies". This technique has drawn significant attention due to its flexibility in design and fabrication. Commercially, AM has the potential to save both money and time while delivering enhanced functionality with respect to conventional subtractive manufacturing techniques. This becomes apparent when highly customised parts with high value and low volume are required. Several researchers have previously reviewed the potential advantages of AM, as well as its positive impact on society [1]. These studies concluded that this process is driving a revolution to manufacturing technology. In the last two decades there has been a dramatic increase in the number of publications associated with Nickel-based materials in AM (see Fig. 2). \footnotetext{\begin{itemize} \item Corresponding author. \end{itemize} E-mail addresses: \href{mailto:salome.sanchez3@nottingham.ac.uk}{salome.sanchez3@nottingham.ac.uk} (S. Sanchez), \href{mailto:peter.smith2@nottingham.ac.uk}{peter.smith2@nottingham.ac.uk} (P. Smith), \href{mailto:xu_zhengkai@simtech.a-star.edu.sg}{xu\_zhengkai@simtech.a-star.edu.sg} (Z. Xu), \href{mailto:gabriele.gaspard@alumni.epfl.ch}{gabriele.gaspard@alumni.epfl.ch} (G. Gaspard), \href{mailto:christopher.hyde@nottingham.ac.uk}{christopher.hyde@nottingham.ac.uk} (C.J. Hyde), \href{mailto:wessel.wits@nl.thalesgroup.com}{wessel.wits@nl.thalesgroup.com} (W.W. Wits), ian.ashcroft@ \href{http://nottingham.ac.uk}{nottingham.ac.uk} (I.A. Ashcroft), \href{mailto:Hao.Chen@nottingham.edu.cn}{Hao.Chen@nottingham.edu.cn} (H. Chen), \href{mailto:adam.clare@nottingham.ac.uk}{adam.clare@nottingham.ac.uk} (A.T. Clare). } \begin{center} \begin{tabular}{|llll|} \hline \multicolumn{2}{|l|}{Nomenclature} & HT & Heat Treatment \\ \multicolumn{1}{|l|}{AB} & As Built & IN & Inconel \\ BD & Build Direction (will be aligned with the Z-axis in Figures, & LPBF & Laser Powder Bed Fusion \\ & unless otherwise specified, see Fig. 1) & OM & Optical Micrography \\ CAD & Computer Aided Design & PBF & Powder Bed Fusion \\ DA & Direct Ageing & PREP & Plasma Rotated Electrode Process \\ EBM & Electron Beam Melting & RA & Recrystallisation Annealing \\ EBSD & Electron Back Scatter Diffraction & RT & Room temperature \\ EDM & Electrical Discharge Machining & SEM & Scanning Electron Microscopy \\ FCC & Face Centred Cubic & SR & Stress Relieved \\ HA & Homogenisation and ageing & ST & Solution treatment \\ HCF & High Cycle Fatigue & STA & Solution treatment and Ageing \\ HIP & Hot Isostatic Pressing & TEM & Transmission Electron Microscope \\ HSA & Homogenisation and solution treatment and ageing & & \\ \hline \end{tabular} \end{center} \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-02(1)} \end{center} Fig. 1. Layout of the build directions and principle planes commonly used when highlighting anisotropy in PBF specimens. Noting the orientation of this primitive is commonly considered to align with the machine coordinate system. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-02} \end{center} Fig. 2. Number of publications on Additive Manufacturing of nickel-based superalloys (from Scopus database) with major milestones [2-8].This shows an increase in publications associated with Nickel-based materials in Additive Manufacturing. There are several prior reviews regarding AM, particularly Powder Bed Fusion (PBF), of metals. The reviews in the literature to date provide a general overview of metal AM techniques, focusing on the processing methods, corresponding microstructures, mechanical properties and their applications with a variety of materials [3,9-12]. Murr et al. compared the two main metal PBF techniques, i.e. Laser Powder Bed Fusion (LPBF) and Electron Beam Melting (EBM), and their processing of commonly used structural materials [13], while Korner et al. reviewed EBM and the process-property links in metallic materials [14]. Song et al. reviewed the differences in microstructures and mechanical properties between LPBF manufactured components and their conventionally manufactured counterparts [15]. Sames et al. reviewed a series of metal AM techniques, focusing on their issues and their mechanisms, as well as creating models to simulate them [16]. Basak and Das reviewed the microstructure evolution in commonly used metal alloys processed with various metal AM techniques [17]. In the literature, review papers on the general mechanical properties of LPBF manufactured components can be found [18-20]. These tend to focus on the microstructure and mechanical performance of additively manufactured Ti-6Al-4V components [21,22] and nickel-based superalloys [23,24]. Wang et al. reviewed LPBF manufactured Inconel (IN) 718 components, mainly concentrating on its microstructure and mechanical properties [25]. This is perhaps the most similar work to the review undertaken here which only represents a fraction of the state-of-the-art now available in the literature. Aboulkhair et al. comprehensively reviewed aluminium alloys [26]. Zhang et al. reviewed the applications of LPBF titanium alloys and of titanium matrix composites in biomedical engineering [27]. It is worth mentioning that there are also review papers on novel design [28] and material systems [29] applied to additive manufacturing. However, there is yet to be presented a comprehensive review on the subject of mechanical properties of PBF Nickel-based superalloys which provides an authoritative and comprehensive resource to scholars in this field. Therefore, the motivation for this paper is to present a comprehensive review of the mechanical properties of PBF nickel-based superalloys. This will provide researchers with a better understanding of the state-of-the-art and the effect of PBF processing parameters on the mechanical properties. A summary of the research undertaken for different mechanical properties will be given at the end of this review to help practitioners to identify what categories of material evaluation have been performed and to identify gaps in research. Finally, both the opportunities and shortfalls of PBF in the processing of Nickel-based superalloys will be discussed. Since our understanding of the fundamental metallurgy and process itself develops at a rapid rate it essential that the community has a point of reference from which to draw. \subsection*{1.1. Powder bed fusion techniques} This paper focuses on the two principle PBF techniques: LPBF and EBM. PBF is one of the most popular AM techniques for metal part fabrication. This process consists of two stages: firstly, the powder is spread uniformly on the working area, then an energy source (a laser\\ beam for LPBF and an electron beam for EBM) selectively melts the powder bed according to a 3D model and hence build the final component [30]. The two main differences between these fabrication methods are their power sources and power transmission systems. Two diagrams representing a typical LPBF and EBM systems are presented in Fig. 3 [9]. Both methods and process nuances have already been properly reviewed by other researchers and Table 1 summarises the relevant differences between the two systems $[9,16,30]$. This method can be used to process a variety of materials, ranging from metals to ceramics, for many applications, such as aerospace, biomedical and automotive. For a broad review of the materials available for all types of AM platform the reader is referred to Bourell et al.'s review of the topic [31]. \subsection*{1.2. Nickel-based superalloys and the role for $A M$} This section will highlight the characteristics of nickel-based superalloys, why they are appropriate for PBF use and examples of their applications, particularly in the aerospace sector. \subsection*{1.2.1. Characteristics of nickel-based superalloys} Nickel-based superalloys, as a family of modern aerospace engine materials [34,35], which possess a combination of high-temperature strength, toughness, creep and oxidation/corrosion resistance. For these reasons, this class of alloys has been widely used in components operating in critical environments [36]. The first generation of nickel-based superalloys, designed for high-temperature applications in jet engines, included Nimonic 75, developed by Henry Wiggin Ltd, UK, in the 1940s [34]. Since then, nickel-based superalloys have been continuously produced, studied and used in building turbine blades, turbine discs, seals, rings, and other components in gas turbines. Nowadays, there are nearly 1.8 tonnes of nickel-based superalloys in a typical jet engine. These materials have greatly contributed to the increase of the continuous operating life of jet engines to above $20,000 \mathrm{~h}$ [37]. While coating technologies (e.g. Zirzonia based thermal barrier coatings, TBCs) have also served to enhance high temperature performance the role of the substrate nickel-based superalloy cannot be overstated. Fig. 4 shows the weldability and therefore the utility of nickel-based superalloys in fabrications. This is a useful indicator of how challenging high integrity AM will be for a given material. Effectively, the process window becomes greatly reduced above the broken red line. Nickel-based superalloys are used in many applications, such as landbased gas turbines, nuclear power plants and chemical containers. A summary of some common applications for nickel-based superalloys are reported in Table 2.\\ Table 1 The major differences between Laser Powder Bed Fusion and Electron Beam Melting. This provides an insight into process characteristics [9,16,30]. \begin{center} \begin{tabular}{|c|c|c|} \hline Aspects & LPBF & EBM \\ \hline Power source & Laser & Electron beam \\ \hline Power range & $20 \mathrm{~W}-1 \mathrm{KW}$ & \begin{tabular}{l} Several KW, much higher \\ than the laser power \\ \end{tabular} \\ \hline \begin{tabular}{l} Energy beam spot \\ size \\ \end{tabular} & $50 \mu \mathrm{m}-180 \mu \mathrm{m}$ & $50 \mu \mathrm{m}-200 \mu \mathrm{m}$ \\ \hline \begin{tabular}{l} Power \\ transmission \\ system \\ \end{tabular} & \begin{tabular}{l} High-frequency scanning \\ mirrors \\ \end{tabular} & \begin{tabular}{l} Electromagnetic lenses and \\ magnetic scan coil \\ \end{tabular} \\ \hline Scan speed range & Up to $15 \mathrm{~m} / \mathrm{s}$ & Up to $10 \mathrm{~m} / \mathrm{s}$ \\ \hline \begin{tabular}{l} Powder bed \\ thickness range \\ \end{tabular} & $20 \mu \mathrm{m}-100 \mu \mathrm{m}$ & $50 \mu \mathrm{m}-200 \mu \mathrm{m}$ \\ \hline \begin{tabular}{l} Powder bed \\ temperature \\ range \\ \end{tabular} & \begin{tabular}{l} Significant substrate heating \\ is not usual and can range \\ from $20^{\circ} \mathrm{C}$ [32] to $975^{\circ} \mathrm{C}$ \\ [33] \end{tabular} & \begin{tabular}{l} Can be very high, slightly \\ below the materials' melting \\ temperature \\ \end{tabular} \\ \hline \begin{tabular}{l} Build chamber \\ condition \\ \end{tabular} & \begin{tabular}{l} Normally filled with \\ protective gases, with an \\ oxygen content less than $0.1 \%$ \\ \end{tabular} & Vacuum, $<10^{2} \mathrm{~Pa}$ \\ \hline \end{tabular} \end{center} Fig. 5 summarises the types of nickel-based superalloys studied in PBF research. It is clear from this figure that IN718 and IN625 are the most explored in PBF studies. Their composition can be found in Table 8 in the Appendix. \subsection*{1.2.2. Nickel-based superalloys and PBF in the aerospace Industry} The aerospace industry represents a significant prize for AM machine producers, since a significant amount of components have a high part value and are produced using high value materials [40]. Indeed, nickel-based aerospace components are characterised by complex geometries and low production volumes. Furthermore, given the characteristic excellent mechanical properties of Nickel-based superalloys, designed to work in safety critical applications, these parts are difficult and expensive to machine with conventional machining techniques [41-44]. As a result, particular attention must be paid to the selection of tooling, coolants, and processing parameters, leading to increasing production costs [40]. On the contrary, PBF's ability to manufacture complex geometries allow the incorporation of new and additional functionalities to components. Hence, this area is an appropriate way to demonstrate the potential of using PBF in conjunction with nickel-based superalloys. Work by Yadroitsev et al. demonstrated the capability of LPBF in producing complex filters constituted of free-form structures from IN625 [45]. The parameters of the unit cell, or even of individual cells, can be \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-03} \end{center} Fig. 3. Diagram of the two main Powder Bed Fusion processes. (a) Laser Powder Bed Fusion. (b) Electron Beam Melting systems. The two main differences (power sources and power transmission systems) are visible. After [9]. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-04} \end{center} Fig. 4. 'Weldability' diagram for a range of nickel-based superalloys as a function of their Ti and Al alloy element composition. Figure from CatchpoleSmith et al. [38]. Weldability is considered poor above the dashed line and deteriorates with increasing Ti and $\mathrm{Al}$ content. This work will inform future research on alloy compositions. easily modified at the modelling stage and optimised for specific applications with features below $1 \mathrm{~mm}$ (Fig. 6). This demonstrates an immediate opportunity for AM technologies which cannot be achieved through conventional machining methods. A further advance was demonstrated by Bernstein et al., who developed and built a prototype turbine blade characterised by a leading edge with inbuilt cooling channels obtained via LPBF (Fig. 7) [42]. The measured average cooling hole size $(0.3945 \mathrm{~mm}$ ) was slightly smaller than the nominal size $(0.5 \mathrm{~mm})$, highlighting tolerancing challenges remaining to be conquered in LPBF. The standard deviation for these hole diameters was small $(0.0206 \mathrm{~mm})$, indicating LPBF has potential as an accurate and effective technique to produce these features. Despite some promising results, this process cannot yet compete with the feature resolution obtained by laser processes or electrical discharge machining (EDM), which are common methods for cooling hole introduction. Indeed, Li et al., amongst others, have demonstrated that by using a laser or EDM drilling process it is possible to obtain holes with a diameter $<150 \mu \mathrm{m}$. However, common industrial processes for civil aircraft require cooling hole diameters in the range 300-500 $\mu \mathrm{m}$ [46]. This is a fundamental technology limitation which governs turbine blade and nozzle guide vain thermodynamic performance. However, focus on AM research for aerospace applications has so far been focused on static components. The reasons behind this becomes evident with exploration of the mechanical properties exhibited by AM materials, a topic explored in much more detail later in this review. Since LPBF has shown promising initial results in the realising complex structures in aero-engine components, more and more companies are expressing intentions to invest in this technology. A review of metal AM use in the commercial aviation industry was written by Gisario et al. [47] and a review of the material characteristics of AM IN718 for high temperature applications was written by Yong et al. [48].\\ Table 2 Common applications of nickel-based superalloys [39]. This shows the numerous areas where additively manufactured nickel-based superalloys could be used. \begin{center} \begin{tabular}{|c|c|c|} \hline Applications & Examples & Typical alloys \\ \hline Aerospace Industry & \begin{tabular}{l} Aircraft gas turbines: disks, \\ combustion chambers, bolts, \\ casings, shafts, exhaust \\ systems, blades, vanes, burner \\ cans, afterburners, thrust \\ reversers, \\ Space vehicles: \\ aerodynamically heated skins, \\ rocket engine parts \\ \end{tabular} & \begin{tabular}{l} IN600, IN601, IN617, \\ IN625, IN706, IN718, \\ IN738, IN754, IN X-750, \\ Nimonic 115, Nimonic \\ 75, Nimonic 80, Nimonic \\ 90, Rene 41, Waspaloy, \\ Hastelloy X \\ \end{tabular} \\ \hline \begin{tabular}{l} Chemical and \\ petrochemical \\ industries \\ \end{tabular} & \begin{tabular}{l} bolts, fans, valves, reaction \\ vessels, tubing, transfer piping, \\ pumps \\ \end{tabular} & \begin{tabular}{l} IN600, IN625, IN690, \\ IN718, IN725, IN925, \\ Rene 41, Waspaloy \\ \end{tabular} \\ \hline Pulp and paper mills & \begin{tabular}{l} tubing, doctor blades, \\ bleaching circuit equipment, \\ scrubbers \\ \end{tabular} & \begin{tabular}{l} Hastelloy G, IN600, \\ IN671, IN706, IN718, \\ Rene 41, Waspaloy \\ \end{tabular} \\ \hline \begin{tabular}{l} Nuclear power \\ systems \\ \end{tabular} & \begin{tabular}{l} control rod drive mechanisms, \\ valve stems, springs, ducting \\ \end{tabular} & \begin{tabular}{l} Hastelloy G, IN600, \\ IN625, IN706, IN718, \\ Rene 41, Waspaloy \\ \end{tabular} \\ \hline Marine architecture & ships, submarines & \begin{tabular}{l} IN600, IN625, IN718, \\ Rene 41, Waspaloy \\ \end{tabular} \\ \hline Electronic Parts & resistors & \begin{tabular}{l} IN706, IN718, \\ Nichrome, Waspaloy \\ \end{tabular} \\ \hline \begin{tabular}{l} Steam turbine power \\ plants \\ \end{tabular} & \begin{tabular}{l} bolts, blades, stack gas \\ reheaters \\ \end{tabular} & IN706, IN X-750 \\ \hline \begin{tabular}{l} Metals processing \\ mills \\ \end{tabular} & \begin{tabular}{l} ovens, furnace, afterburners, \\ exhaust fans \\ \end{tabular} & \begin{tabular}{l} IN600, IN625, IN706, \\ IN718, N06008, \\ Nichrome, Rene 41, \\ Waspaloy \\ \end{tabular} \\ \hline \begin{tabular}{l} Heat-treating \\ equipment and \\ Metal processing \\ \end{tabular} & \begin{tabular}{l} trays, fixtures, conveyor belts, \\ baskets, fans, furnace mufflers, \\ hot-work tools and dies \\ \end{tabular} & \begin{tabular}{l} IN600, IN706, Nimonic \\ 80, Rene 41, Waspaloy, \\ Waspaloy \\ \end{tabular} \\ \hline Automotive industry & \begin{tabular}{l} spark plugs, glow plugs (in \\ diesel engines), catalytic \\ converters, combustion \\ systems \\ Reciprocating engines: \\ turbochargers, exhaust valves, \\ hot plugs, valve seat inserts \\ \end{tabular} & IN625, Waspaloy \\ \hline Medical applications & \begin{tabular}{l} dentistry uses, prosthetic \\ devices \\ \end{tabular} & \begin{tabular}{l} Vitallium, Ni-Cr and \\ Ni-Ti alloys \\ \end{tabular} \\ \hline \begin{tabular}{l} Pollution control \\ equipment \\ \end{tabular} & \begin{tabular}{l} scrubbers, flue gas \\ desulfurization equipment \\ (liners, fans, stack gas \\ reheaters, ducting) \\ \end{tabular} & IN718 \\ \hline \begin{tabular}{l} Coal gasification and \\ liquefaction \\ systems \\ \end{tabular} & \begin{tabular}{l} heat exchangers, repeaters, \\ piping \\ \end{tabular} & IN690 \\ \hline \end{tabular} \end{center} As an example, NASA has tested some LPBF built rocket injectors, demonstrating that these parts can withstand heat and pressures generated during space rocket launches [49]. MTU Aero Engines also announced that the borescope bosses for their PurePower PW1100G-JM engines will now be produced using LPBF [50]. Further, the Netherlands Aerospace Centre together with the University of Twente developed a novel micro-pump assembly for space application, composed of no moving parts, such as hydraulic valves [51]. The manufacturing of this micropump was only possible using LPBF, due to the complex internal features which could not be obtained by other means. GE Aviation played a fundamental role in the introduction of AM to the aerospace industry, in particular through the acquisition of both SLM Solutions and Arcam, two major AM companies specialising in LPBF and EBM, respectively. As a proof of principle, GE Aviation built a working miniature version of a jet engine using entirely LPBF [52]. Even though the scale was far smaller than commercial engines, this prototype was able to reach 33,000 RPM in functional testing and marks a significant step towards a more widespread use of LPBF in aero-engines manufacturing. They also created a sensor housing using only LPBF, which made it the first 3D printed part to be approved for use by the FAA \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-05(1)} \end{center} Fig. 5. Pie chart showing the nickel-based superalloys studied in powder bed fusion research to date, from 290 studies. Inconel 718 and Inconel 625 are the most studied alloys given their level of usage primarily in the aerospace markets where there are immediate opportunities for aerospace.\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-05} Fig. 6. Inconel 625 filter with placement specific pore orientation and cross-sectional area. These samples were manufacture by powder bed fusion, demonstrating the ability of powder bed fusion to manufacture highly optimised geometries with features $<1 \mathrm{~mm}$. This is an exemplary use of Laser Powder Bed Fusion technology [45]. [53]. Other demonstrators, like turbine blades, were printed using EBM IN738LC and were assembled onto a disk for spin pit testing to validate the mechanical integrity and design of the blades [54]. In summary, AM is a promising technique for the manufacturing of nickel-based components. However, the knowledge gaps, which currently restrain PBF from wider scale exploitation, remain significant. Failure to fully understand the microstructural and functional response of AM nickel-based superalloys will severely limit the applications for this technology/material combination. As such, we must obtain a fundamental understanding of the origin of defects in both material and process. \section*{2. PBF process control and quality inspection} In order to efficiently transfer the PBF processes to industry, adequate and solid inspection methods for both the building process and build quality must be selected. Some of the techniques proposed for this purpose are discussed in this section and will provide the scholar with reference methodologies. \subsection*{2.1. Microstructural characterisation} The most common methods for the analysis of PBF nickel-superalloys include Optical Micrography (OM), Scanning Electron Microscopy (SEM) and Electron Back Scatter Diffraction (EBSD), which are described briefly below. All of these are common to well equipped metallography laboratories. In order to use those methods, samples need to be carefully prepared. A review on the preparation of metallic materials has also been written by Zhang et al. [55] which will assist nickel-based superalloy and AM researchers in developing characterisation strategies. Here, the main observations under these techniques are highlighted. Arriving at appropriate specimen conditions often requires appropriate grinding, polishing, and etching. The reader can refer to Zhang et al.'s review [55] on this topic for more information. Porosity in PBF specimens is a classic 'first indicator' to investigate and although the Archimedes principle is suggested by ASTM standards for carrying out porosity measurements for PBF materials [56], OM is typically used by the research community to observed these at a x50 magnification or less (Fig. 8). The lateral resolution of OM is in the order of $200 \mathrm{~nm}$ [57]. If a higher resolution is required (smaller pore size) use of electron microscopy is required. Specimen cross-sections are often analysed using an image analysis software, such as ImageJ, and porosities can be quantified albeit destructively. Perevoshchikova et al. proved that these porosity values were comparable to those obtained with the Archimedes method [58]. In softer materials caution must be taken in order not to obscure pores by material smearing upon polishing. Melt pools can also be identified using OM (Fig. 29) under an appropriate etch. However, microstructure at the grain scale is usually not observed using OM since salient features in PBF nickel-based superalloy specimens are typically 5-30 $\mu \mathrm{m}$. As such, higher resolution imaging techniques, such as SEM, are required if understanding beyond the macro (weld tracks, pores) is required. SEM is widely used to characterise PBF nickel-based superalloys sample microstructures. Along with back scattered electron micrographs, the material's surface topography, grain structure, phases and precipitates can be observed. In fact, the fast heating and cooling cycles produced during PBF, often make precipitates small (in the range of $\mathrm{nm}$ ) which may be beyond the limit of SEM. Since the composition, spatial frequency and size of these are critical in determining alloy performance. Characterisation of these (and controlling their formation) is critical in process. EBSD can provide more detailed information regarding the material's crystallographic texture as an accompaniment to electron microscopy. An EBSD orientation map of an As-Built (AB) LPBF IN718 specimen is given in Fig. 19 [59]. EBSD is exceptionally useful in relating the textural formation in AM with the associated process parameters and material composition. Data sets emerging from EBSD are highly valuable when considering the recrystallisation behaviour of AM specimens. Terner et al. also used EBSD to estimate residual stress in LPBF IN625 by assessing misorientation or strain levels from local misorientation by means of orientation imaging, and found that EBSD was adequate to qualitatively assess residual stress in a material [60]. Allied to EBSD are a number of emergent techniques which make use of laser ultrasonics. Rossin et al. used resonant ultrasound spectroscopy to characterise and detect LPBF part microstructure variability [61]. Further Smith et al. demonstrated the use of Spatially Resolved Acoustic Spectroscopy for characterisation of AM components [62]. While these techniques are in their infancy they have clear potential to be used alongside AM in the production environment. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-06(2)} \end{center} Fig. 7. Example of novel cooling channels in a leading edge. (a) the overview. (b) Computer Aided Design model of the internal structure of the cooling channels [42]. The average measured cooling hole size $(0.39 \mathrm{~mm})$ was slightly smaller than the nominal size $(0.5 \mathrm{~mm})$, highlighting tolerancing challenges still to be conquered in Laser Powder Bed Fusion. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-06(3)} \end{center} Fig. 8. Optical micrographs of a laser powder bed fused specimen [58]. Porosity and other defects are clearly present in the material. These can be easily detected using optical micrography. However, caution should be taken when considering softer materials as smearing can serve to obscure pores. (Process parameters used: $295 \mathrm{~W}, 2250 \mathrm{~mm} / \mathrm{s} 0.11 \mathrm{~mm}$ layer height.) \subsection*{2.2. X-ray diffraction} X-ray diffraction (XRD) can be used not only to determine the crystalline structure of polycrystalline materials but also to measure residual stresses [63]. In Fig. 9, the XRD spectra for an IN718 powder and an AB LPBF specimen are reported, giving a general overview of the phase distribution in the material. As expected, the main phase present was the $\gamma$ face-centred cubic (FCC) NiCr phase [64]. From the peak analysis it is observed that the $\gamma^{\prime}$ and $\gamma$ " peaks can overlap with the $\gamma$, becoming difficult to separate the different contributions. Therefore, other techniques, such as Transmission Electron Microscope (TEM), became necessary to identify and quantify these two precipitates. Some examples of XRD use in literature include Xia et al. who studied the impact of additional strengthening particles (tungsten carbides WC) on LPBF IN718 specimens and used XRD to characterise the phase distribution in some LPBF IN718+WC composites [65]. $\gamma$-Ni, $\mathrm{Ni}_{2} \mathrm{~W}_{4} \mathrm{C}, \mathrm{NbC}$ \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-06(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-06(4)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-06} \end{center} (c) Fig. 9. X-ray Diffraction results for a laser powder bed fused Inconel 718 specimen. (a) Powder. (b) As-built vertical section. (c) As-built horizontal section [64]. As expected, the main phase present is the $\gamma$ face-centred cubic NiCr phase and it is observed that the $\gamma^{\prime}$ and $\gamma^{\prime \prime}$ peaks can overlap with the $\gamma$, becoming difficult to separate the different contributions.\\ and the residual WC particles were found as the main phases. A shift in the diffraction peaks was highlighted, which was probably due to the lattice strain generated by the presence of larger $\mathrm{W}$ atoms. These induced stresses were thought to be responsible for the observed material strengthening. In the study by Raghavan et al., spectra from specimens treated with different HT were compared [66]. The XRD spectra obtained from the specimens (Fig. 10) showed that increasing the solution temperature decreased the number of secondary phases $\left(\gamma^{\prime}\right.$, $\gamma$ ' and carbides). Popovich et al. used XRD to investigate the effect of post-processing techniques on LPBF IN718 specimens [67]. A large amount of $\delta$ and $\gamma$ " phases were generated during the heat treatment (HT), whereas two types of carbides ( $\mathrm{NbC}$ and $\mathrm{TiC}$ ) were produced after Hot Isostatic Pressing (HIP) and HT (Fig. 11). Consistent with prior studies, it also highlighted the limitations and inadequacy of the XRD spectra for quantifying and differentiating phases such as $\delta$ and $\gamma$ ', characterised by overlapping peaks. XRD can be used not only qualitatively, but also quantitatively in determining specimens phase distribution. As an example, in a further study by Popovich et al., XRD was used to analyse the chemical composition of LPBF IN718 specimens at different stages of manufacturing (reported in Table 3) [68]. It was found that the total amount of strengthening particles in the AB specimen slightly increased compared to the original powder, increasing the strength more than predicted. During homogenisation, $\delta$ particles were fully dissolved, partially lowering specimen strength. However, if this was followed by an ageing treatment, the total volume of strengthening phases reached the value of $\sim 33 \mathrm{vol} \%$, namely three times more than $\mathrm{AB}$ specimens. The strengthening effect of these precipitates was confirmed by tensile\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-07(2)} (b) Fig. 10. Comparison between the X-ray Diffraction spectra of two differently heat-treated Laser Powder Bed Fused Inconel 718. (a) Between a $2 \theta$ angle of $20^{\circ}-60^{\circ}$. (b) Between a $2 \theta$ angle of $60^{\circ}-120^{\circ}$ [66]. This showed that increasing the solution temperature decreased the number of secondary phases $\left(\gamma^{\prime}, \gamma^{\prime \prime}\right.$ and carbides). \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-07} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-07(1)} \end{center} (b) Fig. 11. X-ray Diffraction spectra of as-built and post-processed Inconel 718 specimens. (a) Full spectra. (b) An enlarged image [67]. This demonstrates the inadequacy of the X-ray Diffraction spectra for quantifying and differentiating phases such as $\delta$ and $\gamma$ ', characterised by overlapping peaks. Table 3 X-ray Diffraction results for laser powder bed fused Inconel 718 quantitative composition [68]. Illustrating that X-ray Diffraction can be used not only qualitatively, but also quantitatively in determining specimens phase distribution. \begin{center} \begin{tabular}{lll} \hline Sample & Phases & Composition/vol\% \\ \hline Original powder & $\gamma-\mathrm{Ni}^{\prime}$ & 90.0 \\ & $\gamma^{\prime}-\mathrm{Ni}_{3} \mathrm{Al}$ & $3.5-3.9$ \\ & $\gamma^{\prime \prime}-\mathrm{Ni}_{3} \mathrm{Nb}$ & $4.3-4.5$ \\ & $\delta-\mathrm{Ni}_{3} \mathrm{Nb}$ & $1.8-2.0$ \\ AB condition & $\gamma-\mathrm{Ni}^{\prime}$ & 86.8 \\ & $\gamma^{\prime}-\mathrm{Ni}_{3} \mathrm{Al}$ & 1.9 \\ & $\gamma^{\prime \prime}-\mathrm{Ni}_{3} \mathrm{Nb}$ & 8.0 \\ & $\delta-\mathrm{Ni}_{3} \mathrm{Nb}$ & 3.3 \\ Homogenisation & $\gamma-\mathrm{Ni}^{\prime}$ & 90.1 \\ & $\gamma^{\prime}-\mathrm{Ni}_{3}(\mathrm{Al}, \mathrm{Ti})$ & 1.9 \\ & $\gamma^{\prime \prime}-\mathrm{Ni}_{3} \mathrm{Nb}$ & 8.0 \\ & $\gamma-\mathrm{Ni}^{*}$ & 67.3 \\ & $\gamma^{\prime}-\mathrm{Ni}_{3}(\mathrm{Al}, \mathrm{Ti})$ & 8 \\ & $\gamma^{\prime}-\mathrm{Ni}_{3} \mathrm{Nb}$ & 4 \\ & $\delta-\mathrm{Ni}_{3} \mathrm{Nb}$ & 3.5 \\ & $\gamma^{\prime}-\mathrm{Ni}_{3} \mathrm{Al}$ & 17.2 \\ \hline \end{tabular} \end{center} testing: fully HT specimens showed a tensile strength of $1350 \mathrm{MPa}$, compared to $1002 \mathrm{MPa}$ for $\mathrm{AB}$ equivalents. XRD can also be used to measure the residual stresses generated during PBF. Residual stress is commonly observed in PBF as localised stresses are induced upon cooling. Goel et al. observed, using neutron diffraction, that residual stresses were higher in LPBF than in EBM AB IN718 samples [69]. Sanz et al. analysed the residual stresses in some LPBF IN718 specimens and explored how they were affected by different post-processing strategies [70]. These measurements indicated that the stresses, of tensile nature, drastically undermined performance. Shot peening (amongst other methods) can be used to induce high compressive stresses in the surface and hence counterbalance this effect (Fig. 12) [70]. However, much like additional HT intrusive post-processing steps serve to undermine the economic case for using PBF. \subsection*{2.3. X-ray computed tomography} X-ray Computed Tomography (XCT) uses X-rays to take multiple two-dimensional cross-sectional images of an object from different orientations, allowing observation within [71]. With the help of XCT, the built-in defects as well as those generated during the mechanical testing can be characterised without destroying the specimen. In a study by Tillmann et al., the porosity of LPBF IN718 was evaluated using OM and XCT [72]. The comparison between the two measurements showed a significance difference between the two measurements (Table 4). This difference was imputed by the limit of the XCT resolution. In fact, porosities with a diameter smaller than $8 \mu \mathrm{m}$ were not detected by this technique, measuring therefore a smaller amount of porosities. This is a fundamental limitation since defects at this size, and appropriate population, can undermine part integrity significantly. XCT analysis also revealed that a region with a high porosity density was located between skin and core. The formation mechanism of these defects was explained in an previous study [73]. Smith et al. also observed that most of the cavities generated during PBF were found near the surface under the conditions they explored (Fig. 13) [74]. The reconstructed 3D volumes clearly show that the build direction (BD) also has an impact on cavities distribution. However, this may be a result of unfavourable processing conditions as opposed to a general deficiency in processing nickel-based superalloys. $\mathrm{Xu}$ et al. performed a series of staged thermal-mechanical tests, investigating the defects evolution in some LPBF IN718 specimens during creep testing [75]. This was achieved by performing XCT at different stages of the test. Fig. 14 shows the porosity distribution along the specimen length at different stages: before testing, at 7.3\% strain, at $11.5 \%$ strain and after failure. The increase in porosity during creep is straightforward, and the weakest point (Peak 1' in Fig. 14) could also be easily identified through the use of XCT. Having introduced the primary techniques and common observations associated with PBF of Nickel-based superalloys is it now possible \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-08} \end{center} \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-08(1)} \end{center} Fig. 12. Residual stress conditions in specimens treated with different post-processing techniques (heat treatment and shot peening) [70]. This illustrates the use of X-ray Diffraction for residual stress measurements as well as the effect of shot peening for inducing compressive residual stresses.\\ Table 4 Specimens relative density using Optical Micrographs and X-ray Computed Tomography [72]. The comparison between the two measurements showed a significance difference between the two measurements, highlighting the limitation of X-ray Computed Tomography resolution. \begin{center} \begin{tabular}{lll} \hline Optical microscopy (average) & Transverse section & Longitudinal section \\ & $0.13 \pm 0.06 \%$ & $0.09 \pm 0.07 \%$ \\ \hline XCT (average) & Full volume $0.069 \pm 0.012$ vol $\%$ & \\ \hline \end{tabular} \end{center} to more closely inspect the microstructures which result from the process. \subsection*{2.4. PBF process monitoring} Monitoring the PBF building process is necessary to follow and control the process, understand how defects are developed and how they could be removed. This topic has been reviewed extensively by Everton et al. [76]. Robust and widely deployable in-process monitoring capability. Thermal imaging is one of the most used techniques for melt pool characterisation [77,78] and defect distribution [79]. Criales et al. and Arisoy et al. recorded the movement of a single laser scan during a LPBF IN625 building (Fig. 15) [77,80]. From these thermographs, melt pool sizes, particles spattering tendency, thermal gradients, heating and cooling rates were extracted. Spatter generated during laser scanning can create serious defects in the PBF samples surfaces and bulk [81]. Foster et al. demonstrated the validity of using thermal imagining for in-situ monitoring of spatter locations [82]. In the thermographs, spatter trajectories were identified from grey scaling and contour plots as shown in Fig. 16. Tan et al. also showed the potential of using neural-network based image segmentation for spatter extraction during LPBF [83]. Alternative monitoring methods also include high frame rate camera to monitor melt pools [84,85], reflectometer-based instrument to measure the dynamic laser energy absorption during the scan [86] and Back Scattered Electron detection system to record the in operando signal during EBM [87]. XCT has been commonly used as a technique for post build analysis (see section 2.3) but can also carry out online measurements. For example, Leung et al. presented the successful application of XCT in capturing pores generation and spatter distribution in single laser track scanning [88]. Finally, other reviewers categorized the defects generated in PBF specimens while monitoring the building process [89,90]. Overall, in-process monitoring has many benefits in terms of understanding the PBF process and controlling the quality of parts produced. This continues to be an important research area for PBF machine technology. \section*{3. Microstructural observations} AM is a layer-wise technique which differs in many aspects from conventional manufacturing techniques such as casting, forging, or rolling. Hence the microstructures observed here are distinct from rolled or wrought equivalents. A review of the microstructural differences between PBF and conventional materials was written by Song et al. [15]. This different approach generates some characteristic microstructures, leading to distinctive material properties. It is also possible, through various $\mathrm{PBF}$ techniques, to provide a spectrum of microstructures which may be more or less suited for a given application. This section will highlight typical PBF microstructures and show how process parameters give rise to these. A review on the use of LPBF $\gamma^{\prime}$-strengthened nickel-based superalloys was written by Adegoke et al. highlights the effect of process parameters on the microstructure and defects of these alloys [24].\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-09(2)} Fig. 13. X-ray Computed Tomography images of specimen cross-section in 2D and reconstructed 3D volumes. V, D and H indicated the different specimen positioning strategies during the building process (which is shown in the top right-hand corner of the image) [74]. The reconstructed 3D volumes clearly show that the build direction has an impact on cavities distribution.\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-09} Porosity $/ \%$ (a) $1^{\text {st }}$ stage (b) $2^{\text {nd }}$ stage\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-09(1)} (c) $3^{\text {rd }}$ stage (d) $4^{\text {th }}$ stage Fig. 14. Porosity distribution in the specimen at different stages of creep testing [75]. Showing the potential of using X-ray Computed Tomography to determine where failure will occur in a specimen and for investigating evolution during tests. \subsection*{3.1. Typical PBF microstructures} This section will describe typical LPBF and EBM microstructures and how these are linked to process parameters. PBF microstructures have specific process defects. Reviews on these defects and how process parameters affect them have been presented by Malekipour et al. and Grasso et al. [7,91]. Fig. 17 shows some of the typical defects in PBF processes. For more details, the reader is referred to the papers aforementioned. Fig. 18 presents an overview of the microstructures observed in $A B$ LPBF IN718 specimens [92]. The morphology of melt pools can be clearly observed in the XZ plane (Fig. 18a), while the laser scan tracks are recognisable in the XY plane (Fig. 18b) [92]. Fig. 19 shows an EBSD image with individual laser scan tracks with a width of $\sim 75 \mu \mathrm{m}$ on the XY plane [59]. AB samples have a strong $<100\rangle$ crystallographic texture in the build direction [93]. Small equiaxed grains with an average size of $10 \mu \mathrm{m}$ can be recognized at the track the overlapping regions between the tracks [59]. These microstructural differences in the two directions are responsible for the mechanical anisotropy of PBF nickel-based superalloys components, which represents a tremendous challenge to researchers. The dendritic growth directions (yellow arrows in Fig. 18c) follow the build direction (z). However, the growth of dendrites on both sides of the track interfaces does not show any preferential direction (Fig. 18d) [92]. In contrast, another study found that the newly-formed crystals grow into cellular dendrites in a direction which is either parallel to the original direction or rotated by $90^{\circ}$ [94]. This allows the grains to interpenetrate from one layer to another. Chlebus et al. investigated the features in the dendritic (Fig. 20a) and interdendritic (Fig. 20b) regions of AB LPBF IN718 specimens [58]. The fast heating and cooling cycles produced during PBF, results in small interdendritic regions (in the range of $\mathrm{nm}$ ). Microsegregation of some alloying elements, such as $\mathrm{Nb}$, Mo and $\mathrm{C}$, are also produced during the dendrite formation, because of the rapid cooling rate. Some chemical composition inhomogeneities can be observed in Fig. 20a, indicated by arrow 2 [58]. This segregation promotes the formation of NbC carbides and Laves phase in the interdendritic region, as shown in Fig. 20b [58]. EBM specimens show slightly different microstructures compared to\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-10} Fig. 15. Series of thermographs recording the building process of a single track [77]. From these thermographs, melt pool sizes, particles spattering tendency, thermal gradients, heating and cooling rates were extracted, illustrating the utility of thermal graphs to monitor Powder Bed Fusion processes quantitatively. the LPBF equivalents. Kirka et al. analysed the microstructure and chemical compositions of EBM IN718 specimens [96]. Different phases were identified using EDS, their chemical compositions are reported in Table 5. The micrographs of AB EBM specimens (Fig. 21) show that EBM specimens have a lower dislocation density (than LPBF materials), Laves phases and MC carbides in AB state and large disk-shaped $\gamma$ " particles (average size of $80 \mathrm{~nm}$ ). Sames et al. observed a variation in microstructure along the build direction of the AB EBM IN718 specimens (Fig. 22) [97]. The needle-shaped $\delta$ particles at the top were much coarser than those at the bottom. Additionally, the material in this area showed a greater contrast upon etching, indicating a more severe secondary element segregation. Deng et al. provided more detailed information about the precipitate morphologies as well as the microstructural variations occurring during EBM of IN718 [98]. All these results indicated that thermal cycling varied during the build, influenced by the number of layers already deposited. Hence localised HTs are a common phenomenon in PBF and should be considered in-process optimisation. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-11(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-11} \end{center} (b) Fig. 16. Analysis of thermographs. (a) Grayscaling the image captured by infrared cameras. (b) The corresponding contour plot. This shows the useuflness of using infrared cameras to locate spatter generated during the building process [82]. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-11(5)} \end{center} Fig. 17. Summary of defects present in powder bed fusion processes. This can serve as reference for commonly used terminology by practitioners [91].\\ Polonsky et al. studied the presence and morphology of fusion defects in EBM IN718 specimens [99]. Columnar grains, with a primary aspect ratio smaller than 0.2 and oriented in the build direction, were found to surround the defects on the XY plane (Fig. 23a). Instead, the regions above the defects had small equiaxed grains with almost no discernible texture. This shows that defects drive the recrystallisation phenomena and influence the resulting microstructure. Table 6 summarises the commonly observed differences in microstructure between LPBF and EBM and Fig. 24 illustrates some of these differences. Since powder fusion and recrystallisation of the melt pool are the \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-11(3)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-11(4)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-11(6)} \end{center} (c) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-11(2)} \end{center} (d)\\ Fig. 18. Images of as-built Laser Powder Bed Fused Inconel 718 specimens. (a) Side view. (b) Top view. (c) The melt pool boundaries between layers. (d) The melt pool boundaries between adjacent tracks from the side. The arrow in (a) indicates the build direction (BD) and the circle in (b) indicates the plane is perpendicular to the build direction. Melt pools in (a) and laser scan tracks in (b) are clearly visible. The yellow arrows in (c) and (d) represent the dendrite growth direction. They follow the build direction in (c) and have no preferred direction in (d) [92]. These images illustrate the particular anisotropic microstructure resulting from the Powder Bed Fusion processes. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-12(2)} \end{center} Fig. 19. Electron Backscatter Diffraction maps of an as-built Laser Powder Bed Fused Inconel 718 specimen [59]. This shows the dominant direction of grain growth and strong $<100>$ crystallographic texture in the build direction, which is characteristic of Powder Bed Fusion processes. Individual laser scan tracks with small equiaxed grains track the overlapping regions can also be seen in the XY plane (perpendicular to the building direction). central phenomenon in PBF, different building parameters and postprocessing techniques will lead to different characteristics, which can be quantified using the different methods described previously. Exploration of these in more detail is a key concern. \subsection*{3.2. Process parameters driven phenomenon} PBF process parameters naturally have an effect on the print quality and the resulting microstructure. Indeed, using optimised process parameters can suppress the formation of voids and build defects, such as micro-cracks, in LPBF [101-103] and EBM [104]. Review papers on the process-microstructure relationship for LPBF of metallic materials has been presented previously and the reader is directed to these accordingly [105-107]. Kumara et al. also investigated phase transformations of PBF IN718 [108]. An overview of the different effects of build parameters on PBF of nickel-based superalloys microstructure is given below. \subsection*{3.2.1. Powder characteristics driving build quality} Powder quality plays a key role in determining PBF components final quality. Powders can be rotary, gas or water atomized and exhibit different morphologies, particle size distributions, flowability, surface roughness and chemical composition. These can vary from supplier to supplier [109]. Sutton et al. reviewed the most commonly used powder characterisation techniques, paying attention to the impact of powder quality on final material properties [110]. A similar review by Tan et al. also focused on powder characterisation techniques, but with a particular emphasis on powder granulometry [111]. This was identified as a key method to ensure a high performance of the feedstock, leading to high quality and, importantly, dense parts. Studies have found that chemical composition had the strongest impact on the microstructure, as higher content of certain elements resulted in the precipitation of detrimental elements and prevented recrystallisation during $\mathrm{HT}$, which decreased mechanical performance, particularly fatigue [109]. Another work also found that the segregation of alloying elements in LPBF Hastelloy X during solidification resulted in variations in composition which caused cracking in the build direction [112]. By investigating four powders with different contents of various alloying elements, Mancisidor et al. achieved a defect free material [112]. The recyclability of nickel-based powders has been studied by several researchers [113-116]. These generally observed that, as long as the recycled powder is well sieved and stored, little or no difference in properties was found between the specimens manufactured using recycled and fresh powder over $\sim 10$ build cycles. However, it was found that the presence of minor alloying elements in the metal powder, can influence the crack formation mechanism in PBF specimens [117]. Sames et al. evaluated the properties of components manufactured using powders obtained from various production methods, namely gas atomiser, rotary atomiser, and plasma rotated electrode process (PREP) [118]. From the SEM observations, PREP powder showed a smoother surface and almost no internal trapped gas, compared to the others (Fig. 25a-c). These voids in the starting powder particles may lead to an increased number of porosities in the final PBF part, which is the case for \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-12} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-12(1)} \end{center} (b)\\ Fig. 20. Typical examples of Scanning Electron Microscopy images of phenomen in Additively Manufactured nickel-based superalloys. (a) An overview of the interface between adjacent layers. (b) The interdendritic region. Mark 1 indicates the layerlayer melt pool boundary, 2 indicates the dendritic cell tips, 3 and 4 highlight some $\gamma+$ Laves phase eutectic, 5 points a MC carbide [95]. These features are caused by the rapid heating and cooling during the Laser Powder Bed Fusion process. Table 5 Chemical composition of phases marked in Fig. 21 (wt \%) [96]. This provides further proof of the present of Laves phase and MC carbides in the as-built state of Electron beam melted Inconel 718. \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|} \hline Items & Phase & $\mathrm{Ni}$ & $\mathrm{Nb}$ & $\mathrm{Ti}$ & $\mathrm{Fe}$ & $\mathrm{Cr}$ & Mo & Si & C \\ \hline 1 & MC & 0.43 & 89.91 & 6.01 & 0.15 & 0.62 & - & - & 2.73 \\ \hline 2 & Laves & 42.20 & 22.87 & 0.11 & 13.54 & 11.24 & 9.53 & 0.5 & - \\ \hline 3 & MC & 0.4 & 90.38 & 6.31 & 0.17 & 0.7 & 0.71 & - & 0.99 \\ \hline 4 & Laves & 38.8 & 28.55 & 0.23 & 11.70 & 9.40 & 11.19 & 0.77 & - \\ \hline 5 & $\gamma$ matrix & 56.19 & 4.14 & 0.68 & 18.99 & 15.91 & 3.80 & 0.04 & - \\ \hline \end{tabular} \end{center} \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-13(1)} \end{center} Fig. 21. As-built Electron Beam Melted Inconel 718 microstructure and Electron Diffraction Spectroscopy results (quantified in Table 5) [96]. This shows that Laves phases and MC carbides are present in the as-built state, as well as large disk-shaped $\gamma$ ', particles. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-13(2)} \end{center} (a) Top \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-13} \end{center} (b) Bottom Fig. 22. Scanning electron micrographs of an as-built Electron Beam Melted Inconel 718 sample. (a) The top of the sample (few thermal cycles). (b) The bottom of the sample (many thermal cycles) [97]. This shows a variation in microstructure along the build direction, particularly regarding $\delta$ particles which are coarser at the top (a). the powders obtained with the first two methods (Fig. 25d and e). Another study compared powder atomisation methods and found that LPBF parts made using water atomized powder resulted in higher sample porosity than for gas atomized powder. This was thought to be caused by the more irregular morphology of water atomized powder and hence better packing density but this observation is far from conclusive [119]. However, there are still limitations to gas atomized powder, such as hollow and/or satellite balls [120]. It was also found that the laser absorption rate can be increased by increasing the surface roughness of powder particles [120]. In terms of particle size, work has found that the presence of powder particles smaller than $10 \mu \mathrm{m}$ resulted in severe agglomeration and impeded LPBF process through spreading problems [121]. Additionally, a method of rapidly characterising powders (morphology, flowability and size distribution) was developed in order to evaluate the influence of different alloy compositions on LPBF processability [122]. It is clear that many parameters combine to define powder quality, which in turn affects the PBF process and the resulting mechanical properties. Hence, it is essential to understand and control powder quality in order to produce adequate parts. Therefore, the role of the powder is critical in determining both the interaction with the incident energy beam but also in assuring spread-ability upon the powder bed. Furthermore, the recyclability of powder and its effects on mechanical properties and in-situ alloying [123,124] are emerging topics and should be investigated. There is significant opportunity to explore this space further as in many cases the economic viability of PBF processes is driven by new powder cost but also how easily it may be recycled. \subsection*{3.2.2. Controlling build environment} A review of the build environment in PBF was written by Poorganji et al. [125]. All PBF build chambers usually operate under vacuum or an inert gas (e.g. Argon or Nitrogen) in order to avoid oxidation of the part and powder. Traore et al. researched the influence of gas atmosphere on nickel-based superalloys [126]. However, despite processing in an\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-14} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-14(1)} \end{center} (b) Fig. 23. Electron Backscatter Diffraction data of grains surrounding a defect. (a) The columnar grains surrounding the fusion defect. (b) The equiaxed grains above the defect [99]. This shows that defects also affect the build process and resulting microstructural morphology in Powder Bed Fusion processes. Table 6 Summary of common microstructural differences between Laser Powder Bed Fused and Electron Beam Melted material. The reader is advised to use this with caution as process technology evolves. This table is intended as a broad guide only and observations will vary between process configurations. \begin{center} \begin{tabular}{lll} \hline Characteristics & LPBF & EBM \\ \hline Porosity & Both have similar densities in AB or post-processed conditions & \\ Grain & Elongated grain parallel to the & More columnar grain \\ morphology & \begin{tabular}{l} build direction with ill-defined \\ grain boundaries and very fine \\ \end{tabular} & \begin{tabular}{l} formation than LPBF parallel \\ to the build direction. \\ \end{tabular} \\ & \begin{tabular}{l} columnar cellular structures at \\ higher magnifications. \\ \end{tabular} & Perpendicular to the build \\ direction, grains have a more & & \\ \end{tabular} \end{center} Argon atmosphere with $<0.2 \%$, oxidation may still occur resulting in oxide inclusions in built parts and oxide spatter particles, which were in the size range to be recycled [127]. Zhao et al. investigated the role of the build environment on melt pool dynamics in EBM and LPBF [128]. The LPBF build environment, which has the high atmospheric pressure and multiple laser reflections, is the source of more build quality issues, such as vapor recoil pressure on the melt surface, than the EBM build environment. Furthermore, an investigation was conducted on the differences in surface morphology and composition during multicycle EBM with IN718 powder reuse and there was a significant change after exposing the powder to the build chamber environment [129]. Additionally, increased oxidation rates were observed initially for the EBM N06002 alloys in comparison with wrought [130]. A study confirmed $\mathrm{Al}_{2} \mathrm{O}_{3}$ particles were formed in IN718 alloys during PBF, which act as nucleation sites for the precipitation of $\mathrm{Nb} / \mathrm{Ti}$ carbides, leading to the formation of unique core-shell composites with $\mathrm{Al}_{2} \mathrm{O}_{3}$ in the centre and $\mathrm{Ti} / \mathrm{Nb}$ at the periphery [131]. Microstructures can also be influenced by controlling the EBM process temperature [132]. At $915{ }^{\circ} \mathrm{C}$, large $\delta$ needles were formed (Fig. 26a). By increasing SEM magnification (Fig. 26b) a finer $\delta$-phase $(\sim 200 \mathrm{~nm})$ distributed at the grain boundaries, as well as some isolated MC carbides ( $1.5 \mu \mathrm{m}$ ), were observed. On the other hand, the microstructure obtained at $990^{\circ} \mathrm{C}$ appeared relatively clean. Slightly coarser carbides $(\sim 3 \mu \mathrm{m})$ as well as fine $\delta$ particles were also found at the grain boundaries (Fig. 26c and d). Overall, this shows that the build environment needs to be adequately controlled to obtain defect-free and desired microstructures. \subsection*{3.2.3. Energy beam driven phenomenon} Laser parameters, such as the laser power, scan speed, hatch distance and scan strategy, are some of the main factors influencing PBF microstructures. 'Stripe', 'Meander', 'Total fill' and 'Chessboard' (also known as 'Island') strategies are some of the main scan strategies used in LPBF currently (Fig. 27). Different and customised strategies, including multilaser [133], residual heat factor [134] and 'unit-cell' strategies [135], are also being developed in order to obtain and control microstructural characteristics, such as grain morphology, density, defects, cracking, and surface quality. In a study on EBM by Helmer et al., the area energy density E [J $\mathrm{mm}^{-2}$ ] was used as a comparison parameter to evaluate the overall effects of laser power $P[W]$, scan speed $v\left[\mathrm{~m} \mathrm{~s}^{-1}\right]$ and hatch distance $H$ $[\mu \mathrm{m}]$ on grain morphology [136]. To allow a comparison, the energy density applied to two specimens was similar, respectively $1.8 \mathrm{~J} \mathrm{~mm}^{-2}$ for the first and $1.9 \mathrm{~J} \mathrm{~mm}^{-2}$ for the second specimen [94]. As expected, different values of scan speed and hatch distance produced two clearly distinct grain morphologies, as shown in Fig. 28. Additionally, in another study by Karimi et al., it was found that the electron beam focus offset also directly affected the grain morphology [137]. Fernandez-Zelaia et al. also showed that the morphology and texture of the mesoscale can be controlled by the melting sequence [138]. Similar results were found for LPBF processes [139]. Indeed, in a LPBF study, using a flat top laser beam changed grain morphology to a wide and planar geometry with a $150 \%$ increase in grain size, compared to $200 \mathrm{~W}$ Gaussian beam [140]. Sow et al. also compared a $80 \mu \mathrm{m}$ diameter Gaussian laser spot and a $500 \mu \mathrm{m}$ diameter top-hat laser beam and found that the $500 \mu \mathrm{m}$ diameter top-hat laser beam increased productivity, suppressed spatter and produced fully dense IN625 parts [141]. Fig. 29 compared the features of two LPBF IN718 specimens, produced with different laser power (250 W and $950 \mathrm{~W}$ ) [67]. The shape and size of melt pools can be easily recognized in the OM, highlighting a clear influence of the laser power. Indeed, lower power generates smaller melt pools and results in a reduced HT of underlying layers. This, combined with a consequent faster solidification, leads to smaller grains. For LPBF IN738LC, using higher laser power increased the depth of keyholes, causing instability and increasing pore formation due to the periodic collapse of the keyholes [142]. Furthermore, laser volume energy density was found to be the main parameter affecting cracking and porosity. For example, increasing the laser volume energy density resulted in an increase in number and size of cracks in the SRR99 nickel-based superalloy [143] and minimal solidification cracking was observed in IN738LC with narrow melt pools with a strong melt pool overlap [143]. In LPBF René 104 superalloy was built with 3 different strategies (meander, stripe, chessboard) and these were found to have a significant effect on cracking and relative density [142]. The scan strategies with more partitions were shown to increase the emergence of \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-15(7)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-15(4)} \end{center} (b) Fig. 24. Microstructural cubes of hot isostatically pressed Inconel 625. (a) Processed by Electron Beam Melting. (b) Laser Powder Bed Fusion. Showing the differences in microstructure produced by the different processes [100], such as grain morphology and size. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-15(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-15(2)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-15(6)} \end{center} (d) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-15(3)} \end{center} (e) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-15(5)} \end{center} (c) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-15} \end{center} (f) Fig. 25. Powders and corresponding laser powder bed fused specimens [118]. (a) Gas atomized powder at $100 \mu \mathrm{m}$ scale. (b) Rotary atomized powder at $100 \mu \mathrm{m}$ scale. (c) Plasma rotated electrode processed powder at $100 \mu \mathrm{m}$ scale. (d) Gas atomized powder at $200 \mu \mathrm{m}$ scale. (e) Rotary atomized powder at $200 \mu \mathrm{m}$ scale. (f) Plasma rotated electrode processed powder at $200 \mu \mathrm{m}$ scale. From the Scanning Electron Microscopy observations, plasma rotated electrode processed powder showed a smoother surface and almost no internal trapped gas, compared to the others, illustrating the importance and effect of powder types on the resulting microstructure of powder bed fused materials. cracks while the overlapping zone increased the size, number and frequency of cracks [142]. Likewise, residual stresses were shown to be caused by scan-strategy induced microstructure [144] and it was found that a more uniform scan strategy resulted in in less residual stresses [145]. Lee et al. studied the formation of cracks in EBM Mar-M247 parts and concluded that cracks usually form along interdendritic grain boundary at the end of solidification, due to the coexistence of liquid films and residual thermal stresses [102]. Peng et al. demonstrated that EBM induced cracks in DZ125 superalloy, classified as "liquid-state cracks", were also widely spread at interdendritic grain boundary [103]. Similar conclusions about crack formation mechanisms were drawn. It is clear from the above that the formation of cracks, or hot cracking, is an issue in PBF, but with appropriate processing parameters, these can be reduced or eliminated. Chauvet et al. studied the formation of cracks in EBM of non-weldable nickel-based superalloys and found a correlation between hot cracks and high angle grain boundaries [146]. He also found that the presence of a liquid film during the last stage of solidification and thermal stresses trigger hot cracking [146]. Marchese et al. also confirmed that high thermal residual stresses resulted in hot cracks during LPBF of Hastelloy X [147]. Part density is also influenced by the laser parameters. Indeed, a study found that the relationship between density and laser input energy during LPBF of GH3536 was found to comply with a quadratic function and presented an inverted U-shaped distribution [145]. Furthermore, results showed that in a given scanning strategy, the density decreased as the scanning speed increased for a fixed fluence [148].\\ Insufficient laser overlap (large hatch spacing) can deteriorate the surface of materials [149]. Indeed, the laser energy input improved density and surface quality of $\mathrm{Ni}-\mathrm{Cr}-\mathrm{B}-\mathrm{Si}$, with a fine grain microstructure and strengthening precipitates [150]. Attard et al. investigated these effects and produced a controlled functionally graded microstructure by varying the process parameters [151]. Finally, it is clear that the combination of scan strategies, laser power and other process parameters affect the microstructure and final mechanical properties of PBF components. In order to optimise PBF process parameters, some researchers are using the Taguchi regression method $[152,153]$ and Artificial Neural Networks [154]. \subsection*{3.3. Post-processing for component performance} Given the microstructural anisotropy and the defects generated during manufacturing, post-processing is necessary to improve the properties of most AB LPBF components. Post-processing, which includes heat treatments, surface treatments and machining processes, is a prime method to create more favourable microstructures. The purpose of post-processing is to enhance both the form and integrity of the bulk and surface of a component to elevate performance characteristics. While the intent in all AM processes is to create a component within a single step, it is inevitable that, where shortfalls are apparent, additional measures must be taken. This is not entirely inconsistent with established manufacturing routes. For example, it is a pedestrian activity in modern manufacturing to machine a casting. However, in the case of AM \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-16(7)} \end{center} (c) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-16(4)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-16(5)} \end{center} (d) Fig. 26. Scanning Electron micrographs of Electron Beam Melted Inconel 718 built with different base plate temperatures [132]. (a), (b) $915{ }^{\circ} \mathrm{C}$. (c), (d) 990 ${ }^{\circ} \mathrm{C}$. This shows the influence of process temperature on the precipitation of secondary phases. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-16(3)} \end{center} Straight line vector path from each side of the border.\\ Stripes \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-16(1)} \end{center} The area within the border is split into strips and a meander technique is used within each strip.\\ Chessboard \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-16(6)} \end{center} A further advancement on Stripes. Instead, the area is split into squares like a chessboard.\\ Total Fill \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-16} \end{center} The complete area is filled by offsetting Additional Borders. No hatch pattern is applied. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-16(2)} \end{center} Fig. 27. Typical scan strategies as demonstrated in the Renishaw 'QuantAM' material editor. These result in markedly different microstructures and mechanical properties. where subsequent processing is required the business and design case for this technology will be undermined. Hence, while often currently essential, the research community must endeavour to achieve geometrical tolerance and material condition in process. Lim et al. wrote a review on reducing residual stress in metal PBF parts [155] which informs us that while stresses may be substantially reduced by process optimisation they cannot be removed entirely within the PBF process itself. \subsection*{3.3.1. Heat treatments for enhancing mechanical properties} In industry, almost all functional AM parts in mission critical applications are post-processed using HT. HT allows the modification of the microstructure through controlled heating and cooling of materials. By modifying the microstructure, materials can obtain enhanced mechanical properties. Indeed, performance of $A B$ material is poor because of surface integrity defects and unfavourable microstructural formation [156] with secondary phases such as Laves phases that embrittle grain\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-17(2)} Fig. 28. Electron Backscatter Diffraction maps of Electron Beam Melted specimens manufactured with different parameters. (a) Parameters: $\mathrm{E}_{1}=1.8$ $\mathrm{Jmm}^{-2}, \mathrm{v}=2.2 \mathrm{~ms}^{-1}, \mathrm{H}=150 \mu \mathrm{m}$. (b) Parameters: $\mathrm{E}_{2}=1.9 \mathrm{Jmm}^{-2}, \mathrm{v}=8.8 \mathrm{~ms}^{-1}, \mathrm{H}=37.5 \mu \mathrm{m}$ [136]. This illustrates that varying process parameters can result in drastically different microstructures and textures. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-17} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-17(1)} \end{center} (b) Fig. 29. Optical micrographs of etched Laser Powder Bed Fused Inconel 718 specimens manufactured with varying laser power [67]. (a) Power $=250 \mathrm{~W}$. (b) Power $=950 \mathrm{~W}$. The melt pools across layers can be clearly observed and show that a difference in laser power can significantly affect the melt pool shape. boundaries and are detrimental for mechanical properties [95]. Hence, HTs are used to increase grain size [157], obtain a more equiaxed microstructure, dissolve detrimental phases, such as Laves [158],to form strengthening precipitates such as $\delta$-phase, $\gamma^{\prime}$ and $\gamma^{\prime \prime}[156,159]$, and to remove defects [160] to improve mechanical properties. Although sometimes $\mathrm{AB}$ microstructure is more beneficial for certain properties, like in the case of Parizia et al. who found that AB IN625 presented better oxidation resistance than its HT counterpart [161]. There are different types of HT, each giving different microstructures. Usually, samples are first stress relieved, which reduces texture and residual stress in samples [162]. Then, solution treatment (ST) is used to enhance the mechanical properties [163] by dissolving detrimental phases [164, 165]. Subsequently, samples are sometimes aged (single or double ageing) to favour the precipitation of strengthening phases [158]. Samples can be Solution treated then Aged (STA) or Directly Aged (DA). Thermal techniques, such as homogenisation and HIP treatments are also used. Homogenisation is usually used prior to HIP and is similar to stress relief as it reorients columnar grains [157]. Zhao et al. observed that during homogenisation of LPBF IN718, the grains would continue recrystallisation whereas the suction-cast alloy showed abnormal grain growth, which showed the potential of engineering the microstructure of AM materials through HTs to obtain superior mechanical properties than in conventionally manufactured alloys [166]. HIP results in recrystallisation, grain coarsening and change from highly textured columnar grains to randomly oriented equiaxed grains which are larger than after homogenisation [64,157,167,168]. Moreover, a slightly weaker texture is obtained, compared to STA [168], although it is still strong [167]. HIP was also found to be effective at closing defects, resulting in a higher density [169]. Different HT standards exist for conventionally manufactured material, however, as no PBF-specific HT standards have been defined at this time, significant research has gone into exploring the effects of wrought HT and modified HT on PBF microstructure. The Standard Specification for Additive Manufacturing Nickel Alloy (UNS N07718) with Powder Bed Fusion [170] gives guidelines for thermal processing of PBF nickel-based superalloys. For HIPing, components should be processed in an inert atmosphere at no less than 100 $\mathrm{MPa}$, within the range of $1120^{\circ} \mathrm{C}$ and $1185^{\circ} \mathrm{C}$ within $\pm 15^{\circ} \mathrm{C}$, and held for $240 \mathrm{~min} \pm 60 \mathrm{~min}$ followed by cooling under and inert atmosphere below $425^{\circ} \mathrm{C}$ [170]. For HT, it states that components should be solution treated and aged following the AMS2774 standard for Heat Treatment of wrought nickel alloy and cobalt alloy parts [171]. This standard gives the range of possible heat treatments to use for different geometries of nickel-based superalloys. It should be noted that none of these HT are specific for AM, but rather are for conventional manufacturing processes. This highlights the need for the development of PBF specific microstructures. Some practitioners, such as Huang et al. and Aydinöz et al., have started to research this area [163,172]. For example, Huang et al. investigated the effect of solution time, solution temperature, cooling method and ageing process on the mechanical properties of LPBF IN718 and identified that there was a minimum solution time for a given temperature to obtain similar microstructures and mechanical properties (Fig. 30). \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-18} \end{center} Fig. 30. Variation of required solution time for solution temperature for Laser Powder Bed Fused Inconel 718 [163]. This shows that there is a minimum solution time for a given solution temperature which results in similar microstructures and properties. This further highlights the requirement for the development of Powder Bed Fusion specific heat treatments as the standard heat treatments designed for conventionally manufactured materials are unsuitable. 3.3.1.1. Grain structure. Although grain boundaries normally occupy a small fraction of material volume, they play a crucial role in controlling material properties. The sensitivity associated with this behaviour drives significant efforts in optimising processes. Fig. 31 a and b compare the microstructures of an HT and HIP + HT LPBF IN718 specimens [173]. Obvious evidence of the scan strategy used is eliminated in both cases [173]. From the measurements, it was found that the average grain size for HT specimens was $15.5 \pm 2.0 \mu \mathrm{m}$, namely $30 \%$ smaller than HIP + HT equivalents. Holland et al. investigated the evolution of the grain boundary network in AB and HT LPBF IN718 specimens [174]. It was observed that non-specific grain boundaries dominated in the $A B$ specimens, whereas after HT the number of special grain boundaries increased significantly from $9 \%$ to around $60 \%$. This second class of grain boundaries includes twin boundaries and twin-related grain boundaries, able to improve material strength and resistance to intergranular degradation. Another study on LPBF IN625 provided similar results [175]. It was also found that above a certain annealing temperature $\left(1150{ }^{\circ} \mathrm{C}\right)$, these grain boundaries developed in prevalence significantly. "Grain boundary misorientation" is defined as the difference in crystallographic orientations between adjacent grains of the same phase. This microstructural feature can be characterised through EBSD. Gribbin et al. evaluated the misorientation angle (Fig. 32) in some HT and HIP + HT LPBF IN718 specimens [176]. The distribution of the misorientation angles for HT specimens was broad, indicating no preferential growth direction. On the other hand, for HIP + HT specimens, the distribution showed a prominent peak at $60^{\circ}$, indicating an equiaxed grain structure with a high twin content caused by annealing. The detrimental effect on fatigue performance previously observed by Zhang et al. correlates well with the misorientation effect [177]. In a further study by Chauvet et al., it was found that grain boundaries with a high misorientation angle were prone to crack propagation in AB and HT EBM samples (Fig. 33) [146]. Han et al. obtained similar results in this regard [178]. Research also concluded that the difference in interdendritic liquid pressure between the dendrite tip and root, as illustrated in Fig. 34, would cause an insufficient feeding of molten material at the dendrite root, promoting void generation and therefore highly affecting the part hot cracking behaviour. Tomus et al. compared the grain morphology of LPBF Hastelloy X specimens processed with various post-processing techniques [179]. The HT consisted in a single solution step $\left(1175{ }^{\circ} \mathrm{C} / 2 \mathrm{~h}\right)$, while HIP was performed using the same time and temperature, with an applied stress of $150 \mathrm{MPa}$. Fig. 35a,c,e,g display a series of EBSD images illustrating the grain morphology in the XZ plane for AB, HT, HIP and HIP + HT specimens, respectively. HT and HIP effectively reduced the strong texture in the build direction in the AB specimen [180]. Another study found that HIP of LPBF Hastelloy X 'closed' internal cracks, reduced porosity and generated equiaxed grains [181]. This was also observed in CMSX-4 [77]. Fig. 35b,d,f,h illustrates the grain morphology perpendicular to the build direction (XY plane). HIP specimens showed a smaller gran size, because of recrystallisation. Similarly to previous results, individual laser scan tracks and small equiaxed grains were also observed in the XY planes [59]. Further studies on post-processing techniques indicated that HIP cannot be considered as an efficient tool to heal EBM induced cracks [182]. 3.3.1.2. Precipitate formation. Fig. $31 \mathrm{a}$ and $\mathrm{b}$ compare the microstructures of an HT and HIP + HT LPBF IN718 specimens and showed that 'white' precipitates are clearly visible at grain boundaries [173]. EDS observations (Fig. 31) indicated that these are rich in Mo, Nb, W and Si, with stoichiometric ratios of ( $\mathrm{MoNbW}_{5} \mathrm{Si}_{3}$ [173]. However, the precipitates in both specimens were similar in size $(\sim 2.5 \mu \mathrm{m})$ [173]. Similarly, Sames et al. investigated the effects of in-situ HT on $\gamma^{\prime} / \gamma^{\prime}$, phases in EBM IN718 specimens [183]. The $\gamma^{\prime} / \gamma^{\prime \prime}$ phases in the AB specimen showed an elongated disk shape, with a diameter of $\sim 20 \mathrm{~nm}$ and a thickness of $\sim 10 \mathrm{~nm}$ (Fig. 36a). From the micrographs comparison in Fig. 36, both the diameter and thickness of these strengthening particles increased during the in-situ HT. Since strengthening phases have an optimal size range and corresponding mechanical properties, this process was found to be effective in improving material strength [183] by impeding the dislocation movement at the grain boundaries [173]. Furthermore, Divya et al. investigated the HT effects on dislocations and strengthening particles in LPBF CM237LC specimens [184]. In the $\mathrm{AB}$ specimens (Fig. 37a), dislocations entangled and tended to accumulate at the grain boundaries. As shown in Fig. 37d, the HT decreased the dislocation network density, especially at the grain centre. These observations are in accordance with results by Tucho et al. [185]. At the same time, the HT significantly increased the size of the $\gamma^{\prime}$ phase. In fact, before HT, two distinct types of $\gamma^{\prime}$ phase could be observed: one with a size of $\sim 5 \mathrm{~nm}$ (Fig. 37b) and another, much larger, with a size $\sim 50 \mathrm{~nm}$ (Fig. 37c). After the HT (Fig. 37d and e), the primary $\gamma^{\prime}$ particles reached a size of over $500 \mathrm{~nm}$, while the secondary $\gamma^{\prime}$ particles, characterised by a cuboidal morphology, had a size of $\sim 200-400 \mathrm{~nm}$. Fine tertiary $\gamma^{\prime}$ particles were spread in the region between the secondary $\gamma^{\prime}$ particles. The influence on $\gamma^{\prime}$ particles size may the basis of the strengthening mechanisms caused by HT. During a 3-step HT on LPBF Haynes 282, $\gamma^{\prime}$ precipitation was found at $950{ }^{\circ} \mathrm{C}$ during TEM in-situ HT [186]. After $\mathrm{HT}$, the morphology and size of $\gamma^{\prime}$ precipitates were comparable to powder metallurgy samples and annealing twins were present [78]. HT was also optimised for LPBF CMSX-4 to obtain segregation of $\gamma / \gamma$, microstructure [52]. Kuo et al. evaluated the effects of different HT strategies on the $\delta$-phase in LPBF IN718 specimens [92]. In the AB specimen, $\delta$-phase was found distributed parallel to the build direction, segregated in the interdendritic region due to the $\mathrm{Nb}$ segregation during the build process (Fig. 38a). This was hypothesized to be a consequence of $\mathrm{Nb}$ segregation which occurred during the LPBF process. Specimens which underwent a solution treatment and aging (STA) $\left(980^{\circ} \mathrm{C} / 1 \mathrm{~h}\right.$ then $718^{\circ} \mathrm{C} / 8 \mathrm{~h}+621$ ${ }^{\circ} \mathrm{C} / 10 \mathrm{~h}$ ) possessed a much coarser $\delta$-phase than their non-solution-treated equivalents (Fig. 38b and c, respectively). This difference can be related to the dissolution of $\gamma$ " phase and the consequent formation of needle-shaped $\delta$-phase during this first thermal treatment. However, these elongated particles are undesirable since they degrade material mechanical properties, causing " $\delta$-phase \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-19(1)} \end{center} (a)\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-19(2)} \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-19(3)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-19} \end{center} Fig. 31. Microstructure of Laser Powder Bed Fused Inconel 718 specimens with varyinig thermal treatments [173]. (a) Heat treated sample micrsotructure with Electron Diffraction Spectroscopy results of the section squared in (a). (b) Hot Isostatically Pressed and heat treated sample. This shows obvious evidence of the elimination of scan strategy effects through thermal treatments. \section*{embrittlement".} Stoudt et al. presented a time-temperature-transformation diagram for the $\delta$-phase in LPBF and wrought IN625 specimens [187]. The formation of $\delta$-phase during a LPBF process was found to be much faster than in the conventional wrought process. Moreover, it was observed that the stress-relief HT, normally used in industry for IN625 (870 ${ }^{\circ} \mathrm{C} / 1 \mathrm{~h}$, red dot in Fig. 39), would promote the formation of 8 -phase during LPBF and not for wrought equivalents. Zhang et al. also proved that conducting this same $\mathrm{HT}$ at $800{ }^{\circ} \mathrm{C}$ can lead to the nucleation and growth of $\delta$-phase [188]. The calculated activation energy for the growth of the $\delta$-phase was found to be $(131 \pm 0.69) \mathrm{kJ} \mathrm{mol}^{-1}$. Another study also designed a two-step ST with a two-step aging treatment which facilitated the precipitation of $\delta$ phase at the grain boundaries [189]. These results highlight the fact that HT conditions for LPBF processes need to be re-evaluated and distinguished from those used for conventional manufacturing methods. Laves phases are another common precipitate which is known to be detrimental to the mechanical properties of nickel-based superalloys. Indeed, Laves phases subtract $\mathrm{Nb}$ from the two main strengthening phases, namely $\gamma$ " and $\delta$. Pröbstle et al. explained that, in agreement with other studies, only the Laves phases were visible on TEM of $A B$ LPBF IN718 (Fig. 40) because of the rapid heating and cooling cycles, which suppressed other secondary phase precipitation [190]. Hence, HT is necessary to dissolve these and create more wrought like microstructures. TEM micrographs and diffraction patterns for each specimen are shown in Fig. 41. In the solution treated specimen (Fig. 41a), the associated diffraction pattern (Fig. 41b) indicated that there were no secondary phases. This means that the Laves phases, which are commonly observed in $A B$ specimens (Fig. 40), were fully dissolved during the solution treatment. In the HIP specimen, intense recrystallisation occurred (Fig. 41c and d) due to the high temperature and deformation induced, dissolving all substructures present after build. As for the previous case, no secondary phases were revealed from the diffraction patterns, suggesting a complete dissolution of secondary phases during HIP. As stated, a combination of solution treatment and ageing represents one of the most commonly used HT strategies for IN718. Both TEM micrographs and diffraction patterns for this strategy (Fig. 41e and f) revealed the presence of $\gamma$ " particles with a size $\sim 30 \mathrm{~nm}$. Similarly, for HIP + ageing treated specimens, reflections of $\gamma$ ' phases were observed in the diffraction pattern (Fig. 41i). From the TEM micrographs (Fig. $41 \mathrm{~g}$ and h), needle-shaped $\delta$ particles were individuated at the grain boundaries. These observations were similar to those made by Kuo et al. [92]. However, in this case, the needle-shaped $\delta$ precipitates were also found to lower the specimen strength. These results confirmed that the precipitation of $\delta$ particles reduced the amount of $\gamma$ " present in the surrounding area (Fig. $41 \mathrm{~h}$ ). Despite the general consensus that Laves phases are detrimental to \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-20} \end{center} Misorientation Angle \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-20(4)} \end{center} Misorientation Angle [degrees] (a)\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-20(2)} (b) Fig. 32. Electron Backscatter Diffraction maps of Laser Powder Bed Fused Inconel 718 specimens showing grain morphology and misorientation angle distributions [176]. (a) Heat-treated sample. (b) Hot Isostatically Pressed and heat treated sample. The broad distribution in the heat treated sample indicates no preferred growth direction as compared to the Hot Isostatically Pressed and heat treated sample. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-20(3)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-20(1)} \end{center} (b) Fig. 33. This figure demonstrates that grain boundaries with a high misorientation angle are prone to crack propagation in as-built and heat treated Electron Beam Melted samples [146]. (a) An Electron Backscatter Diffraction map showing the cracked grain boundary along the high angle grain boundary (misorientation $>15^{\circ}$ ). (b) The distribution of grain boundary misorientation and cracked grain boundaries. mechanical properties and need to be dissolves, recent studies have found that the size, morphology and distribution of Laves phases can prove beneficial to the mechanical properties of PBF nickel-based superalloys [191-193]. For example, Sui et al. managed to dissolve the sharp corners and grooves of the Laves phase through HT, causing it to change from a long-striped to a granular shape [191]. They then found in another study that the granular Laves phases were more beneficial to the plastic deformation of PBF IN718 than long-striped Laves phases and that a certain amount of Laves phase was the best match between strength and ductility of the sample [192]. Similarly, Xiao et al. found that fine discrete Laves phase improved the tensile properties of LPBF IN718, even outperforming wrought IN718, and had good elongation, \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-21} \end{center} Fig. 34. Illustration of the hot cracking mechanism in the Laser Powder Bed Fusion process [178], demonstrating crack formation and growth within a single melt pool. This shows that the difference in interdendritic liquid pressure between the dendrite tip and the root causes an insufficient feeding of molten material at the dendrite root, promoting void generation and therefore highly affecting the hot cracking behaviour of the part. whereas long-chain-like Laves phase had a more brittle nature and suboptimal properties [193]. Overall, an optimised HT can control the size, shape and distribution of precipitates to cater to mechanical properties desired. More work should be undertaken to explore this aspect. 3.3.1.3. Residual stress. Tucho et al. demonstrated that thermal cycling during the building process induced residual stresses in the material (a common observation in energy beam processes), producing plastic deformation and dislocation networks (Fig. 42a) [185]. However, since these are caused by internal stresses, the dislocation networks can be removed using an appropriate HT, as shown in Fig. 42b. Overall, with the appropriate HT, it is possible to obtain a microstructure which resembles that of a cast nickel-based superalloy, as was the case for LPBF high-strength alloy VZhL21 after progressive stages of post-treatment [194] and Hastelloy X after solution annealing [181]. However, a combination of ductility dip cracking and strain age cracking mechanisms were identified as the primary causes of cracking in LPBF CM247LC following post-build thermal treatments [195]. This shows that HT still requires optimisation in order to obtain a defect free LPBF material. \subsection*{3.3.2. PBF surface integrity following machining} Machining is often required to obtain the desired geometry following PBF. AM components present new machining challenges given material inhomogeneity and intricate geometries. Given the implicit part-to-part variation traditional datum acquisition challenges associated with casting are apparent. However, given the limitations of current class PBF systems it is likely that additional value add to AM components will be derived by machining processes. Machining processes affect materials' microstructure, surface quality and induce residual stress. As stated previously, PBF materials have different microstructures, surface roughness and residual stresses than conventionally cast or wrought nickel-based superalloys. A review on the machinability of conventionally manufactured nickel-based superalloys was conducted by Ezugwu et al. and discusses the issues with the machining of nickel-based alloys and the cause of tool wear and failure [196]. Hence, the effects of machining will be different and it is important to understand their impact in order to control part quality. A study compared the effect of different post-processing techniques namely barrel finishing, ultrasonic shot peening, ultrasonic impact treatment and shot peening - and their effects on surface roughness, hardness and residual porosity [197]. Ultrasonic impact treatment had the best reduction in surface roughness (by $57.4 \%$ ) and in residual porosity (by 84\%), while shot peening improved hardness the most (by 66.5\%) [197]. Shot peening also refined subsurface grains in EBM IN718 and improved surface texture and oxidation performance [198]. A different study also showed that shot peening and ultrasonic impact treatment improved the surface texture parameters and residual stresses of HIP LPBF IN718 [199]. Further, Kuner et al. also found that polishing an AB EBM Hastelloy resulted in slower oxidation kinetics compared to the non-polished sample [200] and Karthick et al. observed a superior surface finish, reduced porosity and improved compressive residual stress in samples that were post-processed using grinding followed by low plasticity burnishing compared to other samples [201]. Furthermore, the use of electropolishing surface treatment with anhydrous electrolyte solution was studied to improve the surface quality of LPBF IN718 [202]. The results clearly indicated the potential benefit of introducing highly regulated electrolyte flow in the polishing of AM metal parts [202]. The effect of LPBF Hastelloy X microstructure on Electromechanical dissolution characteristics was also studied and showed that compared to wrought, LPBF finer grains, denser sub-grain boundaries and dislocations contributed to the formation of a more stable and thicker passivation film [203]. Studies have shown that the PBF microstructure of nickel-based superalloys has implications for the machining process as well. For example, there are peculiar interactions between build orientation and machining strategy [204]. It has been shown that the surface topography and integrity of LPBF IN625 was affected by the relative orientation of cutting direction to the build direction and scan strategy orientation [205]. Indeed, Patel et al. showed that machining with the feed in the build direction generated the greatest cutting force (as shown in Fig. 43) of the orientations tested [205]. Similarly, another study found that feeding the cutter against the build direction resulted in lower peak forces with larger deviations while feeding along the build direction resulted in higher peak forces with lower deviations [206]. Further, LPBF IN718 with HIP and HT were found to have better minimum specific cutting energy, minimum tool wear and minimum surface roughness during milling than wrought IN718 [207]. The peak milling cutting force was found to be dependent upon the feed direction as well as the layer-wise scan rotation employed in fabricating LPBF IN625 [206]. Hence, these studies reveal that, in a similar theme as "Design for Manufacture", the need to select PBF build parameters for post-processing needs to be considered [204]. No studies on the effect of EBM microstructure on machining parameters and vice versa were reported to date. Hence, this area should be developed further. Another development area is the use of hybrid machines which build and machine the part during the same process [208,209]. For example, using a new hybrid method which combines LPBF and Laser Shock Peening, a 95\% decrease in CM247LC cracks was obtained [102]. Hence, more research should be conducted on hybrid machines as they have the potential to further control the microstructure and mechanical properties of PBF materials. \subsection*{3.4. Modelling of nickel-based superalloys in PBF} Numerical modelling is a useful tool to understand the fundamental mechanisms and predict the possible outcomes of PBF processes. A review on the research progress of LPBF nickel-based superalloys simulation was conducted by Qiu et al. [210]. Other reviews investigate multi-scale modelling for PBF [211], the classification of AM modelling [212], microstructure modelling of metal AM materials [213], multi-physics continuum modelling approaches for metal powder [214, 215], have also been presented. Some typical models used in modelling AM microstructure include: thermal modelling, phase field modelling, kinematic modelling and cellular automata. These models can be used individually or coupled with other models to model PBF process and microstructure. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-22} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-22(5)} \end{center} (c) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-22(2)} \end{center} (e) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-22(1)} \end{center} (g) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-22(4)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-22(3)} \end{center} (d)\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-22(6)} (h)\\ Fig. 35. Grain morphology of planes parallel (XZ) and perpendicular (XY) to the build direction with varying thermal treatments [179]. (a) As-built specimen parallel to the build direction. (b) As-built specimen perpendicular to the build direction. (c) Heat treated specimen parallel to the build direction. (d) Heat treated specimen perpendicular to the build direction. (e) Hot Isostatically Pressed specimen parallel to the build direction. (f) Hot Isostatically Pressed specimen perpendicular to the build direction. (g) Hot Isostatically Pressed and heat treated specimen parallel to the build direction. (h) Hot Isostatically Pressed and heat treated specimen perpendicular to the build direction. Heat treatment and Hot isostatic pressing effectively reduced the strong texture in the build direction in as-built specimens. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-23(9)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-23(6)} \end{center} (b) Fig. 36. Transmission Electron Microscopy visualisation of $\boldsymbol{\gamma}$ " precipitates in Electron Beam Melted Inconel 718 specimens [183]. (a) As-built samples (low cooling rate). (b) In-situ heat treated samples. This illustrates the effect of heat treatment on the size of strengthening precipitates, $\gamma$ " in this case. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-23(5)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-23(2)} \end{center} (d) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-23(7)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-23(8)} \end{center} (e) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-23(4)} \end{center} (c) Fig. 37. Dislocations and $\gamma^{\prime}$ phases in Laser Powder Bed Fused CM247LC specimens [184]. (a), (b), (c) As-built. (d), (e) Heat treated This shows that dislocations are entangled and tend to accumulate at the grain boundaries. Heat treatment also decreases the dislocation network density, especially at the grain centre. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-23(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-23(3)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-23} \end{center} (c)\\ Fig. 38. $\delta$-phases in the interdendritic region of Laser Powder Bed Fused Inconel 718 specimens [92]. (a) As-Built with Nb segregation during the Laser Powder Bed Fusion process. (b) Solution + ageing treated. (c) Direct aged. The solution treated and aged specimen contained much coarser $\delta$-phase than their non-solution-treated equivalents due to the dissolution of $\gamma$ " phase and the consequent formation of needle-shaped $\delta$-phase during this first thermal treatment. This shows that thermal treatments affect the presence, size and morphology of precipitates in Powder Bed Fused materials.\\ Thermal models can determine the temperature in the material during AM processes and calculate fluid flow and porosity. These models are some of the most used in AM modelling. Zhang et al. simulated the temperature gradient and the cooling rate at the edge of the melt pool for LPBF IN718 using COMSOL Multiphysics ${ }^{\mathrm{TM}} 5.0$ software [216]. A similar study was conducted by Kirka et al., simulating the thermal profile of a laser track on solidified materials [96]. The results showed that up to five layers underneath can be remelted when a new layer is processed. This can be observed from the dendritic structure refinement found in the last remelted region and its homogenisation in the heat affected zone. Xia et al. also investigated the melt pool temperature contour in the manufacturing process of IN718+WC with LPBF [65]. This helped to explain the formation mechanism of $\mathrm{Ni}_{2} \mathrm{~W}_{4} \mathrm{C}$ primary dendrite and (Nb, $\mathrm{M}) \mathrm{C}$ carbides. Temporal evolution of the temperature \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-24} \end{center} Fig. 39. A time-temperature-transformation diagram for the formation of 8-phase in Laser Powder Bed Fused and wrought Inconel 625 components. The red dot indicates the industry recommended stress-relief Heat Treatment conditions [187]. This shows that the formation of $\delta$-phase is much faster during Laser Powder Bed Fusion than for the wrought process. The stress relief heat treatment (red dot) was shown to promote the formation of $\delta$-phase during Laser powder bed fusion but not for wrought equivalents. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.) distribution for the single bead experiments was simulated for EBM IN718 using FEA with thermal conduction and recoalescence taken into account [217]. Contrary to the experimental observations, the constructed microstructure consisted mostly of equiaxed and mixed grains [217]. Phase field modelling can model both solid and liquid material phases in the same model and is used to observe microstructure evolution (e.g. grain coarsening and dendrite growth). Pinomaa et al. used this method to simulate the kinetics of solidification, including the development of microstructural features including grain morphology, solute distribution, and formation of metastable phases - was able to accurately model temperature distribution, history, thermal gradients and cooling rates of a LPBF nickel-based superalloy [218].\\ Transport phenomena models are also used to study solidification, residual stresses, distoration, defect formation and the evolution of microstructure and properties of AM alloys, as reviewed by Wei et al. [219]. Huynh et al. [220] simulated the stress distribution in novel test pieces to prove desired results could be achieved with customized geometries. These different models can also be used to understand and predict the effect of PBF process parameters on the material microstructure or mechanical properties. For example, Raghavan et al. aimed to create a simulation to predict the effects of various EBM processing parameters on some IN718 specimens microstructure [221]. As expected, both the thermal gradient and solid-liquid interface velocity, generated during melt pool solidification, influenced the final grain morphology (Fig. 44). The same process was then simulated varying some process parameters, such as preheat temperature, spot ON time, beam diameter and spot beam current (Fig. 45), analysing their impact on the morphology produced. Other studies use models to determine the PBF manufacturability of certain nickel-based superalloys, like Yang et al., who determined the feasibility of manufacturing by LPBF nickel-based SX-superalloys by calculating the solidification conditions (temperature field, thermal gradient and solidification speed) of multi-track samples using an established finite element model based on the columnar to equiaxed transition [222]. Using models to optimise PBF parameters is also being researched. A universal and simplified model has been proposed to predict the energy density suitable for LPBF of a variety of metallic materials including nickel-based superalloy, using the relationship between energy absorption and consumption during LPBF [223]. Results confirmed that the model can predict suitable laser energy densities needed for processing materials without tedious trial and error experiments [223]. A full process energy prediction diagram for LPBF GH3536 alloy, based on the simulated molten pool depth and width, is also proposed as a method for the selecting process parameters [224]. Yan et al. also showed that using data-driven multi-scale and multi-physics models can be used to derive process-structure-property relationships for AM and optimise process parameter [225]. Finally, other researchers concentrate their efforts on the simulation of other aspects of the PBF process: powder bed melting [226-234], melt pool fluid dynamics [235,236], phase transitions [237] and microscale thermodynamic and kinetic mechanisms [235,236,238,239]. By simulating the microstructure development during PBF buildings, can provide useful insights to the morphology and distribution of dendrites and precipitates in the final part [240-242]. The simulation of residual stresses has also produced some results [243-246] of note but is also a clearly developing field prime for expansion. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-24(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-24(2)} \end{center} (b) Fig. 40. Transmission Electron Microscopy images showing the laves phase in the as-built Laser Powder Bed Fused Inconel 718 [190]. (a) Bright field. (b) dark field. Only Laves phases are visible due to the fast heating and cooling cycles, which suppressed other secondary phase precipitation during Laser Powder Bed Fusion. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-25(6)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-25(1)} \end{center} (c) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-25(8)} \end{center} (e) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-25(7)} \end{center} (g) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-25(4)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-25(3)} \end{center} (d) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-25(2)} \end{center} (f) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-25} \end{center} (h) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-25(5)} \end{center} (i) Fig. 41. Transmission Electron Microscopy images and diffraction pattern of Laser powder bed fused specimens [172]. (a), (b) Solution treated. (c), (d) Hot Isostatically Pressed. (e), (f) Solution treated and aged. (g)-(i) Hot isostatically pressed and aged. After solution treatment, diffraction shows the dissolution of Laves phases usually present in as-built condition. Following Hot Isostatic Pressing, grains are recrystallised and secondary phases dissolved. After solution and ageing treatment, the Transmission Electron Microscopy images and diffraction patterns reveal the presence of secondary phases, such as $\gamma^{\prime \prime}$. This demonstrates that solution and ageing treatments are able to precipitate secondary phases. The authors would also like to highlight the current gulf in capability between modelling times and the effective process speeds. There remain significant challenges to overcome therefore in arriving at models which are sufficiently computationally efficient to allow 'on-the-fly' model and control architectures to be deployed. \section*{4. Mechanical properties} The characterisation of mechanical properties is essential before AM components can safely be used in applications beyond the static. A review of the mechanical properties of metal AM parts was written by Lewandowski et al. [4] should the reader require broader context. Fig. 46 shows the mechanical properties of PBF nickel-based superalloys studied to date. At the current stage, studies have mainly focused on tensile and hardness performance and less on shear, toughness, fatigue and creep properties. Table 7 captures which studies looked into the different mechanical properties. It also lists the common standards usually used to test the mechanical properties. However, it should be noted that it is common for researchers to adapt standards to specific needs of AM part testing. The controlling factors for the mechanical \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-26(3)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-26(1)} \end{center} (b) Fig. 42. Bright field TEM images showing dislocations in Laser powder bed fused IN718 specimens [185]. (a) As-built. (b) Heat treated. This reveals the presence of Laves phase. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-26(4)} \end{center} Fig. 43. This shows the effect of the PBF build orientation on the subsequent machining. The greatest cutting force is generated when the feed direction is parallel to the build direction, which gives rise to anisotropy at machined faces. properties of LPBF manufactured nickel-based superalloys are analysed as follows. \subsection*{4.1. Tensile properties} Tensile testing is one of the more commonly used mechanical tests which allows the determination of the elasticity modulus, yield strength, ultimate tensile stress, ductility and stress and strain relationship for the material. These are ordinarily undertaken in a uniaxial form and as such, do not closely resemble real world loading cases. Tensile tests can easily be performed at room temperature according to standards listed in Table 7. Since nickel-based superalloys are used at elevated temperatures, their high-temperature tensile properties also need to be carefully considered according to the relevant standards. For example, some studies tested IN718 specimens at $650{ }^{\circ} \mathrm{C}[67,258,277,296]$, IN738LC at $850{ }^{\circ} \mathrm{C}[249,251,267]$, Hastelloy at $750{ }^{\circ} \mathrm{C}[255,284]$, in IN625 at 815 ${ }^{\circ} \mathrm{C}$ [333] and IN625 at $538{ }^{\circ} \mathrm{C}$ [292] and $760{ }^{\circ} \mathrm{C}$ [263,269]. A study found that high temperature tensile tests resulted in intergranular cracking [331]. Some of the most commonly used tensile testing specimen geometries are defined by the ASTM standards (Fig. 47a-c) and the ISO and EN standards (Fig. 47d). Specific countries also have their own standards, which were developed by referring to the ASTM and ISO standards. Using the standard testing procedures allows the AM specimens tensile properties to be obtained and compared between different authors. Gonzalez et al. compared the tensile properties of specimens fabricated with LPBF, EBM and binder jetting [100]. The results indicated that all AM methods surpassed the ASTM F3056-14 standard mechanical properties requirements and that LPBF specimens possessed slightly better performance and isotropy than the other two processes. 4.1.1. The influence of powder composition, quality and thickness on tensile properties Powder quality also plays an important role in the PBF process. Sames et al. compared the tensile properties of specimens manufactured with gas atomized, rotary atomized, and plasma rotated electrode process (PREP) powders [118]. The former two contain trapped gas, leading to increased porosities in the specimen, while PREP powders resulted in higher relative density and thus, higher tensile strength. Nguyen et al. compared the influence of the use of fresh and recycled powder on final properties [115]. Recycled powder particles had a slightly larger \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-26(2)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-26} \end{center} (b) Fig. 44. Thermal gradient and solid-liquid interface velocity generated during melt pool solidification [221]. (a) Relationship between Temperature Gradient (G)/Liquid-solid interface velocity (R) and solidification time. (b) Example of solidification path. This shows that both the thermal gradient and solid-liquid interface velocity, generated during melt pool solidification, influenced the final grain morphology. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-27(3)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-27(2)} \end{center} (c) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-27(1)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-27} \end{center} (d) Fig. 45. Effects of process parameters on grain morphology by controlling the temperature gradient (G) and liquid-solid interface velocity (R) of the melt pool [221]. (a) Preheat temperature. (b) Spot ON time. (c) Beam diameter. (d) Spot beam current. All of these parameters influence the final melt pool morphology. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-27(4)} \end{center} Fig. 46. Pie chart summarising the mechanical properties investigated in powder bed fused nickel-based superalloys research, across 290 papers. Almost half of the studies investigated tensile properties, while a third studied hardness properties. Other properties, especially shear and toughness, have been the subject of limited studies. average size, lower Hall flow rate and lower packing density, with respect to the fresh powder. This was due to particles' tendency to stick together and deform during the building process. Despite the differences in powder quality, their effect on the mechanical properties was insignificant in this case. However, Gasper et al. demonstrated the mechanisms by which Hastelloy $\mathrm{X}$ powders degrade in powder bed fusion [127]. Another study looked at the effect of minor alloying elements on microcrack formation in LPBF Hastelloy $\mathrm{X}$ and the influence of hot cracking on tensile and compressive properties [334]. They found that a reduction in minor alloying elements resulted in the elimination of hot cracking in AB LPBF Hastelloy X but reduced the overall tensile strength by $140 \mathrm{MPa}$ [334]. Similarly, the effect of graphene nanoplatelets (GPNs) reinforced K418 nickel-based superalloy composites, fabricated by LPBF, on the tensile properties was studied [338]. Some work revealed that introducing a $1 \mathrm{wt} \%$ titanium carbide (TiC) nanoparticles in LPBF Hastelloy X eliminated microcracks and increased yield strength by $98 \mathrm{MPa}[340]$. Fabricating metal matrix composites with PBF can augment material properties. Indeed, Yao et al. showed that the tensile strength of AB LPBF IN718 specimens could be improved by adding TiC nano-particles, which hindered dislocation movements [271]. However, the transformation of these particles from $\mathrm{MC}$ to $\mathrm{M}_{23} \mathrm{C}_{6}$ carbides during $\mathrm{HT}$, counteracted the already existing strengthening mechanism, resulting in a reduced tensile strength, and improved ductility. Alternatively, Xia et al. chose WC particles to strengthen the LPBF IN718 specimens [65]. By increasing the scan speed, WC particles and dendritic structures became much finer, increasing the tensile strength and reducing the ductility. Other particles have also been tested to reinforce PBF specimens, including graphene nanoplatelets [262] and carbon nanotubes [266]. Furthermore, the thickness of the powder layer can also affect tensile performance. Indeed, Sufiiarov et al. [274], observed that a thinner layer thickness contributed to better tensile strength since it could provide better bonding between adjacent layers. However, Nayak et al. observed that the tensile performance of $100 \mu \mathrm{m}$ layers was similar to that of thinner layers [416]. Zhou et al. manufactured functionally graded materials with strong bonding between 316L and IN718 powders. A relatively strong shear strength $(581 \pm 11 \mathrm{MPa})$ was obtained, probably due to the strong metallurgical bonds generated during LPBF [265]. Furthermore, Muñoz-Moreno et al. characterised the bulk elastic properties and shear moduli of AB and HT LPBF CM247LC components using resonant ultrasound spectroscopy [311]. The difference between the shear modulus of the $\mathrm{AB}(85 \mathrm{GPa}$ ) and $\mathrm{HT}$ specimens (87-88 GPa) was found to be negligible. Whereas Sabelki et al. found that both HT and build direction affected the torsional properties of LPBF IN718 [342]. The evaluation of shear strength is important for many components working in safety critical scenarios. Nevertheless, the number of studies related to the shear strength of PBF manufactured nickel-based superalloys is very limited. Table 7 Summary of the standards used during mechanical testing and references that investigated these different mechanical properties. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-28(2)} \end{center} 4.1.2. Build orientation results in anisotropic tensile properties and controls the failure mode Build direction is yet another important factor that highly affects specimens mechanical properties and results in microstructural and mechanical anisotropy [417]. Chlebus et al. investigated the tensile properties of specimens built in four directions (Fig. 48) [95]. The results indicated that the specimens built in the $45^{\circ} \times 45^{\circ}$ direction possessed the best tensile strength in both $\mathrm{AB}$ and $\mathrm{HT}$ conditions. Moreover, specimens built in transverse directions were always stronger than the longitudinally built equivalents. This was explained by the angle between the loading direction and grain growth direction, which can greatly affect specimen tensile behaviour. $\mathrm{Ni}$ et al. also compared the tensile strength of longitudinally and transversely built IN718 specimens considering the Schmid factor, which is used to describe the relationship between slip planes and slip direction [259]. This research found that the transversely built specimens possessed better tensile strength, which was in accordance with Chlebus et al. [95]. In another study, the differences in strength between longitudinally and transversely built specimens were believed to be caused by the angle between the loading direction and the orientation of the pores present between adjacent planes, due to lack-of-fusion [260]. The fractures for the longitudinally and transversely built samples were controlled by two different failure modes: the 'opening mode' (Mode I), \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-28(1)} \end{center} Fig. 48. Demonstration of the commonly applied layout of powder bed fused specimens with respect to the machine axis [95]. These layouts, or build orientations, affect the subsequent mechanical properties.\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-28} (d) Fig. 47. Commonly used tensile testing geometries defined by different standards. (a), (b), (c) ASTM standards. (d) ISO/EN standards. Using these geometries and following standards allows comparison between results from different authors.\\ which occurs when the loading direction is normal to the defect/pore orientation, leading to fast failure, and the 'in-plane shear mode' (Mode II), which happens when the loading direction is instead parallel to the pores orientation, typical of transversely built specimens. Similar results were obtained in a series of studies performed at room and elevated temperatures [249,261,263,277,296] with AB [247,257,258,274,276, 279,288] and post-processed specimens [64,92,264]. However, other studies reported that specimens built in the longitudinal direction had better mechanical resistance. Indeed, Tomus et al. found that, when tested at $900{ }^{\circ} \mathrm{C}$, specimens built in the longitudinal direction performed better [117]. This was also the case for LPBF of Haynes 282 at $900{ }^{\circ} \mathrm{C}$ and $20^{\circ} \mathrm{C}$ [104]. However, specimens built in the transverse direction were found to possess higher tensile strength at room temperature and $700{ }^{\circ} \mathrm{C}$. No solid explanation for these observations was given in this study. Similar results obtained by Kirka et al. could not explain the origin of the differences between longitudinally and transversely built specimens [297]. The authors reported that, by increasing the build height, transverse specimens became stronger and more ductile. This might have been caused by the consistent heat input which transformed $\gamma$ " precipitates into brittle $\delta$-phase, in the bottom of the structure. Finally, a comprehensive study relating the build direction to the anisotropic behaviour of LPBF specimens was performed by Hovig et al. [314]. \subsection*{4.1.3. Scan strategies and the effect on tensile properties} The tensile properties of PBF manufactured specimens are mainly controlled by the build parameters, with scan strategy (the path that the energy beam takes) being one of the most important. Indeed, beam power and hatch spacing were found to be the principle factors driving tensile strength [336]. Kirka et al. compared a point heat source fill scan strategy, with the conventional raster scan strategy applied in EBM [298]. The former strategy contributed to a more equiaxed microstructure, almost eliminating tensile strength anisotropy. Based on this result, Zhou et al. applied an improved alternative scanning strategy to a functionally graded component, which resulted in improved tensile strength [265]. Additionally, Geiger et al. compared the effects of three different scan strategies (labelled A, B and C and illustrated in Fig. 49) on LPBF IN738LC components [267]. The EBSD results (Fig. 49) showed that different scan strategies generate different microstructural features, which cannot be fully eliminated by HT. The tensile testing results showed that scan strategy B always developed the highest Young's modulus among all the applied strategies. Moreover, sample tensile properties are also affected by their locations on the building plate and shielding gas flow. \subsection*{4.1.4. Heat treatments for improved tensile properties} In order to produce more homogeneous microstructures and improve components mechanical properties, post-processing is still necessary at the present stage. HTs' main purposes are: decrease anisotropy, dissolve potentially crack-initiating particles and form strengthening precipitates. For more information about typical thermal treatments of PBF nickel-based superalloys and achievable optimum properties, the reader is directed towards the AMS2774 standard for Heat Treatment of wrought nickel alloy and cobalt alloy parts [171], which is currently recommended for thermal treatment of PBF nickel-based superalloys by the Standard Specification for Additive Manufacturing Nickel Alloy (UNS N07718) with Powder Bed Fusion [170]. The effects of HT on PBF specimen tensile properties have been widely studied [66,74,92,95,165, 183,186,249,251,266,283,285,289,311,339,418]. Two standard HTs for conventional materials were applied by Zhang et al. [165]. The first one consisted of two steps: a solution treatment $\left(980^{\circ} \mathrm{C}, 1 \mathrm{~h}\right.$ /air cooling $)+$ double aging $\left(720^{\circ} \mathrm{C}, 8 \mathrm{~h} /\right.$ furnace cooling at $55^{\circ} \mathrm{C} / \mathrm{h}$ to $620^{\circ} \mathrm{C}, 8 \mathrm{~h}$ /air cooling) and the second HT started by a homogenisation treatment $\left(1080{ }^{\circ} \mathrm{C}, 1.5 \mathrm{~h} /\right.$ air cooling $)$, followed by the same two steps as the first HT. Both methods contributed to the recrystallisation of dendritic structures and the precipitation of $\gamma^{\prime}$ and $\gamma^{\prime \prime}$ \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-29(1)} \end{center} As-built\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-29} Heat treated\\ Fig. 49. Electron Backscatter Diffraction results show the effects of scan strategies A, B and C (illustrated on the right-hand side of the figure) on specimens microstructures [267]. This illustrates that different scan strategies generate different microstructural features, which cannot be fully eliminated by heat treatment. The different scan strategies also affect the mechanical properties. The tensile testing results showed that scan strategy B always developed the highest Young's modulus among all the applied strategies. particles. As a result, both HTs increased material tensile strength. It was also observed that in the two-step HT, which is normally used for forged materials, undissolved Laves phases were still present. This illustrated that the three-step HT was better performing, even though it produced slightly lower tensile strength. On the contrary, in a study where LPBF specimens were reinforced with carbon nanotubes [266], the recrystallisation of fine microstructures caused by the HT slightly decreased the yield strength. That was not the case when an in-situ HT technique achieved by taking advantage of the powder bed preheating function in EBM machines - was used. In fact, the results showed that this in-situ HT led to much coarser microstructures, a larger number of intergranular cracks and poorer tensile strength when compared to conventional EBM specimens [183]. Additionally, Sames et al. investigated the impact of cooling rates in EBM [97]. The slowly cooled specimens possessed a much higher tensile strength respect to the fast-cooled equivalents, at the cost of lower elongation. A 3-step HT on LPBF Haynes 282 increased yield strength and ultimate tensile strength [186]. HIP is a commonly used technique to eliminate the micro-porosities and mitigating the impact of microstructural anisotropy. In the literature is usually applied with or compared to other HT processes [64,172, 173,179,250,260,269,284,287,297,418]. For example, Tomus et al. studied the tensile properties of LPBF Hastelloy X specimens in AB, HT, HIP and HIP + HT conditions [179]. The results showed that the HIP improved specimen relative density from $99.2 \%$ to $99.8 \%$ in AB condition to $99.9 \%$. Although it also led to carbides precipitating at the grain boundaries, the amount of carbides was believed to be too low to affect specimen mechanical properties. A study on HIP of IN738 showed that it improved the tensile strength of parts built vertically and horizontally but decreased the strength in the $45^{\circ}$ samples [418]. Furthermore, all post-treated specimens possessed similar tensile strengths, while being lower than the $\mathrm{AB}$. This was attributed to the rearrangement of dislocations during post-processing. Confirmation to that was found in another study by Kreitcberg et al., where HIP LPBF specimens possessed the lowest yield strength (Fig. 50), both at room and elevated temperatures, compared with other LPBF samples tested [269]. From these\\ results, it was concluded that higher porosity lead to higher tensile strength. In a study by Yao et al. [271], the AB and HT tensile test fracture surfaces of LPBF IN718 and TiC reinforced LPBF IN718 specimens were evaluated. All specimens presented dimpled fracture surfaces, indicating a ductile failure. It was also observed that the dimples in the HT specimens were larger and shallower, indicating lower ductility than the $A B$ specimen. The homogenized STA LPBF IN718 samples achieved higher strength than the AMS wrought specifications and good plasticity as well [328]. Although tensile strength is a significant mechanical property in many situations, compressive strength is also of importance. Typically, complex structures including lattices will undergo both tensile and compressive loading in operation. Therefore, AM lattice structures, which are designed for tailored mass, surface area, modulus and strength, are usually characterised with compression tests [216]. Strondl et al. compared the tensile and compressive yield strength of EBM IN718 specimens [96]. The results indicated that the tensile strength was higher than the compressive resistance, especially in HT conditions. Smith et al. demonstrated that LPBF IN718 specimens possessed slightly higher compressive yield stresses [65]. However, none of the studies explained the cause of these differences. The hot compression behaviour of IN718 specimens was studied by Mostafa et al. for LPBF components applications in the forging process [317]. The results indicated that the tensile behaviour is highly strain-rate dependent. Fig. 51 summarises the tensile strength properties for all of the studies reported to date. Each point represents the tensile property results for a sample from a paper. It should be noted that most samples had different processing and post-processing conditions, as well as different \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-30(2)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-30(1)} \end{center} (c) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-30} \end{center} (e) testing procedures. For example, some of the low laying points in the figure correspond to high temperature tensile testing by Popovich et al. [68]. Violin plots require a lot of data and the lesser number of studies available for EBM and Young's Modulus is the reason why they do not have plots. From Fig. 51b, it is clear that EBM nickel-based superalloys are much less investigated than for LPBF. The average tensile strengths of EBM IN718 and IN625 are also slightly below their LPBF counterparts. The research gaps, in terms of which materials are studied, are also apparent in Fig. 51. Fig. 51 was based on Table 9 in the Appendix section which lists the published data for the tensile properties of PBF manufactured nickelbased superalloy. \subsection*{4.2. Hardness properties} Hardness is a measure of a material resistance to localised permanent deformation such as small dents or scratches [420]. Hardness measurements are easy to perform, which is one of the reasons why they are one of the most frequently used mechanical property characterisation tests to provide an indicator to material condition. The processing parameters and HT can significantly influence LPBF specimens hardness [421]. As well as the building parameters, powder proprieties such as particle size distribution, flowability and rheology also affect specimen final hardness [121]. Supports were also shown to have a marginal effect on the local microstructure and hardness due to the low heat input in LPBF [380]. \subsection*{4.2.1. Build orientation and hardness measurements} Yen et al. showed that the build orientation affected hardness \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-30(4)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-30(5)} \end{center} (d) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-30(3)} \end{center} (f) Fig. 50. Tensile stress-strain diagrams obtained at $760{ }^{\circ} \mathbf{C}$ for different thermal treatments [269]. (a) Annealed-wrought. (b) As-built. (c) Stress relieved (SR). (d) Recrystallisation annealed (RA). (e) Solution treated (ST). (f) Hot isostatically pressed. Hot isostatically pressed Laser Powder Bed Fused specimens possessed the lowest yield strength, highlighting the effect of different thermal treatments on the resulting mechanical properties of powder bed fused specimens.\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-31(3)} (c) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-31(2)} \end{center} Fig. 51. Violin plots of research conducted in different nickel-based superalloys on Tensile properties. From 137 papers. (a) Tensile strength of Laser Powder Bed Fused alloys. (b) Tensile strength of Electron Beam Melted alloys. (c) Yield strength of Laser Powder Bed Fused alloys. (d) Elongation of Laser Powder Bed Fused alloys. The white dot is the median, horizontal lines are the mean, the height represents the range, the width is the amount of data, and the vertical lines in the middle are the interquartile range. The materials are arranged in order of how much data was available. These figures include the vast range of sample conditions used in works covered in this review and so direct comparison between studies should be undertaken with caution. This shows that Inconel 718 and Inconel 625 have been studied most comprehensively within the literature. The average tensile strengths of Electron Beam Melted Inconel 718 and Inconel 625 are also slightly below their Laser Powder Bed Fused counterparts. The violin plots were created using free to access matlab code [419]. properties directly [290]. Additionally, Chlebus et al. demonstrated that planes parallel to the build direction had a slightly higher hardness than perpendicular ones [95]. This was imputed to the interfaces overlapping between deposited layers, which tended to crack more easily, as well as weaker grain boundary strengthening planes perpendicular to the build direction. Strößner et al. and Murr et al. both obtained similar results [277,354] while Tomus et al. believed that the difference in hardness between the different planes was so small that it could be ignored [179]. Naturally there are differences in the experimental methods adopted here which are potentially significant in giving rise to these differences in results. Variations in hardness with respect to the build direction, indicated that the hardness is affected by specimens microstructural anisotropy. Indeed, Chauvet et al. demonstrated that the size of $\gamma^{\prime}$ phase decreased gradually along the build direction, due to the thermal history of LPBF [146]. This variation in $\gamma$ ' size was consistent with the measured hardness gradient (Fig. 52a). In the study by Wang et al., hardness was found to decrease with increasing columnar structure width [278].\\ 4.2.2. Energy density, scan strategy, scan speed and hardness properties When evaluating the comprehensive effects of a series of processing parameters, energy density needs to be considered. However, the influences of energy density on final properties are difficult to assess since many studies led to contradictory conclusions based on differing experimental methods. Experimental results, showed in Fig. 53a, by Rong at al. seem to collocate that linear energy density has an optimum in a range between $173 \mathrm{~J} / \mathrm{m}$ and $303 \mathrm{~J} / \mathrm{m}$ [343]. However, in another study (Fig. 53b) by Jia et al., a proportional relationship between the linear energy density (in a range between $180 \mathrm{~J} / \mathrm{m}$ and $330 \mathrm{~J} / \mathrm{m}$ ) and the hardness was observed [344]. The conclusion to that result was that a higher linear energy density resulted in a more homogenous microstructure, thus increasing the hardness [344]. Conversely, other studies found that the materials hardness decreased by increasing linear energy density (in a range between $400 \mathrm{~J} / \mathrm{m}$ and $1200 \mathrm{~J} / \mathrm{m}$ ) [67,281,345,346]. In particular, one study indicated that this was due to a coarsened microstructure [345], while others thought that the cause had to be identified in a finer microstructure and less brittle precipitates [67,281]. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-31(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-31} \end{center} (b) Fig. 52. The size of $\gamma^{\prime}$ phase decreased gradually along the build direction, due to the thermal history of Laser Powder Bed Fusion, which affected the hardness properties. (a) Diagram showing the distribution of $\gamma^{\prime}$ size and hardness along the build direction [146]. (b) Diagram showing the relation between hardness and porosity [353]. This shows that parameters like build height also affect mechanical properties. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-32(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-32} \end{center} (b)\\ Fig. 53. Effects of linear energy density on hardness. (a) By Rong et al. [344]. (b) By Jia et al. [344]. This shows contradicting evidence from the different authors with Rong et al. arguing that linear energy density has an optimum in a range between $173 \mathrm{~J} / \mathrm{m}$ and $303 \mathrm{~J} / \mathrm{m}$ to obtain the best hardness properties [343] and Jia et al. saying that there is a proportional relationship between the linear energy density (in a range between $180 \mathrm{~J} / \mathrm{m}$ and $330 \mathrm{~J} / \mathrm{m}$ ) and the hardness observed [344]. This illustrates some of the variability in findings between different studies, further highlighting repeatability challenges in powder bed fusion.\\ Furthermore, in a study by Yen et al. no direct relationship between the hardness and volumetric energy density was observed [290]. Yang et al. noticed that the melt pool mode, such as keyhole and conduction mode (Fig. 54), in the LPBF process also influenced specimens hardness [369]. Two different microstructures were found in the central and peripheral zone for keyhole mode specimens, while the conduction mode specimens were more uniform. The main features of different melt pool regions are listed in Fig. 54. The central zone in the keyhole mode, which mainly benefited from finer dendritic and $\gamma^{\prime} / \gamma^{\prime}$, precipitates, had a higher hardness ( $249 \mathrm{HV} / 359 \mathrm{HV}$ ) than the marginal mode (249 HV/321 HV) and conduction mode ( $260 \mathrm{HV} / 330 \mathrm{HV})$. Additionally, Gu et al. observed a direct relationship between grain morphology and hardness in the melt pool [368]. The hardness in different locations of the melt pool was measured (Fig. 55b and c) and the results were listed in Fig. 55d. They showed that the top surface of the melt pool was full of fine cellular dendrites and equiaxed grains with an average hardness of $387 \mathrm{HV}$ while the bottom of the melt pool was dominated by unidirectional columnar dendrites with an average hardness of $337 \mathrm{HV}$. The edge of the melt pool, instead, was characterized by the presence of multidirectional columnar dendrites and an average hardness of 340-350 HV. The centre of the melt pool had an average hardness slightly higher of about $363 \mathrm{HV}$. From this, it was concluded that the morphology of the grains in the melt pool, which is controlled by the temperature gradient and the cooling rate, defined the hardness. Lu et al. considered the effects of island scan strategies [286]. A smaller island size contributed to higher solidification rates, meaning that the increased residual heat effectively heat treated the solidified materials, leading to a higher hardness. However, another study found that higher values of hardness and compressive yield strength were obtained from the samples produced using Meander scanning strategies as opposed to an Island approach [148]. Choi et al. indicated that there was an optimum scan speed to produce the highest hardness [59]. The effect of scan speed was studied by Xia et al., the results indicated that higher scan speeds were generally correlated to higher hardness [65]. Indeed, higher scan speeds produced finer primary dendrites and more homogeneously dispersed granular carbides, which both contributed to higher deformation resistance. Furthermore, Choi et al. indicated that there was an optimum scan speed to attain high hardness [59]. However, at higher scan speeds, discontinuous laser tracks and non-fully melted powder were found to increase the porosity. On the other hand, lower scan speeds gave rise to excessive energy input and material vaporization, trapping vapours in the solidified structures. Higher porosity led to lower hardness as the pores would easily collapse when loaded. As proof, LPBF IN625 porosity was found to be consistent with the hardness gradient [59,353]. Rong et al. observed a similar trend but they justify the decrease in hardness at lower scan speeds with a coarsened microstructure [345]. In contrast, Karimi et al. found no direct relationship between EBM specimens hardness and porosity [367]. \subsection*{4.2.3. Optimal heat treatment for improved hardness} Post-processing techniques also have a significant influence on specimens' hardness. Zhang et al. heat treated LPBF IN718 specimens according to standards AMS 5662 and AMS 5383 [165]. Both methods increased the hardness than their $\mathrm{AB}$ equivalents due to the uniform precipitation of $\gamma^{\prime}$ and $\gamma^{\prime \prime}$ strengthening particles during ageing. Similar results and explanations were published in several other studies for LPBF [70,95,184,216,277,287,347,349,350,352] and EBM specimens [294, 295]. Deng et al. suggested that grain size growth might lower the hardness [294], while another study on HT LPBF Nimonic 263 samples by Vilaro et al. explained that the dislocation density and precipitation of $\gamma$ ' particles worked together to influence the hardness [285]. As such, grain size is not the only factor to influence hardness. A study reported that annealing LPBF IN718 at $600^{\circ} \mathrm{C}$ for $2 \mathrm{~h}$ did not affect the grain size or the microhardness of the sample [379]. Whereas under the solid solution process of $950{ }^{\circ} \mathrm{C}$, the fine homogeneous $\delta$-phase and $\gamma^{\prime}$ phase in the grain or near the grain boundary had an important effect on the X-Y surface hardness value (476-500 HV) of the alloy [374]. Under the solid solution process of $950{ }^{\circ} \mathrm{C}$, the fine homogeneous $\delta$-phase and $\gamma$ ' phase in the grain or near the grain boundary had an important effect on the surface hardness value (HV476-500) of LPBF IN718 [375]. Hence, some studies concluded that there was an optimum annealing temperature to obtain the highest hardness [275, 348]. Below the optimum temperature, residual stresses were relieved with no precipitation formation, leading to lower hardness. However, Deng et al. suggested that residual stress could harden the material to some extent [289]. Above the optimum temperature, $\delta$ particles were dissolved, leaving the sample without their strengthening mechanism, decreasing the hardness [348]. Additionally, Tucho et al. demonstrated that the effects of solution HT on the hardness of LPBF IN718 were dependent upon both temperature and hold time [185]. In order to achieve full recrystallisation, a solution temperature higher than $1100{ }^{\circ} \mathrm{C}$ was necessary. It was also found by Sun et al. that HT was not only able to improve EBM IN718 specimens hardness, but also remove any gradient of this along the build direction (Fig. 56) [308]. Other post-processing techniques were found to have positive effects on hardness. An in-situ HT technique applied on EBM IN718 specimens achieved much higher material hardness than an equivalent HT [183]. However, in a study on LPBF IN718 by Tillmann et al., it was observed that while specimen density and microstructure isotropy was highly improved by HIP, hardness decreased compared to the AB equivalents [72]. This was possibly due, at least in part, to grain growth. Murr et al. obtained similar results in their study on EBM of IN625 [292]. They noticed that HIP not only changed the crystallographic structure but also dissolved the $\gamma$ " precipitates, leading to a decrease in hardness. Another technique used different ion irradiation fluences and found that as the \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-33} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-33(4)} \end{center} (b)\\ Fig. 54. Illustration of the heterogeneity of powder bed fusion material properties. (a) Diagram of keyhole mode. (b) Diagram of conduction mode thin wall. (c) Summary of the features in keyhole mode and conduction mode thin walls. (d) Microhardness distribution in the keyhole mode. (e) Microhardness distribution in the conduction mode thin walls [369]. The central zone in the keyhole mode, which mainly benefited from finer dendritic and $\gamma^{\prime} / \gamma^{\prime \prime}$ precipitates, had a higher hardness (249 HV/359 HV) than the marginal mode (249 HV/321 HV) and conduction mode (260 HV/330 HV). \begin{center} \begin{tabular}{|c|c|c|c|c|} \hline \multicolumn{2}{|c|}{\multirow{2}{*}}{Feature} & \multicolumn{2}{|c|}{Keyhole mode} & \multirow{2}{*}{\includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-33(2)} } \\ \hline & & \begin{tabular}{l} Center \\ zone \\ \end{tabular} & \begin{tabular}{c} Marginal \\ zone \\ \end{tabular} & \\ \hline \multicolumn{2}{|c|}{\multirow{3}{*}}{}\{\begin{tabular}{c} PDAS $(\mu \mathrm{m})$ \\ Cooling rate $(\mathrm{K} / \mathrm{s})$ \\ Main precipitates \\ \end{tabular}\} & $0.3 \pm 0.1$ & $3.0 \pm 0.3$ & $2.6 \pm 0.3$ \\ \hline & & $\sim 10^{7}$ & $\sim 10^{4}$ & $\sim 10^{4}$ \\ \hline & & $\gamma^{\prime} / \gamma^{\prime \prime}$ & Laves & Laves \\ \hline \multirow{3}{*}{}\begin{tabular}{l} Area fraction of \\ main precipitates \\ \end{tabular} & Bottom & $22.7 \%$ & $7.6 \%$ & $8.9 \%$ \\ \hline & Middle & $18.5 \%$ & $11.8 \%$ & $10.3 \%$ \\ \hline & Top & $13.9 \%$ & $14.5 \%$ & $16.6 \%$ \\ \hline \multicolumn{2}{|l|}{Texture} & $\langle 001\rangle$ & - & - \\ \hline \end{tabular} \end{center} PDAS indicate the primary dendrite arm spacing (c) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-33(3)} \end{center} (d) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-33(1)} \end{center} (e) fluence was increased, so did the nano-hardness of IN625. Fig. 57 summarises the Hardness properties for all of the studies done to date. Each point represents a hardness value for a sample from a paper. It should be noted that most samples had different processing and post-processing conditions, as well as different testing procedures. The research gaps, in terms of which materials are studied, are also apparent in Fig. 57 since the materials are arranged in order of how much data was available. Data for EBM materials is not shown in this graph as there was not enough data reported to justify a similar type of figure. The reader is directed to Table 10 in the Appendix which compiles published data for the hardness of PBF manufactured nickel-based superalloy, including the EBM data. \subsection*{4.3. Toughness properties} In many structural applications, material fracture toughness, a measure of material resistance to unstable crack propagation, needs to be carefully evaluated. The delay period associated with inhibition of crack propagation is a further key requirement in safety critical applications. However, crack propagation is also evaluated in both creep and fatigue scenarios. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-34(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-34(4)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-34(2)} \end{center} (c) \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|c|c|c|c|} \hline (b) & 1 & 2 & 3 & 4 & 5 & 6 & 7 & 8 & 9 & 10 & 11 & 12 & 13 \\ \hline HV & \begin{tabular}{c} 354. \\ 92 \\ \end{tabular} & \begin{tabular}{c} 366. \\ 91 \\ \end{tabular} & \begin{tabular}{c} 371. \\ 95 \\ \end{tabular} & \begin{tabular}{c} 387. \\ 06 \\ \end{tabular} & \begin{tabular}{c} 379. \\ 03 \\ \end{tabular} & \begin{tabular}{c} 370. \\ 86 \\ \end{tabular} & \begin{tabular}{c} 361. \\ 87 \\ \end{tabular} & 350 & \begin{tabular}{c} 363. \\ 25 \\ \end{tabular} & N.A. & \begin{tabular}{c} 345. \\ 96 \\ \end{tabular} & \begin{tabular}{c} 347. \\ 02 \\ \end{tabular} & \begin{tabular}{c} 349. \\ 04 \\ \end{tabular} \\ \hline 14 & 15 & 16 & 17 & & (c) & 1 & 2 & 3 & 4 & 5 & 6 & 7 & 8 \\ \hline \begin{tabular}{c} 342. \\ 98 \\ \end{tabular} & \begin{tabular}{c} 342. \\ 02 \\ \end{tabular} & \begin{tabular}{c} 339. \\ 90 \\ \end{tabular} & \begin{tabular}{c} 337. \\ 06 \\ \end{tabular} & & HV & 347 & 361 & 385 & 369 & 358 & 341 & 350 & 363 \\ \hline \end{tabular} \end{center} (d) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-34(3)} \end{center} (e) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-34} \end{center} (f)\\ Fig. 55. The difference in hardness and grain morphology between the different regions of the melt pool [368]. (a) Diagram of a melt pool. (b) Cross-section of the melt pool. (c) Longitudinal-section of the melt pool. (d) Hardness measurement results. (e) Grain morphologies in different areas in cross-section. (f) Grain morphologies in different areas in longitudinal-section of the melt pool. The difference in hardness between the different regions concluded that the morphology of the grains in the melt pool, which is controlled by the temperature gradient and the cooling rate, defined the hardness. This also illustrates the heterogeneity of Laser Powder Bed Fused material, not only throughout the sample, but also at melt pool level. \subsection*{4.3.1. Powder layer and build orientation effects on toughness} Ardila et al. and Gruber et al. studied the effect of recycled powder on the fracture toughness of LPBF and EBM IN718 specimens [113,383]. For LPBF, the powder was recycled 14 times, with a Charpy test performed at 5 different junctures between these cycles. Results (Fig. 58) showed that the energy needed to fracture the LPBF IN718 specimens fluctuated around $10 \mathrm{~J}$, with no noticeable influence from the powder degree of recycling. For EBM, excessive oxidation of the recycled powder was identified as the cause for insufficient melting and weak bonding in the specimen, with a consequent reduction in toughness. Sufiiarov et al. investigated the influence of layer thickness on the fracture toughness of LPBF IN718 [274]. The results showed that specimens with a higher layer thickness possessed lower impact strength than thinner ones $\left(59.6 \mathrm{~J} / \mathrm{cm}^{2}\right.$ for $50 \mu \mathrm{m}$ and $83.8 \mathrm{~J} / \mathrm{cm}^{2}$ for $30 \mu \mathrm{m}$, respectively). This highlighted that an increased layer thickness, with more lack-of-fusion defects and cracks in some sections of the specimen, can greatly diminish toughness. In a study by Popovich et al., specimens built in the vertical direction showed a slightly higher impact toughness than the horizontal equivalents $\left(91.3 \pm 4.0 \mathrm{~J} / \mathrm{cm}^{2}\right.$ and $83.8 \pm 3.5 \mathrm{~J} / \mathrm{cm}^{2}$, respectively) [382].\\ Unlike Popovich et al., Hack et al. showed instead that the build direction had no evident impact on LPBF IN625 specimens toughness [381]. A common theme with comparing methodologies highlights the need for uniformity in sample preparation to allow more meaningful comparison between results. \subsection*{4.3.2. As-built specimens have better toughness than heat treated equivalents} Popovich et al. also evaluated the effects of a HT, consisting of homogenisation and ageing (HA), on the impact toughness of LPBF IN718 [68]. The impact toughness of AB specimens $\left(58.7-79.3 \mathrm{~J} / \mathrm{cm}^{2}\right.$ ) was about two times higher than commercial hot rolled and HT equivalents $\left(33-38 \mathrm{~J} / \mathrm{cm}^{2}\right)$. This was also observed by Hack et al., all specimens tested possessed superior impact and fracture toughness than their conventionally manufactured equivalents [381]. A possible explanation is that the specimens became more brittle after HT. Similarly to tensile testing results, the material impact strength presented clear signs of ductility in both $\mathrm{AB}$ and $\mathrm{HT}$ conditions as well as brittle fractures, especially where built-in defects were present [68]. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-35(2)} \end{center} Fig. 56. Hardness distribution in as-built and heat-treated Electron Beam Melted specimens [308]. This shows a reduction in hardness as cooling regimes change with build height. This figure illustrates that although build height affects hardness, heat treatment can reduce or even eliminate that anisotropy. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-35} \end{center} Fig. 57. Violin plots of research conducted in different nickel-based superalloys on Hardness properties, from 79 papers. The white dot is the median, horizontal lines are the mean, the height represents the range, the width is the amount of data, and the vertical lines in the middle are the interquartile range. The materials are arranged in order of how much data was available. These figures include the vast range of sample conditions used in works covered in this review. This shows that Inconel 718 and Inconel 625 have been studied most comprehensively within the literature. The violin plots were created from free-access matlab code [419]. \subsection*{4.4. Fatigue} Components used in critical applications, such as aerospace, are subjected to dynamic cyclic mechanical and thermo-mechanical loads, meaning that their fatigue properties must be carefully considered [420]. This presents a distinct set of performance characteristics as compared to other approaches simulating near static loading conditions. LPBF IN718 damage evolution during monotonic and cyclic loading was monitored and showed accelerated damage evolution in LPBF materials compared to forged [337] as shown in Fig. 59. LPBF process defects result in worse fatigue performance and deteriorate the fatigue crack \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-35(1)} \end{center} Fig. 58. Results from Charpy impact test by Ardila et al. [113]. These results highlight the consistency challenges associated with Powder Bed Fusion apparatus through material recycling procedures. growth behaviour [327]. This is a universally reported observation and much effort has been invested in better accommodating defects or engineering these out of the process. Fig. 60 shows the materials and PBF parameters investigated in Fatigue studies up to date. IN718 is by far the most studied material and the effect of thermal treatments, build orientation and surface condition are some of the most studied parameters. This figure also highlights materials and areas which would require more attention and study, such as the effect of powder composition and laser power on fatigue properties. \subsection*{4.4.1. Grain morphology and orientation affect fatigue performance} A variety of fatigue performance studies focus on the effects of different processing parameters. Among them, the build direction was heavily investigated [176,252,255,296,312,319,384,385,387,402]. For example, two studies compared the fatigue performance of specimens built in the $0^{\circ}, 45^{\circ}$ and $90^{\circ}$ directions [384,385]. Brodin et al. reported that LPBF Hastelloy X specimens built at $90^{\circ}$ possessed the best fatigue strength [384]. The same result was found for HIP EBM specimens (Fig. 61) [296] (see Fig. 62). Regarding fatigue life, LPBF Hastelloy X built at $0^{\circ}$ performed better than those built at $90^{\circ}$, when the load was higher than $600 \mathrm{MPa}$. Under that stress, it was found a negligible difference in performance [255]. Konečná et al. obtained similar results and reported that the large surface roughness of notched $90^{\circ}$ specimens may be the reason for their poor fatigue life [387]. Furthermore, for LPBF nickel-based superalloy $\mathrm{K} 536$, the fatigue performance anisotropy was not apparent at elevated temperatures (between $400{ }^{\circ} \mathrm{C}$ and $600^{\circ} \mathrm{C}$ ) [319]. However, in another study at low strain amplitudes, IN718 specimens built at $45^{\circ}$ possessed longer fatigue lifetimes than the $0^{\circ}$ samples [176]. A possible explanation is that the $45^{\circ}$ built specimens have a longer mean free path for dislocation movement. Grain structure also has an important impact on fatigue performance. Indeed, Zhou et al. and Kirka et al. carefully studied the effect of grain features on the fatigue performance of LPBF IN718 specimens [296, 396]. It was found that the build direction had a lesser influence on specimens with equiaxed grains, whereas specimens with columnar grains could withstand much lower cyclic stress amplitude. It was suggested that a columnar grains structure was more suitable for turbine blades - in order to provide directional preferential performance - while an equiaxed grains structure was more useful in turbine disks to comply with a distinct loading condition. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-36(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-36(2)} \end{center} (b) Fig. 59. A typical fatigue life curves for Laser Powder Bed Fused and forged alloys, demonstrated for Inconel 718 [337]. (a) Strain vs. number of cycles. (b) Stress vs. number of cycles. It is universally accepted that Laser Powder Bed Fusion process defects result in worse fatigue performance and deteriorate the fatigue crack growth behaviour, compared to conventionally manufactured materials. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-36} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-36(5)} \end{center} (b) Fig. 60. Pie charts highlighting the materials and Powder Bed Fusion parameters investigated in Fatigue studies up to date, from 44 papers. (a) Materials studied. (b) Parameters studied. Inconel 718 is clearly the most studied material and the effect of thermal treatments and build orientation on the fatigue performance of powder bed fused nickel-based superalloys have been studied the most. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-36(3)} \end{center} Fig. 61. S-N curves for Electron Beam Melted Inconel 718 of varying textures and orientations, equiaxed and columnar indicate specimens' grain structures that been controlled in the building process [296]. This shows that specimens built at $90^{\circ}$ possess the best fatigue strength. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-36(4)} \end{center} Fig. 62. Comparison of stress amplitude of specimens in different build directions [396]. This shows that the build direction had a lesser influence on specimens with equiaxed grains, whereas specimens with columnar grains could withstand much lower cyclic stress amplitude. \subsection*{4.4.2. Laser driven effects on fatigue performance} The effects of input energy on specimens fatigue resistance have also been investigated [388]. IN718 Specimens manufactured using an input energy of $250 \mathrm{~W}\left(59.5 \mathrm{~J} / \mathrm{mm}^{3}\right)$ possessed less porosity and a higher fatigue resistance, than the ones produced at $950 \mathrm{~W}\left(59.4 \mathrm{~J} / \mathrm{mm}^{3}\right)$. Interestingly, a functionally graded cylinder with a core processed at $950 \mathrm{~W}\left(59.4 \mathrm{~J} / \mathrm{mm}^{3}\right)$, leading to coarse elongated grains, and a shell processed at $250 \mathrm{~W}\left(59.5 \mathrm{~J} / \mathrm{mm}^{3}\right)$, resulting in fine equiaxed grains (Fig. 63). After HT, these cylinders showed the highest fatigue lifetime, which almost double that for HT $250 \mathrm{~W}\left(59.5 \mathrm{~J} / \mathrm{mm}^{3}\right)$ specimen. Kantzos et al. evaluated the influence of hatch spacing and of the corresponding cooling rates on fatigue performance [320]. Increasing the hatch spacing produced some lack of fusion defects, while decreasing the cooling rates resulted in a significant increase in pores size and number. The results highlighted that the fatigue performance was highly sensitive to the changes in porosity of the material. The influence of scan strategies' contour regions were also considered [321]. The lack of fusion defects in the contour and contour-hatch interface significantly limited specimen fatigue performance and could not be eliminated by thermal post-processing. However, it was shown that this contour region could be mechanically removed, leading to significant improvements in fatigue performance. This challenge is well appreciated in the conventional machining world where surface integrity has been widely studied. \subsection*{4.4.3. Heat treatments and machining strategy for fatigue performance} Post-processing can be used to reduce porosity and anisotropy in the material, in order to improve its fatigue performance. Wang et al. applied HIP to eliminate or reduce built-in defects, improving their fatigue limit from $500 \mathrm{MPa}$ for $\mathrm{AB}$ specimens to $550 \mathrm{MPa}$ [255]. In a study by Kanagarajah et al., a two-stage STA was applied to some LPBF IN939 specimens and found that the HT induced some brittleness in the material, dramatically decreasing specimens fatigue life [253]. In fact, it was found that higher brittleness was related to higher sensitivity to crack initiation and crack growth. A standard HT was instead applied to stress-relieved LPBF IN718 specimens. Fig. 64 shows the strain distribution overlaid with grain boundary maps for stress-relieved and HT specimens after fatigue testing [254]. In the first, strain was evenly distributed, with dislocation pileups at the grain boundaries (Fig. 64a). These coupled with finer grains, reduced stress concentration and crack initiation, were considered as the main fatigue-strengthening mechanisms. HT specimens, at the opposite, showed more localised strain (Fig. 64b). Despite these differences, they also showed excellent fatigue resistance, mainly due to big contribution of precipitation hardening, controlled by $\gamma^{\prime}$ and $\gamma$, [254]. The fatigue properties at $455^{\circ} \mathrm{C}$ of homogenized STA samples was studied [328]. Similarly, Popovich et al. compared the thermomechanical fatigue life of $\mathrm{AB}$ and two-stage HT LPBF IN718 specimens [388]. HT specimens showed a higher fatigue resistance due to the dissolution of the brittle Laves phase and the presence of $\delta$-phase on the grain boundaries. Balachandramurthi et al. compared the effects of STA and HIP + STA on EBM specimens fatigue performance [321]. The results showed that the HIP + STA yielded much better fatigue resistance than STA only.\\ This is because HIP + STA closed most of the built-in defects and completely dissolved the $\delta$-phase, which can hinder the precipitation of $\gamma$ " phases. In another study, LPBF specimens possessed better fatigue performance than those produced by EBM, due to the more numerous lack of fusion defects introduced by EBM [397]. However, it was found that HIP and HT were able to effectively close the built-in defects in both LPBF and EBM specimens. Surface finish is another important factor that influences the fatigue performance. In fact, for both LBPF and EBM, the AB surfaces provided more fracture initiation sites than their machined equivalents [397]. Furthermore, it was shown that small-scale specimens show a higher number of surface and near-surface defects which result in reduced mechanical properties, including a $65 \%$ reduction in fatigue strength in the case of Kotzem et al. [422]. It is highly likely that components produced by PBF will require machining ahead of experiencing fatigue-based loading conditions. Surface roughness and built-in defects, such as embedded particles, were identified as the main cause of fatigue initiation points (Fig. 65), limiting the fatigue performance [252,384]. Indeed, Wan et al. observed a $\sim 50 \%$ increase in fatigue strength after surface machining and polishing [423]. Koutiri et al. observed that the use of lower scan speeds and lower power led to higher surface roughness, particularly on the down-skin sides for large building angles [386, 397]. Other polishing processes such as low-stress grinding, have instead been found valid in producing more neat surfaces, thus improving specimens fatigue life [386,397]. Witkin et al. discovered that the AB surfaces of notched fatigue specimen always contained critical-sized defects which may lead to faster fatigue failure than what estimated using notch stress concentration calculations [392]. A modified HT was developed in order to reduce LPBF IN718 anisotropy in fatigue performance [401]. The effects on fatigue performance of surface preparation technique (mechanical or electromechanical polishing) was investigated for LPBF IN718 [402]. The effect of dry or emulsion cutting conditions on the fatigue performance of LPBF IN718 were also investigated and using a dry machining condition resulted in better surface roughness and more compressive residual stress, leading to more cycles to failure [373]. \subsection*{4.4.4. Fatigue crack growth} Konečná et al. and Kim et al. compared the fatigue crack growth rate of conventional (such as rolled or forged material) and LPBF IN718 specimens (Fig. 66) [390]. The LPBF manufactured specimens and their conventional equivalents had similar crack growth resistance in the high $\Delta \mathrm{K}$ region [390,424], whereas in the intermediate $\Delta \mathrm{K}$ region, the LPBF samples had much higher fatigue crack growth rate than their conventionally built counterparts [424]. However, the threshold stress intensity factor was about $1 \times 10^{-7} \mathrm{~mm} /$ cycle for LPBF samples, much lower than for conventional equivalents, highlighting the poor fatigue resistance of LPBF manufactured materials. Three main factors were found to be responsible: the lower boron content, the finer microstructure and the presence of residual stress in the LPBF manufactured materials. Poulin et al. investigated the influence of build direction and postprocessing on the crack propagation behaviour of LPBF IN625 \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-37} \end{center} Fig. 63. Diagram of a functionally graded structure with fine equiaxed grains at the core encased by columnar grains. This showcases the degree of grain engineering possible by Powder Bed Fusion to obtain desired mechanical properties [388]. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-38(4)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-38(2)} \end{center} (b)\\ 0.02 Fig. 64. The strain distribution overlaid with grain boundary maps of specimens after fatigue testing [254]. (a) Stress relieved specimens. (b) Heat treated specimens. Arrows indicate the loading direction. In the stress-relieved state, strain was evenly distributed, with dislocation pileups at the grain boundaries which reduced stress concentration and crack initiation and were considered as the main fatigue-strengthening mechanisms. The heat treated specimens showed more localised strain but showed excellent fatigue resistance, demonstrating the beneficial effects of heat treatment on Fatigue performance of Powder Bed Fused specimens. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-38(3)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-38} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-38(1)} \end{center} (c) Fig. 65. Scanning Electron Microscopy image of the subsurface defect and Fatigue initiation sites in as-built Laser powder bed fused Inconel 625 [252]. (a) Low magnification. (b) High magnification. (c) Embedded particle on the as-built surface. Surface roughness and built-in defects, such as embedded particles, were identified as the main cause of fatigue initiation points. specimens [312]. The fatigue crack growth results (Fig. 67) indicated that the crack growth rate of stress-relieved LPBF specimens was highly dependent on the build direction. Similar results by Ma et al. showed the dependence of crack growth rate on build direction at room and elevated temperatures [425]. Furthermore, in a study by Brynk et al., LPBF IN718 specimens built at $45^{\circ}$ were found to have the best fatigue crack growth resistance [385]. Post-processing HIP was also successfully used to eliminate the anisotropic behaviour. Finally, the fatigue crack growth rate of LPBF IN625 specimens seemed comparable to the wrought equivalents. Additionally, the effect of HT and loading direction on dwell-fatigue crack propagation resistance of LPBF IN718 at $550{ }^{\circ} \mathrm{C}$ and $2160 \mathrm{~s}$ dwell holding period was investigated and showed that a creep mechanism was dominant [400,426]. The fatigue crack growth rate of LPBF IN625 was found to increase as the stress ratio increases [399]. Horizontal samples had a faster fatigue crack growth rate than vertical samples [399]. LPBF IN738 was built using different laser powers and scan speeds and it was found that small grains present along large grain boundaries act as crack initiation sites or affect crack propagation path [427]. \subsection*{4.5. Creep properties} Creep resistance defines component performance in hightemperature conditions. However, the creep performance of LPBF nickel-based superalloys has not been fully studied yet. Studies have focused mainly on IN718 (Fig. 68a) and the effect of thermal treatments and build orientation (Fig. 68b) on the creep properties. Fig. 68 also highlights the areas which require more research, such as the effect of surface condition and complex geometries on the creep properties of PBF nickel-based superalloys. Investigations to date also report significant shortfalls with respect to counterparts machined from wrought material. Creep testing can be categorized into two types: tensile loading tests [75,249,251,384,388,405-407] and compressive loading tests [190, 275,295]. The former can normally provide information on creep fracture. For example, Brodin et al. observed the creep fracture surface of LPBF Hastelloy $\mathrm{X}$ tested at $815{ }^{\circ} \mathrm{C}$ with a tensile load [384]. Compressive creep tests are performed to analyse specimens creep rate. The tests were usually interrupted when a predetermined plastic deformation was achieved. Some non-standard creep specimens, such as 2 bar specimens are also used [75]. Small punch creep specimens are another commonly used specimen geometry, these tests are shown schematically in Fig. 69. Wang et al. used a small punch creep test to compare the creep performance of forged, cast and LPBF material [411]. The results (Fig. 70) indicated that the AB LPBF specimens possessed similar creep lifetime to the forged ones, but far lower than the cast specimen. Laves phase at the grain boundaries may have led to poor creep performance in LPBF specimens. Moreover, the local tensile stress induced by the experimental loading in the fracture region was believed to be another reason for the reduced time to rupture. Another study also found that \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-39} \end{center} Fig. 66. Diagram showing the differences in fatigue crack growth between Laser Powder Bed Fused and conventionally manufactured Inconel 718 specimens [390]. The Laser Powder Bed Fused specimens and their conventional equivalents had similar crack growth resistance in the high $\Delta \mathrm{K}$ region [390,424], whereas in the intermediate $\Delta \mathrm{K}$ region, the Laser Powder Bed Fused samples had much higher fatigue crack growth rate than their conventionally built counterparts. This highlights that Powder Bed Fusion has the potential to have similar mechanical performance to conventional methods. LPBF IN718 had a lower creep ductility than their wrought equivalent due to oxide contamination from powder surface [428]. Moreover, new AM 'specific' approaches for creep testing inspection have been investigated. Xu et al. conducted a new staged thermalmechanical testing method which looked at the defect evolution in the LPBF IN718 specimens during creep [75]. The idea was to interrupt the testing and characterise the defects using X-ray computed tomography. For example, Fig. 14 shows the porosity distribution in the specimens at the different stages where the test was stopped. This allowed the characterisation of the porosity accumulation during creep and the localisation of the position of the weakest point. \subsection*{4.5.1. $L P B F$ process parameters affect creep performance and failure} The effect of LPBF building orientation on specimens creep performance has been widely studied [249,251,384,405,406]. Rickenbacher et al. considered the creep performance of horizontally and vertically built LPBF IN738LC specimens [249]. The results indicated that the vertically built specimens had better creep resistance, with respect to cast equivalents. Hautfenne et al., Kuo et al. and Kunze et al. obtained similar results, explaining that vertical samples had the stress applied parallel to the columnar elongated grains, behaving similarly to the creep resistance strengthening mechanisms in directionally solidified and single crystal superalloys [251,405,406,429]. Small Punch Creep test of LPBF CM247LC with different process parameter (beam power, layer thickness and energy density) and build orientations ( $30^{\circ}$ and $90^{\circ}$ ) found that $90^{\circ}$ samples performed better for creep deformation [412]. Sanchez et al. also noted that the build orientation and the stress state were responsible for the different types of failure modes of creep speciemens [414].Shassere et al. also studied the influence of microstructure on creep performance of EBM IN718 specimens [409]. The results indicated that specimens with columnar grains possessed better creep performance than the ones with equiaxed grains, especially when the loading direction was parallel to the columnar grain growth direction. In fact, having the grain boundaries transverse to the loading direction was found to be particularly detrimental on creep resistance. Laser power is also an important factor which can affect creep performance. Popovich et al. investigated the creep performance of LPBF IN718 specimens manufactured with two different laser power, namely $250 \mathrm{~W}\left(59.5 \mathrm{~J} / \mathrm{mm}^{3}\right)$ and $950 \mathrm{~W}\left(59.4 \mathrm{~J} / \mathrm{mm}^{3}\right)$. In general, ductile fracture (Fig. 71a) dominated in specimens produced with lower power, except in some regions rich in brittle precipitates (Fig. 71b) [388]. On the other hand, specimens manufactured using higher power failed before reaching the required testing conditions. This was mainly due to the presence of a large number of built-in defects (Fig. $71 \mathrm{c}$ and d) caused by the excessive power. Sanchez et al. also showed that using multiple-laser scan strategies does not adversely affect the creep performance of LPBF IN718, with multi-laser vertically built samples even performing similarly to wrought material [415]. Assessing functional performance of functionally graded materials produced by AM is an interesting emergent research area. Popovich et al. analysed and compared the performance of post-treated LPBF IN718 specimens and functionally graded specimens [388]. \subsection*{4.5.2. Heat treatment for enhanced creep performance} To improve LPBF specimens creep performance, post-processing is still necessary at the present stage. Multiple studies compared the effects of HT on LPBF IN718 specimens [190,275,406,408,413]. Pröbstle et al. characterised the creep performance of cylindrical specimens subjected to different HTs, including direct ageing and STA [190]. The post-treated specimens showed an improved creep performance. As a confirmation, Hautfenne et al. [405] proved that the use of a solution temperature higher than $1000{ }^{\circ} \mathrm{C}$ could contribute to better creep performance. Using a solution temperature of $1000{ }^{\circ} \mathrm{C}$ followed by a two-stage ageing treatment resulted in better creep resistance than specimens directly aged [190,275]. However, when the solution temperature was decreased below $1000{ }^{\circ} \mathrm{C}$, direct aged specimens performed better [190,406]. Two possible reasons were identified to explain these results [190]. The first might be related to the strengthening phases composition. In nickel-based superalloys, Nb content is critical since it forms the main strengthening precipitates. When solution treated at $1000{ }^{\circ} \mathrm{C}, \delta$-phased dissolves, releasing $\mathrm{Nb}$ in the surrounding matrix and hence allowing more $\gamma$ " precipitation. On the contrary, when treated at $930^{\circ} \mathrm{C}$, more $\delta$-phase is formed at the expense of $\gamma^{\prime \prime}$ phase. But since these precipitates contribute more to creep resistance than the $\delta$ particles, specimens treated at $1000{ }^{\circ} \mathrm{C}$ perform better. The second possible reason might have been the size of $\gamma$, (Fig. 72), with the largest average size of these precipitates was found in specimens treated at $1000{ }^{\circ} \mathrm{C}(13.4 \pm 5.8 \mathrm{~nm})$, followed by direct aged ones $(9.4 \pm 3.2 \mathrm{~nm})$ and lastly the $930{ }^{\circ} \mathrm{C}$ treated ones $(9.1 \pm 5.8 \mathrm{~nm})$ [190]. Wang et al. however, noted that using STA resulted in a shorter creep life than using homogenisation ad ageing treatments [430]. Another study found that creep life and ductility was improved after HT by adding Y (yttrium) as Y-oxide (yttria) precipitated around the $\delta$-phase instead of Al-oxides, which impeded $\delta$-phase precipitate growth and improved creep properties [413].(Table 11) Similarly, Davies et al. tried two HT strategies with different solution temperatures $\left(1150{ }^{\circ} \mathrm{C}\right.$ and $1275^{\circ} \mathrm{C}$ ) to improve the creep performance of LPBF C263 and found that the higher solution temperature increased creep resistance [410]. This was because a higher solution temperature generated a more equiaxed microstructure, smaller average local misorientation, shorter random grain boundary network segment length and carbides ( $\mathrm{MC}$ and $\mathrm{M}_{6} \mathrm{C}$ ) precipitation at grain boundaries. It was \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-40(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-40(4)} \end{center} (c) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-40(2)} \end{center} (e) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-40(6)} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-40} \end{center} Stress intensity factor range, $\Delta \mathrm{K}\left(\mathrm{MPa}^{*} \mathrm{~m}^{1 / 2}\right), \mathrm{R}=0.1$ (d) Fig. 67. Fatigue crack propagation diagrams of specimens in various heat treated conditions and build orientations [312]. They indicate that the crack growth rate of stress-relieved Laser Powder Bed Fused specimens was highly dependent on the build direction. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-40(5)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-40(3)} \end{center} (b) Fig. 68. Pie chart of the materials and the powder bed fusion parameters investigated in the creep studies to date, from 21 papers. (a) The materials studied. (b) The parameters studied (the label "Basic" refers to studies which only tested the material in one condition, focusing on performance without looking into the effect of a certain parameter.) Inconel 718 is the main material investigated and the effect of thermal treatments and build direction have been studied the most. concluded that shorter random grain boundary network segment length, meant shorter potential intergranular crack paths. In turn carbides could hinder grain boundaries deformation, futher improving creep resitance. However, despite HT increasing creep life of LPBF IN718 samples, Sanchez et al. scan lines were apparent on the fracture surface of vertically built HT samples (Fig. 73b-b',c-c'), showing that despite post-processing, an AM specific failure still occurred [414]. In summary, effort must be directed to understand and optimise postprocessing to improve creep performance of LPBF manufactured nickelbased superalloys. Published data for the creep testing of LPBF manufactured nickel-based superalloys are compiled in Table 12 for the convenience of the reader. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-41(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-41} \end{center} (b) Fig. 69. Schematic showing tests on two non-standard specimens. (a) Small Punch test. (b) Two Bar Specimen test. This illustrates the potential of small and non-conventional specimen testing in Powder Bed Fusion. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-41(2)} \end{center} Fig. 70. Creep deflection vs creep time results for specimens in different conditions (Forged-N: loading direction parallel to the forging direction; Forged-P: loading direction perpendicular to the forging direction) [411]. Indicating that the as-built laser powder bed fused specimens possessed similar creep lifetime to the forged ones, but far lower than the cast specimen. \section*{5. Toward more appropriate testing procedures for AM} Conventional approaches to mechanical testing of AM components are costly and can serve to undermine the use case for AM. As such alternate approaches are required. Small specimen testing techniques have the capability to characterise a localised mechanical response while using only a small volume of materials [429]. For example, a small specimen testing method called "small punch test" exhibits potential in characterising LPBF nickel-based superalloys mechanical performance [430]. In fact, this method has been accepted as an approach to estimate mechanical properties from small quantities of materials only. The test is performed by deforming a disc specimen, typically $8 \mathrm{~mm}$ in diameter and $0.5 \mathrm{~mm}$ thick, under a hemispherical punch with a radius of 1-1.25 $\mathrm{mm}$ [431]. Two test configurations are available: constant load, which shows a creep-type response and is termed small punch creep test (SPCT); and constant displacement rate, which exhibits a tensile-type behaviour and is called small punch tensile test (SPTT) [432]. The specimen will deform biaxially and the data obtained can either be plotted on a displacement versus fracture time chart for SPCT [433] or load versus displacement chart for SPTT [434]. Both data sets can be correlated to the equivalent bulk mechanical properties through various means [419-422]. The small punch test method has been widely applied to evaluate the mechanical performance of various aerospace nickel-based superalloys [423-428]. Since only a small volume of material is required, this test offers a feasible option to study AM components mechanical properties when traditional testing methods are not possible, due to complex design geometries [429-431]. The anisotropic nature of PBF components has been highlighted on numerous occasions in this review. This is a major concern for the structural integrity of these materials. AB LPBF components usually exhibit microstructural and hence mechanical discrepancies between the build direction and the scanned planar direction. Even though various strategies have been employed to minimise this tendency, it is inevitable to have anisotropic LPBF materials, mainly because of the layer-upon-layer approach. The community is exploring scanning and building strategies to ameliorate anisotropy however this will require maturation of process technology and control in unison with the development of our materials understanding. It is generally observed that the material is mechanically weaker in the build direction ( $\mathrm{Z}$ direction) as compared to in the $\mathrm{X}-\mathrm{Y}$ plane. This is thought to be caused by poor bonding at the interface between individual layers. These anisotropies can also be evaluated through small specimen tests [431,432]. Small specimens can be sampled from different planes, allowing a localised investigation of the anisotropic behaviour of LPBF nickel-based superalloys. Using this technique, rapid adjustments on the process parameters can be done to minimise this tendency. This may allow researchers to have a better understanding of worst case scenarios. However, there are applications where anisotropy is sought [433]. For example, one AM sector that can greatly benefit from LPBF components anisotropy is lattice structure design [434,435]. These structures allows for lightweight design and good mechanical properties in specific load conditions. PBF capability in manufacturing lattice structures outweighs traditional manufacturing processes. To date, studies concerning the effectiveness of lattice structure on bulk mechanical properties are few and future work in this area would benefit the AM community. In addition more complex testing configurations (e.g. triaxial approaches) are required in order to validate component performance in these cases. Using more complex designs and lattice structures is quite clearly a research trend in AM [436]. For metal specific approaches a review of design and mechanical properties of metal lattice structures was written by Hanks et al. [437]. Here it is critical to ensure that methods are devised which can appropriately assess both the structural and material properties of AM components. A shortfall of many contributions made within the AM literature. \section*{6. Conclusions} Regardless of the unique advantages PBF boasts compared to more traditional manufacturing processes it is widely recognized within the literature that morphological defects and suboptimal microstructures limit the performance of Nickel-based superalloys in current class PBF approaches. This is not confined to Nickel-based superalloys alone and remains a broader materials problem. However, the high sensitivity of $\mathrm{Ni}$ alloy performance to precipitate, phase, texture and grain size makes these systems particularly challenging when compared to Ti or Fe based alloys. Further, the primary application areas (including within turbomachinery) mean that manufacturers require enhanced surety of the performance of these materials. This is particularly relevant for dynamic components subject to both thermal and mechanical loading cycles. As such many of the defects characteristic of PBF processes cannot be tolerated.\\ \includegraphics[max width=\textwidth, center]{2024_04_13_99ab48aa81fd824b9ddbg-42(3)} Fig. 71. Creep fracture surfaces of Laser Powder Bed Fused Inconel 718 specimen built with varying laser power [388]. (a),(b) Laser power of $250 \mathrm{~W}$ (59.5 $\mathrm{J} / \mathrm{mm}^{3}$ ). (c),(d) Laser power of $950 \mathrm{~W}\left(59.4 \mathrm{~J} / \mathrm{mm}^{3}\right)$. This highlights the catastrophic failure associated with lack of fusion. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-42(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-42} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-42(2)} \end{center} (c) Fig. 72. Transmission Electron Microscopy images showing the $\gamma$ " phase after different thermal treatments [190]. (a) Direct aged (DA). (b) Solution treated at $1000{ }^{\circ} \mathrm{C}$ then aged (SHT1000 ${ }^{\circ} \mathrm{C}$ ). (c) Solution treated at $930{ }^{\circ} \mathrm{C}$ then aged $\left(\mathrm{SHT} 930^{\circ} \mathrm{C}\right.$ ) conditions. Using a solution temperature of $1000^{\circ} \mathrm{C}$ followed by a two-stage ageing treatment resulted in better creep resistance than specimens directly aged, but when the solution temperature was decreased below $1000{ }^{\circ} \mathrm{C}$, direct aged specimens performed better. From the literature there is a desire to understand the influence of PBF process parameters - such as power density, scanning strategy and build direction - on specimen final mechanical properties. However, there is little evidence that product direct from machine will remotely match the performance of 'machined from wrought' equivalents. Hence there has been a significant body of work evaluating post-processing strategies which includes both thermomechanical techniques (to recover microstructural and in-built defects) and surface processing techniques to address stress concentration issues. This comes in the context of a rapidly developing machine tool market for AM which is seemingly improving month-to-month. It is also evident that the extensive campaigns for mechanical evaluation of PBF nickel-based superalloys through traditional test methods is time-consuming and costly. Furthermore, PBF specimens often exhibit evolving microstructures and properties throughout the building process, while most of the traditional testing standards were designed for monolithic materials. Hence, these testing methods might not reflect the localised mechanical discrepancies in LPBF materials to equip designers with appropriate information. In some regards, particularly pre-HT, AM components are best considered as a continuous fabrication (a single weld constituting the whole component) as opposed to a monolithic and uniform component. They should therefore be analysed as such. This review has highlighted a suite of Ni based materials which are \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-43(4)} \end{center} (a") \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-43(6)} \end{center} (d') \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-43} \end{center} (d")\\ 90 ${ }^{\circ}$ Meander HT \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-43(2)} \end{center} (b') \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-43(3)} \end{center} (b") \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-43(5)} \end{center} (e') \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-43(7)} \end{center} LD \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-43(1)} \end{center} Fig. 73. Fracture surfaces of Laser Powder Bed Fused Inconel 718 after creep testing [414]. (a) - (a') $90^{\circ}$ Meander as-built; Heat treated (b) - (b') $90^{\circ}$ Meander; (c) - (c') $90^{\circ}$ Stripe; (d) - (d') 45 Stripe; (e) - (e") $0^{\circ}$ Stripe and (f) - (f') Wrought Inconel 718 specimens. All images oriented in line with the Loading Direction (LD). All surfaces showed signs of ductile fracture and apparent scan lines on the heat treated $90^{\circ}$ samples, which indicates that heat treatments are still not optimised for Powder Bed Fused materials.\\ explored in the literature. Many of these can be considered the 'low hanging fruit' in that they readily consolidate in PBF to realise components. However, the pallet of materials processable by AM is continually developing and much of this resides in proprietary knowledge and is not committed to the academic literature. The wider exploitation of materials within AM will require the development of several approaches to accelerate this. Indeed the opportunities for materials development for PBF platforms may be inferred throughout this review. Machine tool technology is also proving to be a limiting factor. Process control strategies commonly observed in even low cost conventional machine tools are not yet available for AM systems. As such the immaturity of PBF systems is apparent from many of the studies reported. The current state-of-the-art does indeed deliver parts but the repeatability and reliability of these is still very much a work in progress. The shortfalls of current platforms often make the use of PBF a tenuous decision. Similarly, to materials development this paper has highlighted numerous opportunities for machine tool innovation which will enhance the utility of nickel-based materials in AM. \section*{7. Research outlook} Based upon our assessment of the state-of-the-art, the authors propose the following key themes (summarised in Fig. 74) as being of primary interest to proliferate the manufacture of $\mathrm{Ni}$ superalloy components using PBF: \begin{itemize} \item Heat treatments within-process - post-processing of any description with reduce the already tenuous business case for deploying AM. As such there is a real need to arrive at desired integrity and microstructure within the process. This will be derived from enhanced process understanding and control. Where 'post-processing' cannot be achieved in process, which will be the case for some time, we must consider efficient methods for installing the properties and performance required by specification. These may not always be consistent with methods for processing wrought equivalents but will allow designers to make better and wider use of PBF Ni alloys. \item Enhanced thermal management - Advanced scan strategies, 'new' energy beam profiles and in-process monitoring have the potential to overcome cracking and stochastic pore formation in difficult to process materials. They may also allow users to introduce microstructure by design. These approaches present additional challenges as they are so rich in data generation that the role of the computer scientist will be important in gathering and interpreting this. This is analogous to process control which is far more mature in more established manufacturing technologies. \item Modelling - Allied to the development which will emerge in process control, it is essential that we develop higher fidelity but computationally lighter modelling approaches. To relate process to part performance there will be a significant need to predict recrystallisation phenomenon over longer ranges which deal with part geometry at scale. In addition, machine tools must respond to complex events in process and as such 'on-the-fly' modelling approaches will be required to develop truly adaptive processes. \item Design - Exploitation of PBF (and all forms of AM for that matter) hinges upon designing for these techniques. It is foolish to expect that PBF will simply replace the preferred method of manufacturing for all components. Typically, such components will incorporate many processes whose effect upon tolerancing and microstructure are well \end{itemize} \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-44} \end{center} Fig. 74. Graphical representation of future research trends required to develop the next generation Powder Bed Fused nickel-based superalloys. Here an exemplar gas turbine blade is used to illustrate a vision for convergence of several new manufacturing capabilities.\\ understood. Our understanding of PBF in this regard is somewhat more juvenile and a holistic approach for PBF introduction is required. \begin{itemize} \item Build environment - It is apparent throughout the work presented here that build environment (pressure and species) drives the consolidation phenomena amongst other factors. In the opinion of the authors this feature of modern LPBF (this is less of an issue in EBM systems or course) is primed for further development and will be a feature of machine tools capable of achieving superior material properties in Nickel-based superalloys. \item Standardisation - Through groups such as the ASTM F42 committee, significant contributions have been made in standardisation descriptors and taxonomy. Much work remains however to allow direct comparison between machine technologies and resulting product. The reader will encounter this difficulty when assessing the extended appendices, the authors have compiled. It proves most difficult to draw meaningful conclusions without comparing likewith-like. \item Alloy formulation - This review has highlighted much work which attempts to process powdered specimens of existing and widely used Nickel-based superalloys. However, there is a tremendous opportunity within the Ni superalloy metallurgical fraternity to modify alloys specifically for AM processing. There is also an opportunity in this regard to devise new methods to enhance the longevity of $\mathrm{Ni}$ based materials through processing and devise reuse strategies in which alloy performance can be assured. The ability to reuse/recycle metal powder for PBF will therefore be hugely important in driving down process costs. \item The role of AI - Artificial intelligence as a discipline, while not new, is proving to be useful in materials and process design for AM. Through sophisticated decision making from suboptimal data sets it is possible to advance process and material development at pace. The authors propose that the role of AI will become more prevalent in AM given the experimental space which emerges when seeking to optimise composition and process parameter sets. Clearly automation and evaluation will be critical in driving this area when coupled to both experimental and modelling efforts. \end{itemize} The authors have had the privilege of exploring the AM literature from first efforts with powder bed fusion through to the state-of-the-art. It is fitting that Nickel-based superalloys were explored early in the development of AM but presented significant challenges as compared to the more readily processible Titanium and Iron based alloys which 'weld' well under PBF conditions. Many of the metallurgical challenges presented by nickel-based alloys persist and can trace their heritage back to the conception of these alloys as a family. As such we may suggest that AM practitioners continue to try to find new solutions to these established problems. The rate of literature contributions to this domain is clearly increasing as researcher access to PBF techniques expands. As such, this review will age much faster than a nickel-based superalloy in service but the authors would hope that this contribution is of some value to our community providing a reference to common challenges and baseline performance while inspiring the pursuit of new research directions. We trust the work is of value and would welcome scientific dialogue on all topics contained here. \section*{Declaration of competing interest} The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. \section*{Appendix A. Composition of IN718 and IN625} Table 8 below summarises the chemical composition of the most used nickel-based superalloys in PBF research: IN718 and IN625. Table 8 Summary of the chemical composition ( $w t \%$ ) of the most used nickelbased superalloys in PBF (from CES EduPack software, Granta Design Limited, Cambridge, UK, 2009). \begin{center} \begin{tabular}{lll} \hline Elements (wt\%) & IN718 & IN625 \\ \hline $\mathbf{N i}$ & $50-55$ & $58-69$ \\ $\mathbf{C r}$ & $17-21$ & $20-23$ \\ $\mathbf{F e}$ & $11-25$ & $<5$ \\ Mo & $2.8-3.3$ & $8-10$ \\ $\mathbf{N b}$ & $2.4-2.8$ & $3-4$ \\ $\mathbf{T a}$ & $2.4-2.8$ & 0.15 \\ $\mathbf{T i}$ & $0.65-1.2$ & $<0.4$ \\ $\mathbf{A l}$ & $0.2-0.8$ & $<0.4$ \\ $\mathbf{C o}$ & $<0.1$ & $<1$ \\ $\mathbf{M n}$ & $<0.35$ & $<0.5$ \\ $\mathbf{S i}$ & $<0.35$ & $<0.5$ \\ $\mathbf{C u}$ & $<0.3$ & $<0.1$ \\ $\mathbf{C}$ & $<0.08$ & $<0.015$ \\ $\mathbf{P}$ & $<0.015$ & $<0.015$ \\ $\mathbf{S}$ & $<0.015$ & \\ $\mathbf{B}$ & $<0.006$ & \\ \end{tabular} \end{center} Below are a summary of the mechanical properties reported in the literature. Appendix B. Tensile properties of PBF Nickel-based superalloys Table 9 Summary of tensile properties of PBF Nickel-based superalloys (' $\sim$ ' indicates that the data was obtained from the bar chart). \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample Condition & Yield strength (MPa) & \begin{tabular}{l} Tensile strength \\ (MPa) \\ \end{tabular} & Elongation (\%) & \begin{tabular}{l} Young's modulus \\ (GPa) \\ \end{tabular} & \begin{tabular}{l} Ref \\ \end{tabular} \\ \hline \multirow[t]{8}{*}{LPBF} & C263 & & Small punch tensile test & $90^{\circ} / \mathrm{HT} 1 / 20^{\circ} \mathrm{C}$ & 818 & 1100 & & & [273] \\ \hline & & & & $90^{\circ} / \mathrm{HT} 1 / 780^{\circ} \mathrm{C}$ & 401 & 409 & & & \\ \hline & & & & $0^{\circ} / \mathrm{HT} 1 / 20^{\circ} \mathrm{C}$ & 870 & 1045 & & & \\ \hline & & & & $0^{\circ} / \mathrm{HT} 1 / 780^{\circ} \mathrm{C}$ & 513 & 480 & & & \\ \hline & & & & $90^{\circ} / \mathrm{HT} 2 / 20^{\circ} \mathrm{C}$ & 843 & 886 & & & \\ \hline & & & & $90^{\circ} / \mathrm{HT} 2 / 780^{\circ} \mathrm{C}$ & 489 & 589 & & & \\ \hline & & & & $0^{\circ} / \mathrm{HT} 2 / 20^{\circ} \mathrm{C}$ & 590 & 1078 & & & \\ \hline & & & & $0^{\circ} / \mathrm{HT} 2 / 780^{\circ} \mathrm{C}$ & & & & & \\ \hline LPBF & CM247LC & ASTM E8/E8M & & $90^{\circ}$ & $\sim 790.8$ & $\sim 1012.7$ & $\sim 5.53$ & & [270] \\ \hline \multirow[t]{5}{*}{LPBF} & CM247LC & & & $\mathrm{AB}$ & & & & 220 & [311] \\ \hline & & & & STA $\left(1210^{\circ} \mathrm{C}\right)$ & & & & 226 & \\ \hline & & & & STA $\left(1230^{\circ} \mathrm{C}\right)$ & & & & 227 & \\ \hline & & & & STA $\left(1240^{\circ} \mathrm{C}\right)$ & & & & 226 & \\ \hline & & & & STA $\left(1260^{\circ} \mathrm{C}\right)$ & & & & 227 & \\ \hline \multirow[t]{3}{*}{LPBF} & EP718 & & & $0^{\circ} / \mathrm{AB}$ & 586 & 845 & 27 & & [173] \\ \hline & & & & $0^{\circ} / \mathrm{SA}$ & 1046 & 1301 & 6 & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}+\mathrm{SA}$ & 1025 & 1306 & 24.4 & & \\ \hline \multirow[t]{4}{*}{LPBF} & FGH100L & & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & $\mathrm{AB}$ & $761 \pm 16$ & $874 \pm 12$ & $3.25 \pm 0.16$ & & [324] \\ \hline & & & & HT & $1094 \pm 14$ & $1232 \pm 21$ & $6.35 \pm 0.28$ & & \\ \hline & & & & HIP & $879 \pm 21$ & $1146 \pm 20$ & $10.17 \pm 0.11$ & & \\ \hline & & & & $\mathrm{HIP}+\mathrm{HT}$ & $1155 \pm 17$ & $1307 \pm 15$ & $5.59 \pm 0.25$ & & \\ \hline \multirow[t]{5}{*}{LPBF} & FGH4096 M & & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & $\mathrm{AB}$ & 958.22 & 1204.13 & 24.97 & & [341] \\ \hline & & & & DA & 1459.46 & 1595.56 & 5.49 & & \\ \hline & & & & $\operatorname{STA}\left(1050^{\circ} \mathrm{C}\right)$ & 1039.86 & 1299.75 & 15.68 & & \\ \hline & & & & STA $\left(1130^{\circ} \mathrm{C}\right)$ & 1006.36 & 1322.02 & 14.44 & & \\ \hline & & & & Double Aging & 1037.86 & 1325.72 & 10.74 & & \\ \hline LPBF & GH648 & & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & 325 W/annealed & 890 & & $\sim 40$ & & [143] \\ \hline \multirow{2}{*}{LPBF} & Hastelloyx & & Strain rate $1.5 \mathrm{~mm} / \mathrm{min}$ & $0^{\circ} / \mathrm{AB}$ & $480 \pm 10$ & $620 \pm 15$ & $40 \pm 1$ & $149 \pm 9$ & [178] \\ \hline & & & & $0^{\circ} / \mathrm{HIP}$ & $350 \pm 6$ & $560 \pm 9$ & $41 \pm 4$ & $150 \pm 5$ & \\ \hline \multirow[t]{9}{*}{LPBF} & Hastelloyx & & & $\mathrm{AB}$ (built in the centre of the building plate)/ & $\sim 815.5$ & $\sim 936.5$ & $\sim 35.5$ & & [255] \\ \hline & & & & & & & & & \\ \hline & & & & \begin{tabular}{l} $\mathrm{AB}$ (built in the corner of the building plate)/ \\ $\mathrm{RT}$ \\ \end{tabular} & $\sim 812.5$ & $\sim 924.5$ & $\sim 34.5$ & & \\ \hline & & & & HIP (centre)/RT & $\sim 557$ & $\sim 839.5$ & $\sim 30$ & & \\ \hline & & & & HIP (corner)/RT & $\sim 556.5$ & $\sim 841.5$ & $\sim 29.5$ & & \\ \hline & & & & $\mathrm{AB}$ (centre) $/ 750^{\circ} \mathrm{C}$ & $\sim 544$ & $\sim 756$ & $\sim 22.5$ & & \\ \hline & & & & $\mathrm{AB}$ (corner) $/ 750^{\circ} \mathrm{C}$ & $\sim 544$ & $\sim 757.5$ & $\sim 21.5$ & & \\ \hline & & & & HIP (centre) $/ 750^{\circ} \mathrm{C}$ & $\sim 416$ & $\sim 729$ & $\sim 18.5$ & & \\ \hline & & & & HIP (corner) $/ 750^{\circ} \mathrm{C}$ & $\sim 412.5$ & $\sim 726$ & $\sim 19$ & & \\ \hline \multirow[t]{8}{*}{LPBF} & Hastelloyx & ASTM E8/E8M & Strain rate $0.05 \mathrm{~s}^{-1}$ & $0^{\circ} / \mathrm{AB}$ & $\sim 650.15$ & 698.59 & 8.79 & & [179] \\ \hline & & & & $90^{\circ} / \mathrm{AB}$ & 601.26 & 784.98 & 27.67 & & \\ \hline & & & & $0^{\circ} / \mathrm{HT}$ & 413.73 & 672.30 & 22.5 & & \\ \hline & & & & $90^{\circ} / \mathrm{HT}$ & 417.14 & 717.37 & 36.98 & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}$ & 468.85 & 807.51 & 39.83 & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}$ & 431.16 & 754.93 & 41.90 & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}+\mathrm{HT}$ & 430.53 & 777.47 & 49.14 & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}+\mathrm{HT}$ & 400.33 & 739.91 & 55.6 & & \\ \hline \multirow[t]{4}{*}{LPBF} & Original Hastelloy X (OHX) & ASTM E21 & & $\mathrm{OHX} / \mathrm{RT}$ & $\sim 727.88$ & 889.38 & 22.63 & & [284] \\ \hline & Modified Hastelloy X (MHX) & & & MHX/RT & 723.45 & 882.74 & 24.74 & & \\ \hline & & & & OHX/1033K & 384.96 & 482.30 & 46.11 & & \\ \hline & & & & MHX/1033K & 400.44 & 502.21 & 15.68 & & \\ \hline \multirow[t]{12}{*}{LPBF} & Hastelloy X & ASTM E8 & Strain rate $0.015 \mathrm{~mm} / \mathrm{min}$ & $90^{\circ} / \mathrm{AB} / \mathrm{RT}$ & $663 \pm 12$ & $773 \pm 9$ & 22 & & [180] \\ \hline & & & & $90^{\circ} / \mathrm{HT} / \mathrm{RT}$ & 420 & 723 & 42 & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP} / \mathrm{RT}$ & 440 & 730 & 48 & & \\ \hline & & & & $0^{\circ} / \mathrm{AB} / \mathrm{RT}$ & $792 \pm 1$ & $923 \pm 9$ & 25 & & \\ \hline & & & & $0^{\circ} / \mathrm{HT} / \mathrm{RT}$ & 370 & $793 \pm 20$ & $50 \pm 2$ & & \\ \hline & & & & $0^{\circ} /$ HIP $/$ RT & 400 & $800 \pm 10$ & 45 & & \\ \hline & & & Strain rate: $0.005 \mathrm{~mm} / \mathrm{min}$ & $90^{\circ} / \mathrm{AB} / 750^{\circ} \mathrm{C}$ & $386 \pm 19$ & 453 & 11 & & \\ \hline & & & & $90^{\circ} / \mathrm{HT} / 750^{\circ} \mathrm{C}$ & 270 & 420 & 40 & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP} / 750^{\circ} \mathrm{C}$ & 270 & 410 & 38 & & \\ \hline & & & & $0^{\circ} / \mathrm{AB} / 750^{\circ} \mathrm{C}$ & $460 \pm 17$ & $543 \pm 17$ & 12 & & \\ \hline & & & & $0^{\circ} / \mathrm{HT} / 750^{\circ} \mathrm{C}$ & 230 & $450 \pm 2$ & 38 & & \\ \hline & & & & & & & & & t page) \\ \hline \end{tabular} \end{center} Table 9 (continued) \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample Condition & Yield strength (MPa) & \begin{tabular}{l} Tensile strength \\ (MPa) \\ \end{tabular} & Elongation (\%) & \begin{tabular}{l} Young's modulus \\ (GPa) \\ \end{tabular} & Ref \\ \hline & & & & $0^{\circ} / \mathrm{HIP} / 750^{\circ} \mathrm{C}$ & 230 & $440 \pm 2$ & 38 & & \\ \hline \multirow{2}{*}{LPBF} & Hastelloy X & ASTM E8 & Strain rate $2 \mathrm{~mm} / \mathrm{min}$ & $\mathrm{AB}$ & $730 \pm 20$ & & $14 \pm 1$ & & [334] \\ \hline & & & & $0.2 \%$ less $\mathrm{Si}, \mathrm{Mn}$ and $\mathrm{C}$ (than above)/AB & $590 \pm 5$ & & $37 \pm 2$ & & \\ \hline \multirow[t]{2}{*}{LPBF} & Hastelloy $\mathrm{X}$ & & Strain rate $1.33 \times 10^{-3} \mathrm{~s}^{-1}$ & Hastelloy $\mathrm{X} / \mathrm{AB}$ & 584 & 734 & $19 \pm 2.8$ & & [340] \\ \hline & & & & Hastelloy $\mathrm{X}+1 \mathrm{wt} \% \mathrm{TiC}$ & $682 \pm 5.6$ & $849 \pm 1$ & $15 \pm 4.2$ & & \\ \hline LPBF & Hastelloy X & ASTM-E8/E8M and ASTM-E21 & Strain rate $0.05 \mathrm{~s}^{-1}$ & \begin{tabular}{l} Various build orientations and high \\ tempertaure tensile test \\ \end{tabular} & See Fig. 5 in the origin & 1 reference paper & & & [117] \\ \hline \multirow[t]{6}{*}{LPBF} & Haynes $\mathbb{8} 230 ®$ & DIN 50125 A-5x25 & & $0^{\circ} /$ Energy density $116 \mathrm{~J} / \mathrm{mm}^{3}$ & $798 \pm 5$ & $1102 \pm 3$ & $28 \pm 1$ & $205 \pm 4$ & [279] \\ \hline & & & & $90^{\circ} / 116 \mathrm{~J} / \mathrm{mm}^{3}$ & $656 \pm 4$ & $941 \pm 2$ & $32 \pm 3$ & $152 \pm 1$ & \\ \hline & & & & $0^{\circ} / 77 \mathrm{~J} / \mathrm{mm}^{3}$ & $794 \pm 6$ & $1087 \pm 5$ & $25 \pm 2$ & $201 \pm 2$ & \\ \hline & & & & $90^{\circ} / 77 \mathrm{~J} / \mathrm{mm}^{3}$ & $681 \pm 2$ & $979 \pm 5$ & $25 \pm 4$ & $165 \pm 3$ & \\ \hline & & & & $0^{\circ} / 66 \mathrm{~J} / \mathrm{mm}^{3}$ & $798 \pm 10$ & $1077 \pm 11$ & $21 \pm 2$ & $201 \pm 3$ & \\ \hline & & & & $90^{\circ} / 66 \mathrm{~J} / \mathrm{mm}^{3}$ & $702 \pm 6$ & $991 \pm 19$ & $16 \pm 4$ & $179 \pm 4$ & \\ \hline \multirow[t]{3}{*}{Micro laser aided AM} & IN100 & ASTM E8/E8M & & $\mathrm{AB}$ & $\sim 821.95$ & 1029.27 & 8.98 & & [280] \\ \hline & & & & $\mathrm{SA} / \mathrm{RT}$ & 956.10 & 1048.78 & 4.96 & & \\ \hline & & & & $\mathrm{SA} / 600^{\circ} \mathrm{C}$ & 904.88 & 1063.41 & 5.00 & & \\ \hline \multirow[t]{2}{*}{LPBF} & IN625 & & Strain rate, $10^{-3} \mathrm{~s}^{-1}$ & $10^{\circ} / \mathrm{AB}$ & 711 & 976 & \begin{tabular}{l} $3.0 \% \mathrm{RA}$ \\ $35 \%$ \\ \end{tabular} & 177 & [252] \\ \hline & & & & $35^{\circ} / \mathrm{AB}$ & 727 & 971 & $35 \% \mathrm{RA}$ & 179 & \\ \hline \multirow[t]{16}{*}{LPBF} & IN625 & & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & $0^{\circ} / \mathrm{AB} / \mathrm{RT}$ & 127 & תו & $\sim 27.3$ & & [263] \\ \hline & & & & $90^{\circ} / \mathrm{AB} / \mathrm{RT}$ & & & $\sim 36.9$ & & \\ \hline & & & & $0^{\circ} / \mathrm{SR} / \mathrm{RT}$ & & & $\sim 30$ & & \\ \hline & & & & $90^{\circ} / \mathrm{SR} / \mathrm{RT}$ & & & $\sim 39.3$ & & \\ \hline & & & & $0^{\circ} / \mathrm{ST} / \mathrm{RT}$ & & & $\sim 48.3$ & & \\ \hline & & & & $90^{\circ} / \mathrm{ST} / \mathrm{RT}$ & & & $\sim 52.7$ & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP} / \mathrm{RT}$ & & & $\sim 53.0$ & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP} / \mathrm{RT}$ & & & $\sim 58.7$ & & \\ \hline & & & & $0^{\circ} / \mathrm{AB} / 760^{\circ} \mathrm{C}$ & & & $\sim 6.0$ & & \\ \hline & & & & $90^{\circ} / \mathrm{AB} / 760^{\circ} \mathrm{C}$ & & & $\sim 22.1$ & & \\ \hline & & & & $0^{\circ} / \mathrm{SR} / 760^{\circ} \mathrm{C}$ & & & $\sim 12.3$ & & \\ \hline & & & & $90^{\circ} / \mathrm{SR} / 760^{\circ} \mathrm{C}$ & & & $\sim 37.7$ & & \\ \hline & & & & $0^{\circ} / \mathrm{ST} / 760^{\circ} \mathrm{C}$ & & & $\sim 15.0$ & & \\ \hline & & & & $90^{\circ} / \mathrm{ST} / 760^{\circ} \mathrm{C}$ & & & $\sim 12.3$ & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP} / 760^{\circ} \mathrm{C}$ & & & $\sim 24.8$ & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP} / 760^{\circ} \mathrm{C}$ & & & $\sim 21.6$ & & \\ \hline \multirow{2}{*}{LPBF} & IN625 & & Strain rate $\sim 10^{-3} \mathrm{~s}^{-1}$ & $90^{\circ} / \mathrm{HIP}$ & 360 & 880 & 58 & & [264] \\ \hline & & & & $0^{\circ} / \mathrm{HIP}$ & 380 & 900 & 58 & & \\ \hline \multirow[t]{3}{*}{LPBF} & IN625 & ASTM E8/E8M & Strain rate $0.5 \mathrm{~mm} / \mathrm{min}$ & $\mathrm{AB}$ & $641.5 \pm 23.5$ & $878.5 \pm 1.5$ & $30 \pm 2$ & $196 \pm 12$ & [266] \\ \hline & & & & $\mathrm{AB}$ carbon nanotube strengthened & $788 \pm 29$ & $998 \pm 34$ & $19.1 \pm 0.1$ & $378 \pm 12$ & \\ \hline & & & & HT carbon nanotube strengthened & $585 \pm 10$ & $1000 \pm 3$ & $31.5 \pm 0.5$ & $293 \pm 5$ & \\ \hline \multirow[t]{26}{*}{LPBF} & IN625 & & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & $90^{\circ} / \mathrm{AB}$ & $\sim 0.60$ & 0.82 & 36.6 & & [269] \\ \hline & & & & $90^{\circ} / \mathrm{SR}$ & 0.55 & 0.81 & 39 & & \\ \hline & & & & $90^{\circ} / \mathrm{RA}$ & 0.48 & 0.78 & 43.1 & & \\ \hline & & & & $90^{\circ} / \mathrm{ST}$ & 0.4 & 0.79 & 52.5 & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP} \mathrm{V}$ & 0.35 & 0.78 & 58.7 & & \\ \hline & & & & $0^{\circ} / \mathrm{AB}$ & 0.65 & 0.85 & 27.5 & & \\ \hline & & & & $0^{\circ} / \mathrm{SR}$ & 0.58 & 0.89 & 29.8 & & \\ \hline & & & & $0^{\circ} / \mathrm{RA}$ & 0.51 & 0.83 & 37.5 & & \\ \hline & & & & $0^{\circ} / \mathrm{ST}$ & 0.4 & 0.81 & 48.2 & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}$ & 0.35 & 0.78 & 53.1 & & \\ \hline & & & & $45^{\circ} / \mathrm{AB}$ & 0.67 & 0.87 & 30.8 & & \\ \hline & & & & $45^{\circ} / \mathrm{SR}$ & 0.62 & 0.93 & 34.7 & & \\ \hline & & & & $45^{\circ} / \mathrm{RA}$ & 0.52 & 0.83 & 41.7 & & \\ \hline & & & & $45^{\circ} / \mathrm{ST}$ & 0.40 & 0.83 & 52.9 & & \\ \hline & & & & $45^{\circ} / \mathrm{HIP}$ & 0.36 & 0.80 & 53.8 & & \\ \hline & & & & $90^{\circ} / \mathrm{AB} / 760^{\circ} \mathrm{C}$ & 0.35 & 0.36 & 22.4 & & \\ \hline & & & & $90^{\circ} / \mathrm{SR} / 760^{\circ} \mathrm{C}$ & 0.36 & 0.39 & 37.9 & & \\ \hline & & & & $90^{\circ} / \mathrm{RA} / 760^{\circ} \mathrm{C}$ & 0.30 & 0.32 & 36.6 & & \\ \hline & & & & $90^{\circ} / \mathrm{ST} / 760^{\circ} \mathrm{C}$ & 0.26 & 0.31 & 12.4 & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP} / 760^{\circ} \mathrm{C}$ & 0.23 & 0.34 & 21.5 & & \\ \hline & & & & $0^{\circ} / \mathrm{AB} / 760^{\circ} \mathrm{C}$ & 0.36 & 0.37 & 5.9 & & \\ \hline & & & & $0^{\circ} / \mathrm{SR} / 760^{\circ} \mathrm{C}$ & 0.37 & 0.40 & 12.2 & & \\ \hline & & & & $0^{\circ} / \mathrm{RA} / 760^{\circ} \mathrm{C}$ & 0.32 & 0.34 & 11.7 & & \\ \hline & & & & $0^{\circ} / \mathrm{ST} / 760^{\circ} \mathrm{C}$ & 0.26 & 0.32 & 15.1 & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP} / 760^{\circ} \mathrm{C}$ & 0.21 & 0.35 & 25.0 & & \\ \hline & & & & $45^{\circ} / \mathrm{AB} / 760^{\circ} \mathrm{C}$ & 0.39 & 0.41 & \begin{tabular}{l} 7.0 \\ 7.0 \\ \end{tabular} & & \\ \hline \end{tabular} \end{center} \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample Condition & Yield strength (MPa) & \begin{tabular}{l} Tensile strength \\ (MPa) \\ \end{tabular} & Elongation (\%) & \begin{tabular}{l} Young's modulus \\ (GPa) \\ \end{tabular} & Ref \\ \hline & & & & $45^{\circ} / \mathrm{SR} / 760^{\circ} \mathrm{C}$ & 0.40 & 0.41 & 11.1 & & \\ \hline & & & & $45^{\circ} / \mathrm{RA} / 760^{\circ} \mathrm{C}$ & 0.34 & 0.35 & 9.9 & & \\ \hline & & & & $45^{\circ} / \mathrm{ST} / 760^{\circ} \mathrm{C}$ & 0.28 & 0.33 & 11.1 & & \\ \hline & & & & $45^{\circ} / \mathrm{HIP} / 760^{\circ} \mathrm{C}$ & 0.24 & 0.35 & 23.5 & & \\ \hline LPBF & IN625 & ASTM E8/E8M & Strain rate $0.015 \mathrm{~min}^{-1}$ & $0^{\circ}$ & $\sim 765.86$ & $\sim 1068.99$ & $\sim 33.86$ & $\sim 185.60$ & [272] \\ \hline \multirow[t]{2}{*}{LPBF} & IN625 & ISO-7500/1 & Strain rate $1 \mathrm{~mm} / \mathrm{min}$ & $0^{\circ}$ & $800 \pm 20$ & $1030 \pm 50$ & $\sim 8-10$ & & [282] \\ \hline & & & & $90^{\circ}-1$ & $720 \pm 30$ & $1070 \pm 60$ & $\sim 8-10$ & & \\ \hline \multirow[t]{2}{*}{LPBF} & IN625 & EN-10002/ISO-6892 & & $0^{\circ}-1-2-1$ & $734 \pm 3$ & $1036 \pm 3$ & $36 \pm 0.3$ & $200 \pm 3$ & [288] \\ \hline & & & & $90^{\circ}$ & $579 \pm 5$ & $888 \pm 6$ & $40 \pm 1$ & $159 \pm 5$ & \\ \hline \multirow[t]{4}{*}{LPBF} & IN625 & ASTM E8/E8M & Strain rate $8.10^{-3} \mathrm{~s}^{-1}$ & $\mathrm{AB}$ & $783 \pm 23$ & $1041 \pm 36$ & $33 \pm 1$ & & [303] \\ \hline & & & & DA & $1012 \pm 54$ & $1222 \pm 56$ & $23 \pm 1$ & & \\ \hline & & & & STA & $722 \pm 7$ & $1116 \pm 6$ & $35 \pm 5$ & & \\ \hline & & & & ST & $396 \pm 9$ & $883 \pm 15$ & $55 \pm 1$ & & \\ \hline \multirow[t]{2}{*}{LPBF} & IN625 & & & $0^{\circ}-1-2$ & $396 \pm 33$ & $906 \pm 28$ & $62.34 \pm 1.98$ & $561 \pm 14$ & $[100]$ \\ \hline & & & & $90^{\circ}-$ & $349 \pm 5$ & $842 \pm 29$ & $56.3 \pm 6.24$ & $539 \pm 58$ & \\ \hline \multirow{4}{*}{LPBF} & IN625 & & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & $0^{\circ} / \mathrm{SR}$ & $718 \pm 13$ & $1069 \pm 11$ & $37 \pm 2$ & $202 \pm 3$ & [312] \\ \hline & & & & $90^{\circ} / \mathrm{SR}$ & $685 \pm 87$ & $1009 \pm 56$ & $43 \pm 5$ & $195 \pm 12$ & \\ \hline & & & & \begin{tabular}{l} 0 \\ $0^{\circ} / \mathrm{HIP}$ \\ \end{tabular} & $442 \pm 6$ & $933 \pm 14$ & $43 \pm 4$ & \begin{tabular}{l} $212 \pm 7$ \\ \end{tabular} & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}$ & $425 \pm 9$ & $923 \pm 21$ & $46 \pm 5$ & $214 \pm 7$ & \\ \hline \multirow[t]{5}{*}{LPBF} & IN625 & & $0.5 \mathrm{mN}$ at $1 / 10$ of layer thickness & AB/0 dpa & & & & $\sim 225$ & [376] \\ \hline & & & & $\mathrm{AB} / 0.1 \mathrm{dpa}$ & & & & $\sim 220$ & \\ \hline & & & & AB/0.5 dpa & & & & $\sim 220$ & \\ \hline & & & & $\mathrm{AB} / 1 \mathrm{dpa}$ & & & & $\sim 210$ & \\ \hline & & & & $\mathrm{AB} / 3 \mathrm{dpa}$ & & & & $\sim 215$ & \\ \hline \multirow[t]{3}{*}{LPBF} & IN625 & & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & $\mathrm{AB}$ & $652 \pm 10$ & $925 \pm 13$ & $32 \pm 3$ & $145 \pm 4$ & [325] \\ \hline & & & & $\mathrm{AB}+90^{\circ} \mathrm{C} / 1 \mathrm{~h}$ & $567 \pm 15$ & $869 \pm 7$ & $38 \pm 1$ & $142 \pm 11$ & \\ \hline & & & & $\mathrm{AB}+1100^{\circ} \mathrm{C} / 1 \mathrm{~h}$ & $409 \pm 14$ & $886 \pm 11$ & $56 \pm 5$ & $114 \pm 8$ & \\ \hline \multirow[t]{2}{*}{LPBF} & IN625 & & $815^{\circ} \mathrm{C}$ Average load: $1021 \mathrm{~N}$ & ST & & 355.6 & & & [329] \\ \hline & & & $815^{\circ} \mathrm{C}$ Average load: $1147 \mathrm{~N}$ & Welding zone & & 392.5 & & & \\ \hline \multirow[t]{2}{*}{LPBF} & IN625 & & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & Along the laser scanning & 686.7 & 967.5 & 11.5 & & [416] \\ \hline & & & & Normal to laser scanning & 747.9 & 1077.8 & 9 & & \\ \hline \multirow[t]{12}{*}{LPBF} & IN625 & ISO 6892-1:2009 & Strain rate $2.5 \times 10^{-4} \mathrm{~s}^{-1}$ & Build orientation on X-axis; $90^{\circ}$ interlayer & $560 \pm 5$ & $877 \pm 8$ & $39 \pm 3$ & & [417] \\ \hline & & & & \begin{tabular}{l} Build orientation on X-axis; $67^{\circ}$ interlayer \\ rotation \\ \end{tabular} & $619 \pm 11$ & $962 \pm 11$ & $43 \pm 2$ & & \\ \hline & & & & \begin{tabular}{l} Build orientation on X-axis; $45^{\circ}$ interlayer \\ rotation \\ \end{tabular} & $627 \pm 11$ & $991 \pm 6$ & $42 \pm 2$ & & \\ \hline & & & & \begin{tabular}{l} Build orientation on Y-axis; $90^{\circ}$ interlayer \\ rotation \\ \end{tabular} & $559 \pm 5$ & $874 \pm 10$ & $38 \pm 5$ & & \\ \hline & & & & \begin{tabular}{l} Build orientation on Y-axis; $67^{\circ}$ interlayer \\ rotation \\ \end{tabular} & $616 \pm 11$ & $946 \pm 11$ & $42 \pm 3$ & & \\ \hline & & & & \begin{tabular}{l} Build orientation on Y-axis; $45^{\circ}$ interlayer \\ rotation \\ \end{tabular} & $630 \pm 9$ & $993 \pm 5$ & $43 \pm 2$ & & \\ \hline & & & & \begin{tabular}{l} Build orientation on Z-axis; $90^{\circ}$ interlayer \\ rotation \\ \end{tabular} & $518 \pm 5$ & $814 \pm 4$ & $50 \pm 41$ & & \\ \hline & & & & \begin{tabular}{l} Build orientation on Z-axis; $67^{\circ}$ interlayer \\ rotation \\ \end{tabular} & $546 \pm 10$ & $825 \pm 3$ & $53 \pm 1$ & & \\ \hline & & & & \begin{tabular}{l} Build orientation on Z-axis; $45^{\circ}$ interlayer \\ rotation \\ \end{tabular} & $551 \pm 6$ & $824 \pm 5$ & $53 \pm 1$ & & \\ \hline & & & & \begin{tabular}{l} Build orientation at $45^{\circ} ; 90^{\circ}$ interlayer \\ rotation \\ \end{tabular} & $551 \pm 6$ & $870 \pm 7$ & $48 \pm 1$ & & \\ \hline & & & & \begin{tabular}{l} Build orientation at $45^{\circ} ; 67^{\circ}$ interlayer \\ rotation \\ \end{tabular} & $583 \pm 5$ & $910 \pm 7$ & $48 \pm 2$ & & \\ \hline & & & & \begin{tabular}{l} Build orientation at $45^{\circ} ; 45^{\circ}$ interlayer \\ rotation \\ \end{tabular} & $643 \pm 5$ & $990 \pm 9$ & $46 \pm 2$ & & \\ \hline \multirow[t]{2}{*}{LPBF} & IN625 & ASTM E8M & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & HIP and tested at RT & $459.5 \pm 6.5$ & $926.0 \pm 12.7$ & $74.9 \pm 2.9$ & & [438] \\ \hline & & & & HIP and tested at $650^{\circ} \mathrm{C}$ & $246.2 \pm 3.7$ & $637.5 \pm 15.4$ & $38.7 \pm 4.6$ & & \\ \hline LPBF & IN625 & ISO 6892-1:2009 & & Different notches & \multicolumn{3}{|c|}{See the original reference paper} & & [439] \\ \hline LPBF & IN625 & & & \begin{tabular}{l} Various temperatures $\left(20^{\circ} \mathrm{C}, 540^{\circ} \mathrm{C}, 760^{\circ} \mathrm{C}\right.$, \\ $815^{\circ} \mathrm{C} 870^{\circ} \mathrm{C} 950^{\circ} \mathrm{C}$ and $100{ }^{\circ} \mathrm{C}$ \\ \end{tabular} & \multicolumn{3}{|c|}{See Figs. 6 and 7 in the original reference paper} & & [331] \\ \hline \multirow{2}{*}{LPBF} & IN718 & Chinese grain boundaries/T 228 & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & $0^{\circ}$ & 780 & 1069.6 & 30.9 & & [247] \\ \hline & & & & $90^{\circ}-1$ & 634 & 980 & & & \\ \hline \multirow[t]{4}{*}{LPBF} & IN718 & & Strain rate $4.25 \times 10^{-4} \mathrm{~s}^{-1}$ & $\mathrm{AB} / \mathrm{RT}$ & 677 & 1023 & 28.1 & & [250] \\ \hline & & & & SA/RT & 1271 & 1425 & 18.6 & & \\ \hline & & & & $\mathrm{AB} / 650^{\circ} \mathrm{C}$ & 594 & 862.0 & 25.1 & & \\ \hline & & & & $\mathrm{SA} / 650^{\circ} \mathrm{C}$ & 1042 & 1142 & 10.1 & & \\ \hline \end{tabular} \end{center} \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample Condition & Yield strength (MPa) & \begin{tabular}{l} Tensile strength \\ (MPa) \\ \end{tabular} & Elongation (\%) & \begin{tabular}{l} Young's modulus \\ (GPa) \\ \end{tabular} & Ref \\ \hline \multirow[t]{12}{*}{LPBF} & IN718 & ASTM E8/E8M & & \begin{tabular}{l} Top left corner (TL), laser focus $3 \mathrm{~mm}$ above \\ the building plate $(+3 \mathrm{~mm})$ \\ \end{tabular} & 1234 & 1455 & 14.4 & $19.2 \% \mathrm{RA}$ & [315] \\ \hline & & & & $\mathrm{TL}$, laser focus on the building plate $(0 \mathrm{~mm})$ & 1207 & 1455 & 14.5 & $16.7 \% \mathrm{RA}$ & \\ \hline & & & & \begin{tabular}{l} TL, laser focus $3 \mathrm{~mm}$ below the building \\ plate $(-3 \mathrm{~mm})$ \\ \end{tabular} & 1207 & 1372 & 3.6 & $12.2 \% \mathrm{RA}$ & \\ \hline & & & & Middle left (ML), $+3 \mathrm{~mm}$ & 1179 & 1475 & 13.9 & $25.4 \% \mathrm{RA}$ & \\ \hline & & & & ML, $0 \mathrm{~mm}$ & 1200 & 1448 & 15.5 & $30.3 \%$ RA & \\ \hline & & & & ML, $-3 \mathrm{~mm}$ & 1213 & 1434 & 10.2 & $18.0 \% \mathrm{RA}$ & \\ \hline & & & & Top right corner (TR), $+3 \mathrm{~mm}$ & 1213 & 1420 & 6.9 & $14.4 \% \mathrm{RA}$ & \\ \hline & & & & RT, $0 \mathrm{~mm}$ & 1213 & 1427 & 9.3 & $15.1 \%$ RA & \\ \hline & & & & RT, $-3 \mathrm{~mm}$ & 1220 & 1406 & 5.7 & 10.1\%RA & \\ \hline & & & & Middle right (MR), $3 \mathrm{~mm}$ & 1207 & 1448 & 12.2 & $18.0 \% \mathrm{RA}$ & \\ \hline & & & & MR, $0 \mathrm{~mm}$ & 1213 & 1455 & 10.8 & 13.1\%RA & \\ \hline & & & & $\mathrm{MR},-3 \mathrm{~mm}$ & 1179 & 1358 & 4.4 & $13.9 \% \mathrm{RA}$ & \\ \hline LPBF (DMLS) & IN718 & ASTM-E8/E8M & & Lattice dog bone specimen & 1005.8 & 1220.8 & 22.3 & 166.5 & [220] \\ \hline \multirow[t]{7}{*}{LPBF} & IN718 & & Strain rate $4 \times 10^{-3} \mathrm{~s}^{-1}$ & АВ & 580 & 845 & & 100.5 & $[172]$ \\ \hline & & & & ST & 535 & 870 & & & \\ \hline & & & & SA & 1240 & 1400 & & & \\ \hline & & & & HIP & 430 & 875 & & & \\ \hline & & & & HIP + Aging & 1100 & 1315 & & & \\ \hline & & & & Arc-PVD + HIP & 420 & 815 & & & \\ \hline & & & & Arc-PVD + HIP + Aging & 1185 & 1300 & & & \\ \hline \multirow[t]{5}{*}{LPBF} & IN718 & ASTM-E8/E8M & Strain rate $2 \mathrm{~mm} / \mathrm{min}$ & $A B$ & $596 \pm 30$ & $943 \pm 8$ & $35 \pm 1$ & $170 \pm 9$ & [256] \\ \hline & & & & HT A & $924 \pm 11$ & $1186 \pm 2$ & $25 \pm 5$ & $158 \pm 17$ & \\ \hline & & & & нт в & $951 \pm 3$ & $1210 \pm 23$ & $23 \pm 1$ & $195 \pm 2$ & \\ \hline & & & & HT C & $1158 \pm 14$ & $1339 \pm 30$ & $7 \pm 1$ & $138 \pm 6$ & \\ \hline & & & & HT D & $558 \pm 7$ & $933 \pm 3$ & $43 \pm 1$ & $170 \pm 7$ & \\ \hline \multirow{8}{*}{LPBF (DMLS)} & IN718 & ASTM-E8/E8M & Strain rate $102 \mu \mathrm{m} / \mathrm{s}$ & $0^{\circ} /$ Powder I \#1 & 1070 & 1316 & 18 & 195 & [257] \\ \hline & & & & $90^{\circ} /$ Powder I \#2 & 1082 & 1331 & 20 & 199 & \\ \hline & & & & $90^{\circ} /$ Powder I \#5 & 1071 & 1322 & 20 & 198 & \\ \hline & & & & $0^{\circ} /$ Powder I \#6 & 1059 & 1293 & 17 & 191 & \\ \hline & & & & $0^{\circ} /$ Powder II \#2 & 789 & 1059 & 31 & 174 & \\ \hline & & & & $90^{\circ} /$ Powder II \#3 & 868 & 1162 & 26 & 182 & \\ \hline & & & & $0^{\circ} /$ Powder II \#4 & 787 & 1034 & 31 & 185 & \\ \hline & & & & $90^{\circ} /$ Powder II \#5 & 854 & 1148 & 26 & 172 & \\ \hline \multirow[t]{6}{*}{LPBF} & IN718 & ASTM-E8/E8M & $24^{\circ} \mathrm{C}$ or $650^{\circ} \mathrm{C}$ & $0^{\circ} / \mathrm{STA} / 24^{\circ} \mathrm{C}$ & 1295 & 1484 & & & [258] \\ \hline & & & & $45^{\circ} / \mathrm{STA} / 24^{\circ} \mathrm{C}$ & 1368 & 1521 & & & \\ \hline & & & & $90^{\circ} / \mathrm{STA} / 24^{\circ} \mathrm{C}$ & 1240 & 1398 & & & \\ \hline & & & & $0^{\circ} / \mathrm{STA} / 650^{\circ} \mathrm{C}$ & 1033 & 1139 & & & \\ \hline & & & & $45^{\circ} / \mathrm{STA} / 650^{\circ} \mathrm{C}$ & 1124 & 1187 & & & \\ \hline & & & & $90^{\circ} / \mathrm{STA} / 650^{\circ} \mathrm{C}$ & 978 & 1114 & & & \\ \hline \multirow[t]{2}{*}{LPBF} & IN718 & & & $90^{\circ} / \mathrm{AB}$ & $711 \pm 14$ & $1110 \pm 11$ & $24.5 \pm 1.1$ & & [259] \\ \hline & & & & $0^{\circ} / \mathrm{AB}$ & $858 \pm 12$ & $1167 \pm 10$ & $21.5 \pm 1.3$ & & \\ \hline \multirow[t]{9}{*}{LPBF} & IN718 & & Strain rate $3 \times 10^{-3} \mathrm{~s}^{-1}$ & $0^{\circ} / \mathrm{RT}$ & 1186 & 1440 & 18.5 & & [261] \\ \hline & & & & $450^{\circ} \mathrm{C}$ & 1033 & 1216 & 12.4 & & \\ \hline & & & & $650^{\circ} \mathrm{C}$ & 870 & 1011 & 3.6 & & \\ \hline & & & & $90^{\circ} / \mathrm{RT}$ & 1180 & 1400 & 20.4 & & \\ \hline & & & & $450^{\circ} \mathrm{C}$ & 1026 & 1160 & 15.9 & & \\ \hline & & & & $650^{\circ} \mathrm{C}$ & 860 & 992 & 14.2 & & \\ \hline & & & & $45^{\circ} / \mathrm{RT}$ & 1190 & 1450 & 16.9 & & \\ \hline & & & & $450^{\circ} \mathrm{C}$ & 1080 & 1255 & 12.8 & & \\ \hline & & & & $650^{\circ} \mathrm{C}$ & 855 & 1074 & 5.8 & & \\ \hline \multirow[t]{5}{*}{LPBF} & IN718 & ASTM-E8/E8 & Strain rate $0.5 \mathrm{~mm} / \mathrm{min}$ & $\mathrm{AB}$ & $\sim 614$ & $\sim 957$ & $\sim 28.9$ & & $[66]$ \\ \hline & & & & SA (Standard) & $\sim 1211$ & $\sim 1391$ & $\sim 18.0$ & & \\ \hline & & & & SA (single step ageing) & $\sim 1211$ & $\sim 1391$ & $\sim 18.9$ & & \\ \hline & & & & SA (solution at $1100^{\circ} \mathrm{C}$ ) & $\sim 1142$ & $\sim 1304$ & $\sim 17.8$ & & \\ \hline & & & & $\mathrm{SA}$ (solution at $1200^{\circ} \mathrm{C}$ ) & $\sim 822$ & $\sim 1121$ & $\sim 21.5$ & & \\ \hline \multirow[t]{2}{*}{LPBF} & IN718 & ASTM-E8/E8M & Strain rate $10^{-4} \mathrm{~s}^{-1}$ & Raw powder & $1210 \pm 25$ & $1404 \pm 32$ & $18.5 \pm 1.6$ & & [115] \\ \hline & & & & Recycled powder & $1178 \pm 31$ & $1369 \pm 35$ & $17.4 \pm 1.7$ & & \\ \hline \multirow{4}{*}{LPBF (DMLS)} & IN718 & ASTM-E8/E8M & & $90^{\circ} / \mathrm{SA}$ & 1215 & & & 165 & $[74]$ \\ \hline & & & & $45^{\circ} / \mathrm{SA}$ & 1305 & & & 215 & \\ \hline & & & & $0^{\circ} / \mathrm{SA}$ & 1290 & & & 195 & \\ \hline & & & & $\mathrm{SA}+\mathrm{HIP}$ & 1125 & & & 200 & \\ \hline \multirow{2}{*}{LPBF (DMLS)} & IN718 & ISO 6892-1 & & $\mathrm{AB}$ & 666 & 1065 & 24 & & [268] \\ \hline & & & & Rolled with $15 \%$ deformation & 694 & 1144 & 22 & & \\ \hline \end{tabular} \end{center} \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample Condition & Yield strength (MPa) & \begin{tabular}{l} Tensile strength \\ (MPa) \\ \end{tabular} & Elongation (\%) & \begin{tabular}{l} Young's modulus \\ (GPa) \\ \end{tabular} & Ref \\ \hline & & & & $30 \%$ def. & 826 & 1446 & 20 & & \\ \hline & & & & $50 \%$ def. & 925 & 1594 & 17 & & \\ \hline & & & & $15 \%$ def. + SA & 753 & 1405 & 13 & & \\ \hline & & & & $30 \%$ def. + SA & 886 & 1578 & 8 & & \\ \hline & & & & $50 \%$ def. + SA & 977 & 1623 & 5 & & \\ \hline \multirow[t]{4}{*}{LPBF} & IN718 & ASTM-E8/E8M & & $\mathrm{AB}$ & 646.50 & 940.10 & 35.47 & & [271] \\ \hline & & & & SA & 1211.33 & 1408.50 & 14.83 & & \\ \hline & IN718/TiC & & & $\mathrm{AB}$ & 774.26 & 1029.00 & 12.32 & & \\ \hline & & & & SA & 1144.00 & 1380.86 & 9.08 & & \\ \hline \multirow[t]{4}{*}{LPBF} & IN718 & & & $\mathrm{AB} / 20^{\circ} \mathrm{C}$ & $569-646$ & $851-1002$ & $9.8-31.7$ & & [68] \\ \hline & & & & $\mathrm{HA} / 20^{\circ} \mathrm{C}$ & 1160 & 1350 & 17.6 & & \\ \hline & & & & $\mathrm{AB} / 1000^{\circ} \mathrm{C}$ & 112 & 114 & $47.4-53.5$ & & \\ \hline & & & & $\mathrm{HA} / 1000^{\circ} \mathrm{C}$ & 113 & 116 & 58.1 & & \\ \hline \multirow[t]{4}{*}{LPBF} & IN718 & ISO 6892-1 & & $0^{\circ} / 50 \mu \mathrm{m}$ layer thickness & 646 & 1049 & 27.2 & & [274] \\ \hline & & & & $90^{\circ} / 50 \mu \mathrm{m}$ layer thickness & 609 & 949 & 31.7 & & \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-50} \\ \hline & & & & $0^{\circ} / 30 \mu \mathrm{m}$ layer thickness & 807 & 1051 & 21.9 & & \\ \hline & & & & $90^{\circ} / 30 \mu \mathrm{m}$ layer thickness & 675 & 957 & 27.7 & & \\ \hline \multirow{14}{*}{LPBF} & IN718 & EN 10002 & & $0^{\circ} / \mathrm{AB} / \mathrm{RT}$ & $816 \pm 24$ & $1085 \pm 11$ & $19.1 \pm 0.7$ & & [277] \\ \hline & & & & $90^{\circ} / \mathrm{AB} / \mathrm{RT}$ & $737 \pm 4$ & $1010 \pm 10$ & $20.6 \pm 2.1$ & & \\ \hline & & & & $0^{\circ} / \mathrm{HT} 1 / \mathrm{RT}$ & $1227 \pm 1$ & $1447 \pm 10$ & $10.1 \pm 0.6$ & & \\ \hline & & & & $90^{\circ} / \mathrm{HT1} / \mathrm{RT}$ & $1136 \pm 16$ & $1357 \pm 5$ & $13.6 \pm 0.2$ & & \\ \hline & & & & $0^{\circ} / \mathrm{HT} 2 / \mathrm{RT}$ & $1222 \pm 26$ & $1417 \pm 4$ & $15.9 \pm 1.0$ & & \\ \hline & & & & $90^{\circ} / \mathrm{HT} 2 / \mathrm{RT}$ & $1186 \pm 23$ & $1387 \pm 12$ & $17.4 \pm 0.4$ & & \\ \hline & & & & $0^{\circ} / \mathrm{HT} 1 / 450^{\circ} \mathrm{C}$ & 1100120 & $\sim 1287.00$ & $\sim 13.06$ & & \\ \hline & & & & $90^{\circ} / \mathrm{HT} 1 / 450^{\circ} \mathrm{C}$ & & 1224.22 & 15.77 & & \\ \hline & & & & $0^{\circ} / \mathrm{HT} 2 / 450^{\circ} \mathrm{C}$ & & 1228.70 & 14.75 & & \\ \hline & & & & $90^{\circ} / \mathrm{HT} 2 / 450^{\circ} \mathrm{C}$ & & 1183.86 & 15.43 & & \\ \hline & & & & $0^{\circ} / \mathrm{HT} 1 / 650^{\circ} \mathrm{C}$ & & 1107.62 & 8.78 & & \\ \hline & & & & $90^{\circ} / \mathrm{HT} 1 / 650^{\circ} \mathrm{C}$ & & 1029.15 & 18.8 & & \\ \hline & & & & $0^{\circ} / \mathrm{HT} 2 / 650^{\circ} \mathrm{C}$ & & 1134.53 & 12.61 & & \\ \hline & & & & $90^{\circ} / \mathrm{HT} 2 / 650^{\circ} \mathrm{C}$ & & 1121.08 & 17.57 & & \\ \hline \multirow[t]{8}{*}{LPBF} & IN718 & & & $90^{\circ} / \mathrm{AB}$ & $572 \pm 44$ & $904 \pm 22$ & $19 \pm 4$ & $162 \pm 18$ & [95] \\ \hline & & & & $0^{\circ} / \mathrm{AB}$ & $643 \pm 63$ & $991 \pm 62$ & $13 \pm 6$ & $193 \pm 24$ & \\ \hline & & & & $45^{\circ} / \mathrm{AB}$ & $590 \pm 15$ & $954 \pm 15$ & $20 \pm 1$ & $200 \pm 23$ & \\ \hline & & & & $\mathrm{D} 45^{\circ} / \mathrm{AB}$ & $723 \pm 55$ & $1117 \pm 45$ & $16 \pm 3$ & $208 \pm 48$ & \\ \hline & & & & $90^{\circ} / \mathrm{SA}$ & $1074 \pm 42$ & $1320 \pm 6$ & $19 \pm 2$ & $163 \pm 30$ & \\ \hline & & & & $0^{\circ} / \mathrm{SA}$ & $1159 \pm 32$ & $1377 \pm 66$ & $8 \pm 6$ & $199 \pm 15$ & \\ \hline & & & & $45^{\circ} / \mathrm{SA}$ & $1152 \pm 24$ & $1371 \pm 5$ & $15 \pm 5$ & $188 \pm 19$ & \\ \hline & & & & $\mathrm{D} 45^{\circ} / \mathrm{SA}$ & $1241 \pm 68$ & $1457 \pm 55$ & $14 \pm 5$ & $209 \pm 44$ & \\ \hline \multirow[t]{8}{*}{LPBF} & IN718 & Nano indentation & & $90^{\circ} /$ top & & & & 188.9 & [278] \\ \hline & & & & $90^{\circ} /$ middle top & & & & 210.6 & \\ \hline & & & & $90^{\circ}$ /middle bottom & & & & 211.1 & \\ \hline & & & & $90^{\circ} /$ bottom & & & & 202.8 & \\ \hline & & & & $0^{\circ} /$ top & & & & 193.7 & \\ \hline & & & & $0^{\circ} /$ middle top & & & & 193.4 & \\ \hline & & & & $0^{\circ} /$ middle bottom & & & & 201.2 & \\ \hline & & & & $0^{\circ}$ /bottom & & & & 196.2 & \\ \hline \multirow{6}{*}{LPBF} & IN718 & & Strain rate $1.5 \mathrm{~mm} / \mathrm{min}$ & $90^{\circ} / 250 \mathrm{~W} / \mathrm{AB}$ & $668 \pm 16$ & $1011 \pm 27$ & $22 \pm 2$ & $173 \pm 13$ & [281] \\ \hline & & & & $90^{\circ} / 950 \mathrm{~W} / \mathrm{AB}$ & $531 \pm 9$ & $866 \pm 33$ & $21 \pm 7$ & $113 \pm 3$ & \\ \hline & & & & \begin{tabular}{l} Zone 1-250 W Matrix and Zone 2 - two lines \\ of $950 \mathrm{~W}$ \\ \end{tabular} & $574 \pm 6$ & $873 \pm 14$ & $13 \pm 2$ & $136 \pm 13$ & \\ \hline & & & & \begin{tabular}{l} Zone 1-950 W Matrix and Zone 2 - two lines \\ of $250 \mathrm{~W}$ \\ \end{tabular} & $591 \pm 14$ & $920 \pm 23$ & $15 \pm 3$ & $131 \pm 3$ & \\ \hline & & & & \begin{tabular}{l} Zone 1-250 W Matrix and Zone 2 - four lines \\ of $950 \mathrm{~W}$ \\ \end{tabular} & $585 \pm 7$ & $880 \pm 17$ & $14 \pm 1$ & $155 \pm 11$ & \\ \hline & & & & \begin{tabular}{l} Zone $1-950 \mathrm{~W}$ Matrix and Zone 2 - four lines \\ of $250 \mathrm{~W}$ \\ \end{tabular} & $586 \pm 16$ & $920 \pm 7$ & $18 \pm 2$ & $137 \pm 12$ & \\ \hline \multirow[t]{2}{*}{LPBF} & IN718 & grain boundaries/T 228-2002 & & $\mathrm{AB}$ & $889-907$ & $1137-1148$ & $19.2-25.9$ & 204 & [283] \\ \hline & & & & SA & 1102-1161 & $1280-1358$ & $10-22$ & 201 & \\ \hline \multirow[t]{2}{*}{LPBF} & IN718 & & & HSA & 1046 & 1371 & 12.3 & & [216] \\ \hline & & & & HA & 1174 & 1451 & 13.5 & & \\ \hline \multirow[t]{3}{*}{LPBF} & IN718 & & & $\mathrm{AB}$ & 849 & 1126 & 22.8 & & [165] \\ \hline & & & & SA & 1084 & 1371 & 10.1 & & \\ \hline & & & & HSA & 1046 & 1371 & 12.3 & & \\ \hline \multirow[t]{2}{*}{LPBF} & IN718 & & Strain rate $2 \mathrm{~mm} / \mathrm{s}$ & Island size $2 \times 2 \mathrm{~mm}^{2}$ & $804.0 \pm 49.5$ & $1076.5 \pm 28.9$ & $16.85 \pm 0.07$ & & [286] \\ \hline & & & & $3 \times 3 \mathrm{~mm}^{2}$ & $800.5 \pm 7.80$ & $1075.0 \pm 8.50$ & $21.05 \pm 0.21$ & & \\ \hline \end{tabular} \end{center} \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample Condition & Yield strength (MPa) & \begin{tabular}{l} Tensile strength \\ (MPa) \\ \end{tabular} & Elongation (\%) & \begin{tabular}{l} Young's modulus \\ (GPa) \\ \end{tabular} & Ref \\ \hline & & & & $5 \times 5 \mathrm{~mm}^{2}$ & $770.5 \pm 2.10$ & $1064.5 \pm 3.50$ & $22.35 \pm 0.21$ & & \\ \hline & & & & $7 \times 7 \mathrm{~mm}^{2}$ & $772.5 \pm 2.20$ & $1065.0 \pm 1.40$ & $25.25 \pm 0.35$ & & \\ \hline \multirow[t]{4}{*}{LPBF (DMLS)} & IN718 & ASTM-E8/E8M & Strain rate $0.0185 / \mathrm{min}$ & Hт & 1170 & 1380 & 9.0 & 188 & [287] \\ \hline & & & & $\mathrm{HIP}+\mathrm{HT}$ & 1090 & 1310 & 8.2 & 210 & \\ \hline & & & & $\mathrm{HT}+$ shot peen & 1110 & 1340 & 4.3 & 196 & \\ \hline & & & & $\mathrm{HIP}+\mathrm{HT}+$ shot peen & 1080 & 1350 & 7.5 & 220 & \\ \hline \multirow[t]{10}{*}{LPBF} & IN718 & & Strain rate $0.10 \% / \mathrm{s}$ & $0^{\circ} / \mathrm{AB}$ & $\sim 789.19$ & 1075.09 & 31.16 & 200.38 & [289] \\ \hline & & & & $90^{\circ} / \mathrm{AB}$ & 627.03 & 995.22 & 36.05 & 204.19 & \\ \hline & & & & $0^{\circ} / \mathrm{DA}$ & 1356.76 & 1523.55 & 15.24 & 203.24 & \\ \hline & & & & $90^{\circ} / \mathrm{DA}$ & 1194.59 & 1425.26 & 15.51 & 205.15 & \\ \hline & & & & $0^{\circ} / \mathrm{SA}$ & 1232.43 & 1492.83 & 18.91 & 211.83 & \\ \hline & & & & $90^{\circ} / \mathrm{SA}$ & 1167.57 & 1400.68 & 23.67 & 197.52 & \\ \hline & & & & $0^{\circ} / \mathrm{HA}$ & 1248.65 & 1449.83 & 19.73 & 220.42 & \\ \hline & & & & $90^{\circ} / \mathrm{HA}$ & 1178.38 & 1376.11 & 25.31 & 208.97 & \\ \hline & & & & $0^{\circ} / \mathrm{HSA}$ & 1248.65 & 1443.68 & 19.86 & 205.15 & \\ \hline & & & & $90^{\circ} / \mathrm{HSA}$ & 1200 & 1382.25 & 23.95 & 205.15 & \\ \hline \multirow[t]{4}{*}{LPBF} & WC/IN718 & & & Scan speed $400 \mathrm{~mm} / \mathrm{s}$ & & 1299.6 & 22.12 & & [65] \\ \hline & & & & $500 \mathrm{~mm} / \mathrm{s}$ & & $\sim 1339.94$ & $\sim 20.94$ & & \\ \hline & & & & $600 \mathrm{~mm} / \mathrm{s}$ & & $\sim 1408.84$ & $\sim 20.50$ & & \\ \hline & & & & $700 \mathrm{~mm} / \mathrm{s}$ & & 1464.6 & 19.74 & & \\ \hline \multirow[t]{2}{*}{LPBF} & IN718 & & Strain rate $1 \mathrm{~mm} / \mathrm{min}$ & $0^{\circ} / \mathrm{AB}$ & $912 \pm 15.2$ & $1072 \pm 38.6$ & $11.27 \pm 2.14$ & & [299] \\ \hline & & & & $90^{\circ} / \mathrm{AB}$ & $1102 \pm 34.5$ & $889 \pm 20.5$ & $30.42 \pm 1.68$ & & \\ \hline \multirow[t]{12}{*}{LPBF} & IN718 & & & $90^{\circ} / \mathrm{HIP}+\mathrm{SA} / 31.75 \mathrm{~J} / \mathrm{mm}^{3}$ & $\sim 1128.18$ & $\sim 1340.16$ & $\sim 19.82$ & & [300] \\ \hline & & & & $90^{\circ} / \mathrm{HIP}+\mathrm{SA} / 39.7 \mathrm{~J} / \mathrm{mm}^{3}$ & $\sim 1085.78$ & $\sim 1376.00$ & $\sim 19.96$ & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}+\mathrm{SA} / 40.8 \mathrm{~J} / \mathrm{mm}^{3}$ & $\sim 1100.99$ & $\sim 1373.54$ & $\sim 20.10$ & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}+\mathrm{SA} / 47.6 \mathrm{~J} / \mathrm{mm}^{3}$ & $\sim 1124.13$ & $\sim 1376.50$ & $\sim 19.71$ & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}+\mathrm{SA} / 51 \mathrm{~J} / \mathrm{mm}^{3}$ & $\sim 1096.59$ & $\sim 1354.00$ & $\sim 18.62$ & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}+\mathrm{SA} / 61.2 \mathrm{~J} / \mathrm{mm}^{3}$ & $\sim 1099.76$ & $\sim 1349.60$ & $\sim 17.61$ & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}+\mathrm{SA} / 31.75 \mathrm{~J} / \mathrm{mm}^{3}$ & $\sim 1137.78$ & $\sim 1389.21$ & $\sim 19.45$ & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}+\mathrm{SA} / 39.7 \mathrm{~J} / \mathrm{mm}^{3}$ & $\sim 1132.70$ & $\sim 1386.67$ & $\sim 23.58$ & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}+\mathrm{SA} / 40.8 \mathrm{~J} / \mathrm{mm}^{3}$ & $\sim 1140.32$ & $\sim 1394.29$ & $\sim 25.25$ & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}+\mathrm{SA} / 47.6 \mathrm{~J} / \mathrm{mm}^{3}$ & $\sim 1132.70$ & $\sim 1386.67$ & $\sim 23.63$ & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}+\mathrm{SA} / 51 \mathrm{~J} / \mathrm{mm}^{3}$ & $\sim 1114.92$ & $\sim 1373.97$ & $\sim 18.26$ & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}+\mathrm{SA} / 61.2 \mathrm{~J} / \mathrm{mm}^{3}$ & $\sim 1120$ & $\sim 1384.13$ & $\sim 23.25$ & & \\ \hline \multirow[t]{10}{*}{LPBF} & IN718 & ASTM E8 & Strain rate $0.13 \mathrm{~cm} / \mathrm{min}$ & HT \# 554 & $698.2 \pm 15.2$ & $995.2 \pm 12.8$ & $33.21 \pm 1.10$ & & [301] \\ \hline & & & & HT \# 528 & $1204.1 \pm 8.6$ & $1392 \pm 8.9$ & $17.32 \pm 0.71$ & & \\ \hline & & & & HT \# 527 & $1268.5 \pm 27.0$ & $1739.5 \pm 17.7$ & $15.44 \pm 2.00$ & & \\ \hline & & & & HT \# 522 & $1237.8 \pm 13.4$ & $1379.3 \pm 10.4$ & $19.49 \pm 0.54$ & & \\ \hline & & & & HT \# 553 & $859.5 \pm 22.9$ & $1171.4 \pm 12.8$ & $34.34 \pm 1.52$ & & \\ \hline & & & & нT \# 515 & $1124.4 \pm 18.9$ & $1330.8 \pm 21.4$ & $21.34 \pm 0.80$ & & \\ \hline & & & & HT \# 514 & $1200.6 \pm 9.5$ & $1330.8 \pm 21.4$ & $20.78 \pm 0.25$ & & \\ \hline & & & & НТ \# 509 & $1203.3 \pm 5.5$ & $1390.2 \pm 8.1$ & $21.96 \pm 0.37$ & & \\ \hline & & & & HT \# 507 & $1087.2 \pm 7.5$ & $1384.7 \pm 6.2$ & $23.36 \pm 0.62$ & & \\ \hline & & & & HT \# 506 & $1110.9 \pm 7.4$ & $1395.7 \pm 4.2$ & $23.61 \pm 0.36$ & & \\ \hline \multirow[t]{4}{*}{LPBF} & IN718 & ASTM E8M & Strain rate $10^{-4} \mathrm{~s}^{-1}$ & No WC & $771 \pm 4$ & $1073 \pm 1$ & $29.4 \pm 0.6$ & $159 \pm 12$ & [302] \\ \hline & & & & $+5 \% \mathrm{WC}$ & $980 \pm 16$ & $1195 \pm 12$ & $21.6 \pm 1.5$ & $215 \pm 17$ & \\ \hline & & & & $+10 \% \mathrm{WC}$ & $1078 \pm 14$ & $1287 \pm 15$ & $15.1 \pm 1.1$ & $223 \pm 15$ & \\ \hline & & & & $+15 \% \mathrm{WC}$ & $974 \pm 21$ & $1104 \pm 19$ & $7.8 \pm 1.1$ & $225 \pm 21$ & \\ \hline \multirow[t]{9}{*}{LPBF} & IN718 & ASTM E8/E8M and ASTM E21 & Strain rate $0.005 / \mathrm{min}$ & Sample orientation: XYZ/RT & 694.77 & 1007.04 & 28.64 & 134.61 & [306] \\ \hline & & & & $\mathrm{YZX} / \mathrm{RT}$ & 662.90 & 1004.21 & 27.92 & 127.01 & \\ \hline & & & & $\mathrm{ZXY} / \mathrm{RT}$ & 604.43 & 912.63 & 33.63 & 82.16 & \\ \hline & & & & $\mathrm{XYZ} / 200^{\circ} \mathrm{C}$ & 702.82 & 1016.54 & 25.08 & 120.70 & \\ \hline & & & & $\mathrm{YZX} / 200^{\circ} \mathrm{C}$ & 634.35 & 968.86 & 24.51 & 121.23 & \\ \hline & & & & $\mathrm{ZXY} / 200^{\circ} \mathrm{C}$ & 539.90 & 850.76 & 31.43 & 90.27 & \\ \hline & & & & $\mathrm{XYZ} / 350^{\circ} \mathrm{C}$ & 667.97 & 955.15 & 26.45 & 108.44 & \\ \hline & & & & $\mathrm{YZX} / 350^{\circ} \mathrm{C}$ & 615.30 & 931.03 & 24.98 & 121.64 & \\ \hline & & & & $\mathrm{ZXY} / 350^{\circ} \mathrm{C}$ & 517.93 & 786.57 & 31.39 & 70.03 & \\ \hline \multirow[t]{4}{*}{LPBF} & IN718 + Nano TiC & & Simulation & $225 \mathrm{~J} / \mathrm{m}$ & & & & 117.506 & [307] \\ \hline & & & & $250 \mathrm{~J} / \mathrm{m}$ & & & & 120.824 & \\ \hline & & & & $275 \mathrm{~J} / \mathrm{m}$ & & & & 207.491 & \\ \hline & & & & $300 \mathrm{~J} / \mathrm{m}$ & & & & 225.402 & \\ \hline \multirow[t]{5}{*}{LPBF} & IN718 & ISO 6892-2 & Loading rate $1.5 \mathrm{~mm} / \mathrm{min} 250 \mathrm{~W} /$ & $668 \pm 16$ & $1011 \pm 27$ & $22 \pm 2$ & $173 \pm 13$ & & $[67]$ \\ \hline & & & $\mathrm{AB}$ & $950 \mathrm{~W} / \mathrm{AB}$ & $531 \pm 9$ & $866 \pm 33$ & $21 \pm 5$ & $113 \pm 3$ & \\ \hline & & & & $250 \mathrm{~W} / \mathrm{HT}$ & $875 \pm 11$ & $1153 \pm 4$ & $17 \pm 2$ & $190 \pm 11$ & \\ \hline & & & & $950 \mathrm{~W} / \mathrm{HT}$ & $668 \pm 7$ & $884 \pm 80$ & $7 \pm 2$ & $138 \pm 5$ & \\ \hline & & & & $250 \mathrm{~W} / \mathrm{HIP}$ & $645 \pm 6$ & $1025 \pm 14$ & $38 \pm 1$ & $188 \pm 8$ & \\ \hline \end{tabular} \end{center} \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|c|} \hline \multirow[t]{2}{*}{Technique} & Material & \multirow[t]{2}{*}{Standard} & Test condition & Sample Condition & Yield strength (MPa) & \begin{tabular}{l} Tensile strength \\ (MPa) \\ \end{tabular} & \multicolumn{2}{|c|}{Elongation (\%)} & \begin{tabular}{l} Young's modulus \\ (GPa) \\ \end{tabular} & Ref \\ \hline & & & & $950 \mathrm{~W} / \mathrm{HIP}$ & $481 \pm 11$ & $788 \pm 12$ & $34 \pm 3$ & & $183 \pm 19$ & \\ \hline & IN718 zone 1-250 W Matrix; & & Strain rate $2 \mathrm{~mm} / \mathrm{min}$ & $250 \mathrm{~W} / \mathrm{HIP}+\mathrm{HT}$ & $1145 \pm 16$ & $1376 \pm 14$ & $19 \pm 1$ & & $190 \pm 6$ & \\ \hline & zone 2 - two lines of $950 \mathrm{~W}$ & & & $950 \mathrm{~W} / \mathrm{HIP}+\mathrm{HT}$ & $1065 \pm 20$ & $1272 \pm 12$ & $15 \pm 4$ & & $188 \pm 20$ & \\ \hline & & & & $\mathrm{AB}$ & $574 \pm 6$ & $873 \pm 14$ & $13 \pm 2$ & & $136 \pm 13$ & \\ \hline & & & & HT & $704 \pm 8$ & $920 \pm 53$ & $4 \pm 2$ & & $167 \pm 13$ & \\ \hline & & & & HIP & $500 \pm 6$ & $817 \pm 16$ & $21 \pm 1$ & & $187 \pm 8$ & \\ \hline & & & & HIP + HT & $1041 \pm 47$ & $1154 \pm 68$ & $7 \pm 1$ & & $196 \pm 15$ & \\ \hline & IN718 Cylindrical samples & & & $250 \mathrm{~W} / \mathrm{AB} / 650^{\circ} \mathrm{C}$ & $650 \pm 11$ & $845 \pm 9$ & $28 \pm 4$ & & & \\ \hline & & & & $950 \mathrm{~W} / \mathrm{AB} / 650^{\circ} \mathrm{C}$ & $543 \pm 2$ & $782 \pm 6$ & $31 \pm 6$ & & & \\ \hline & & & & $250 \mathrm{~W} / \mathrm{HIP} / 650^{\circ} \mathrm{C}$ & $626 \pm 8$ & $857 \pm 14$ & $29 \pm 1$ & & & \\ \hline & & & & $950 \mathrm{~W} / \mathrm{HIP} / 650^{\circ} \mathrm{C}$ & $479 \pm 5$ & $665 \pm 7$ & $28 \pm 2$ & & & \\ \hline & & & & $250 \mathrm{~W} / \mathrm{HIP}+\mathrm{HT} / 650^{\circ} \mathrm{C}$ & $942 \pm 11$ & $1078 \pm 8$ & $20 \pm 2$ & & & \\ \hline & & & & $950 \mathrm{~W} / \mathrm{HIP}+\mathrm{HT} / 650^{\circ} \mathrm{C}$ & $872 \pm 13$ & $1005 \pm 12$ & $17 \pm 4$ & & & \\ \hline \multirow[t]{5}{*}{LPBF} & IN718 & & Strain rate $\sim 10^{-3} \mathrm{~s}^{-1}$ & $90^{\circ} / \mathrm{HIP}+$ annealed $/ \mathrm{Ar}$ & 850 & 1140 & 28 & & & [64] \\ \hline & & & & $0^{\circ} / \mathrm{HIP}+$ annealed $/ \mathrm{Ar}$ & 890 & 1200 & 28 & & & \\ \hline & & & & $0^{\circ} / \mathrm{AB} / \mathrm{N} 2$ & 830 & 1120 & 25 & & & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}+$ annealed $/ \mathrm{N} 2$ & 880 & 1140 & 30 & & & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}+$ annealed $/ \mathrm{N} 2$ & 930 & 1200 & 27 & & & \\ \hline \multirow[t]{12}{*}{LPBF} & IN718 Lattice structure & & & Body Centred Cubic (BCC) $2 \mathrm{~mm}$ & 7.28 & & 27 & & 354.61 & [309] \\ \hline & & & & BCC $3 \mathrm{~mm}$ & 1.56 & & & & 56.73 & \\ \hline & & & & BCC $4 \mathrm{~mm}$ & 0.52 & & & & 26.50 & \\ \hline & & & & Body Centred Cubic-Z Struts (BCCZ) $2 \mathrm{~mm}$ & 16.21 & & & & 859.47 & \\ \hline & & & & BCCZ $3 \mathrm{~mm}$ & 4.13 & & & & 522.85 & \\ \hline & & & & BCCZ $4 \mathrm{~mm}$ & 1.83 & & & & 276.58 & \\ \hline & & & & Face Centred Cubic (FCC) $2 \mathrm{~mm}$ & 8.97 & & & & 610.71 & \\ \hline & & & & FCC $3 \mathrm{~mm}$ & 2.16 & & & & 150.92 & \\ \hline & & & & FCC $4 \mathrm{~mm}$ & 0.86 & & & & 54.14 & \\ \hline & & & & Face Centred Cubic-Z Struts (FCCZ) $2 \mathrm{~mm}$ & 16.21 & & & & 1267.26 & \\ \hline & & & & FCCZ $3 \mathrm{~mm}$ & 4.86 & & & & 639.84 & \\ \hline & & & & FCCZ $4 \mathrm{~mm}$ & 2.16 & & & & 365.80 & \\ \hline \multirow[t]{5}{*}{LPBF} & IN718 & ASTM-E8/E8M & Strain rate $0.01 \mathrm{~s}^{-1}$ & Nominal & 973.82 & 1265.84 & 18.01 & & & [320] \\ \hline & & & & Increased hatch spacing & 956.46 & 1256.55 & 12.29 & & & \\ \hline & & & & Decreased hatch spacing & 904.85 & 1229.98 & 20.94 & & & \\ \hline & & & & Increased cooling rate & 910.36 & 1233.97 & 20.05 & & & \\ \hline & & & & Decreased cooling rate & 940.98 & 1253.89 & 21.29 & & & \\ \hline LPBF & IN718 & & & Material A/90 $90^{\circ} / \mathrm{HT}$ & 1284 & 1432 & 29.2 & \begin{tabular}{l} $29.7 \%$ \\ $\mathrm{RA}$ \\ \end{tabular} & & [313] \\ \hline & & & & Material $\mathrm{A} / 0^{\circ} / \mathrm{HT}$ & 1329 & 1499 & 31.8 & \begin{tabular}{l} $30.7 \%$ \\ $\mathrm{RA}$ \\ \end{tabular} & & \\ \hline & & & & Material $\mathrm{B} / 90^{\circ} / \mathrm{HT}$ & 1227 & 1366 & 13.6 & $8.6 \%$ & & \\ \hline & & & & & & & & RA & & \\ \hline & & & & Material $\mathrm{B} / 0^{\circ} / \mathrm{HT}$ & 1300 & 1467 & 27.3 & 20\%RA & & \\ \hline LPBF & IN718 & & Strain rate $5 \times 10^{-4} \mathrm{~s}^{-1}$ & $\mathrm{AB}$ & $\sim 559.82$ & $\sim 781.96$ & & & & [322] \\ \hline & & & & SA & $\sim 1016.63$ & $\sim 1147.41$ & & & & \\ \hline & & & & НА & $\sim 986.18$ & $\sim 1163.53$ & & & & \\ \hline & & & & HSA & $\sim 1082.92$ & $\sim 1152.78$ & & & & \\ \hline LPBF & IN718 & & & AB & & 1021-1035 & $31-34$ & & & [335] \\ \hline & & & & STA & & 1428 & 14 & & & \\ \hline LPBF & IN718 & & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & Homogenisation + AMS 5663 STA & $1211 \pm 24$ & $1406 \pm 21$ & $13.6 \pm 4$ & & $191 \pm 3.1$ & [337] \\ \hline LPBF & IN718 (AMS 5662) & & Strain rate $8.3 \times 10^{-3} \mathrm{~s}^{-1}$ & Test at $20^{\circ} \mathrm{C}$ & 1175 & 1364 & $20.9 \pm 0.5$ & & & [339] \\ \hline & & & & Test at $550^{\circ} \mathrm{C}$ & 1064 & 1176 & $16.9 \pm 0.3$ & & & \\ \hline & & & & Test at $600^{\circ} \mathrm{C}$ & 1042 & 1185 & $16.0 \pm 0.7$ & & & \\ \hline & & & & Test at $6500^{\circ} \mathrm{C}$ & 1025 & 1126 & $19.5 \pm 1.0$ & & & \\ \hline & & & & Test at $700^{\circ} \mathrm{C}$ & 957 & 1011 & $19.3 \pm 1.3$ & & & \\ \hline LPBF & IN718 & Compressive test & Strain rate $0.005 \mathrm{~mm} / \mathrm{min}$ & \begin{tabular}{l} Island $/ 30^{\circ}$ interlayer rotation $/ 500 \mathrm{~mm} / \mathrm{s}$ \\ scan speed \\ \end{tabular} & 595 & & & & & [148] \\ \hline & & & & \begin{tabular}{l} scan speed \\ Island $/ 30^{\circ}$ interlayer rotation $/ 700 \mathrm{~mm} / \mathrm{s}$ \\ \end{tabular} & 580 & & & & & \\ \hline & & & & \begin{tabular}{l} scan speed \\ Island $/ 30^{\circ}$ interlayer rotation $/ 1000 \mathrm{~mm} / \mathrm{s}$ \\ \end{tabular} & 580 & & & & & \\ \hline & & & & scan speed & & & & & & \\ \hline & & & & Meander $/ 90^{\circ}$ interlayer rotation $/ 500 \mathrm{~mm} / \mathrm{s}$ & 680 & & & & & \\ \hline & & & & \begin{tabular}{l} scan speed \\ Meander $/ 90^{\circ}$ interlayer rotation $/ 700 \mathrm{~mm} / \mathrm{s}$ \\ \end{tabular} & 660 & & & & & \\ \hline & & & & \begin{tabular}{l} scan speed \\ Meander $/ 90^{\circ}$ interlayer rotation $1000 \mathrm{~mm} /$ \\ \end{tabular} & & & & & & \\ \hline & & & & \begin{tabular}{l} Meander $/ 90^{\circ}$ interlayer rotation $/ 1000 \mathrm{~mm} /$ \\ s scan speed \\ \end{tabular} & & & & & & \\ \hline \end{tabular} \end{center} Table 9 (continued) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-53} \end{center} \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample Condition & Yield strength (MPa) & \begin{tabular}{l} Tensile strength \\ (MPa) \\ \end{tabular} & \multicolumn{2}{|c|}{Elongation (\%)} & \begin{tabular}{l} Young's modulus \\ (GPa) \\ \end{tabular} & Ref \\ \hline & & & & & & & & & 225 & \\ \hline & & & & $\mathrm{XY}$ (triple) $/ \mathrm{HT} / 23^{\circ} \mathrm{C}$ & & & & & \begin{tabular}{l} 210 \\ 197 \\ \end{tabular} & \\ \hline & & & & & & & & & 228 & \\ \hline & & & & & & & & & 204 & \\ \hline & & & & $\mathrm{XY}$ (single)/HT $/ 850^{\circ} \mathrm{C}$ & & & & & 141 & \\ \hline & & & & & & & & & \begin{tabular}{l} 164 \\ 154 \\ \end{tabular} & \\ \hline & & & & $\mathrm{Z} / \mathrm{AB} / 23^{\circ} \mathrm{C}$ & & & & & \begin{tabular}{l} 154 \\ 141 \\ \end{tabular} & \\ \hline & & & & & & & & & 138 & \\ \hline & & & & $\mathrm{Z} / \mathrm{HT} / 23^{\circ} \mathrm{C}$ & & & & & 196 & \\ \hline & & & & & & & & & 200 & \\ \hline & & & & $\mathrm{Z} / \mathrm{HT} / 850^{\circ} \mathrm{C}$ & & & & & 143 & \\ \hline \multirow[t]{3}{*}{LPBF} & IN738LC & EN 10002 & & $90^{\circ}$ & $765 \pm 10$ & & & & $141 \pm 3$ & [276] \\ \hline & & & & $0^{\circ}$ & $853 \pm 16$ & & & & $141 \pm 4$ & \\ \hline & & & & $\mathrm{H} 45^{\circ}$ & $893 \pm 4$ & & & & $215 \pm 7$ & \\ \hline \multirow[t]{3}{*}{LPBF} & IN738LC & ASTM E8 & Strain rate $4 \% / \mathrm{min}$ & $\mathrm{AB}$ & 895 & 1010 & $1.6 \pm 0.2$ & & & [323] \\ \hline & & & & HIP & & 1010 & & & & \\ \hline & & & & \begin{tabular}{l} $\mathrm{HIP}+\mathrm{DA} / \mathrm{RT}$ \\ $\mathrm{HIP}+\mathrm{DA} / 850^{\circ} \mathrm{C}$ \\ \end{tabular} & & $720+1$ & $14.4+1$ & & & \\ \hline LPBF & IN939 & & & \begin{tabular}{l} HIP + DA $/ 850^{\circ} \mathrm{C}$ \\ Various HTs \\ \end{tabular} & \begin{tabular}{l} $560 \pm 1$ \\ See Fig. 7 in the origina \\ \end{tabular} & \begin{tabular}{l} $720 \pm 1$ \\ 1 reference paper \\ \end{tabular} & $14.4 \pm 1$ & & & [253] \\ \hline \multirow[t]{6}{*}{LPBF} & Invar 36 & & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & $0^{\circ} / 3200 \mathrm{~mm} / \mathrm{s}$ & $\sim 397.7$ & $\sim 509.7$ & $\sim 15.3$ & & & [260] \\ \hline & & & & $90^{\circ} / 3200 \mathrm{~mm} / \mathrm{s}$ & $\sim 352.5$ & $\sim 381.7$ & $\sim 3.25$ & & & \\ \hline & & & & $90^{\circ} / 3200 \mathrm{~mm} / \mathrm{s} / \mathrm{HIP}$ & $\sim 288.3$ & $\sim 346.3$ & $\sim 4.7$ & & & \\ \hline & & & & $90^{\circ} / 1000 \mathrm{~mm} / \mathrm{s}$ & $\sim 342.5$ & $\sim 453$ & $\sim 29.5$ & & & \\ \hline & & & & $0^{\circ} / 1000 \mathrm{~mm} / \mathrm{s}$ & $\sim 400$ & $\sim 536.5$ & $\sim 19.5$ & & & \\ \hline & & & & $0^{\circ} / 1000 \mathrm{~mm} / \mathrm{s} / \mathrm{HT}$ & $\sim 318$ & $\sim 443.5$ & $\sim 30.5$ & & & \\ \hline \multirow[t]{2}{*}{LPBF} & K418 & & Strain rate $1 \mathrm{~mm} / \mathrm{min}$ & No graphene nanoplatelts (GNPs) & 912 & 1200 & & & & [338] \\ \hline & & & & $10.1 \mathrm{wt} \%$ GNPs & 1018 & 1078 & & & & \\ \hline \multirow[t]{10}{*}{LPBF} & к536 & ASTM E8/E8M and ASTM E21 & Strain rate $0.005 / \mathrm{min}$ & $0^{\circ} / \mathrm{RT} / \mathrm{SR}$ & $338 \pm 1.7$ & $800 \pm 2.1$ & $36.9 \pm 0.1$ & $38.4 \%$ & $193.3 \pm 3.8$ & [319] \\ \hline & & & & $90^{\circ} / \mathrm{RT} / \mathrm{SR}$ & $325 \pm 2.5$ & $775 \pm 1.2$ & $41.3 \pm 0.7$ & $46.3 \%$ & $188.7 \pm 1.2$ & \\ \hline & & & & $0^{\circ} / 400^{\circ} \mathrm{C} / \mathrm{SR}$ & $250 \pm 17.9$ & $683 \pm 6.7$ & $40.7 \pm 0.9$ & \begin{tabular}{l} RA \\ $41.4 \%$ \\ RA \\ \end{tabular} & $143.8 \pm 21.4$ & \\ \hline & & & & $90^{\circ} / 400^{\circ} \mathrm{C} / \mathrm{SR}$ & $243 \pm 13.1$ & $667 \pm 0.7$ & $45.2 \pm 0.8$ & \begin{tabular}{l} $45.7 \%$ \\ $\mathrm{RA}$ \\ \end{tabular} & $129.1 \pm 31.1$ & \\ \hline & & & & $0^{\circ} / 500^{\circ} \mathrm{C} / \mathrm{SR}$ & $249 \pm 2.6$ & $670 \pm 2.6$ & $40.3 \pm 2.0$ & \begin{tabular}{l} $40.9 \%$ \\ RA \\ \end{tabular} & $118.3 \pm 11.5$ & \\ \hline & & & & $90^{\circ} / 500^{\circ} \mathrm{C} / \mathrm{SR}$ & $237 \pm 5.0$ & $657 \pm 4.0$ & $43.3 \pm 0.7$ & \begin{tabular}{l} $47.3 \%$ \\ RA \\ \end{tabular} & $114.4 \pm 7.1$ & \\ \hline & & & & $0^{\circ} / 600^{\circ} \mathrm{C} / \mathrm{SR}$ & $242 \pm 4.5$ & $635 \pm 2.1$ & $41.6 \pm 0.8$ & \begin{tabular}{l} $42.4 \%$ \\ $\mathrm{RA}$ \\ \end{tabular} & $117.2 \pm 12.6$ & \\ \hline & & & & $90^{\circ} / 600^{\circ} \mathrm{C} / \mathrm{SR}$ & $228 \pm 4.6$ & $616 \pm 0.7$ & $44.6 \pm 2.8$ & \begin{tabular}{l} $46.0 \%$ \\ RA \\ \end{tabular} & $103.9 \pm 1.4$ & \\ \hline & & & & $0^{\circ} / 700^{\circ} \mathrm{C} / \mathrm{SR}$ & $218 \pm 1.7$ & $482 \pm 3.2$ & $27.2 \pm 0.4$ & \begin{tabular}{l} $26.7 \%$ \\ $\mathrm{RA}$ \\ \end{tabular} & $129.8 \pm 27.9$ & \\ \hline & & & & $90^{\circ} / 700^{\circ} \mathrm{C} / \mathrm{SR}$ & $213 \pm 4.4$ & $463 \pm 3.1$ & $32.0 \pm 1.8$ & \begin{tabular}{l} $32.4 \%$ \\ $\mathrm{RA}$ \\ \end{tabular} & $104.4 \pm 30.9$ & \\ \hline \multirow[t]{10}{*}{LPBF} & к536 & & Strain rate $0.005 \mathrm{~mm} / \mathrm{min}$ & $0^{\circ} / \mathrm{RT}$ & $338 \pm 1.7$ & $800 \pm 2.1$ & $36.9 \pm 0.1$ & & $193 \pm 3.8$ & [319] \\ \hline & & & & $90^{\circ} / \mathrm{RT}$ & $325 \pm 2.5$ & $775 \pm 1.2$ & $41.3 \pm 0.7$ & & $188.7 \pm 1.2$ & \\ \hline & & & & $0^{\circ} / 400^{\circ} \mathrm{C}$ & $250 \pm 17.9$ & $683 \pm 6.7$ & $40.7 \pm 0.9$ & & $143.8 \pm 21.4$ & \\ \hline & & & & $90^{\circ} / 400^{\circ} \mathrm{C}$ & $243 \pm 13.1$ & $667 \pm 0.7$ & $45.2 \pm 0.8$ & & $129.1 \pm 31.1$ & \\ \hline & & & & $0^{\circ} / 500^{\circ} \mathrm{C}$ & $249 \pm 2.6$ & $670 \pm 2.6$ & $40.3 \pm 2$ & & $118.3 \pm 11.5$ & \\ \hline & & & & $90^{\circ} / 500^{\circ} \mathrm{C}$ & $237 \pm 5$ & $657 \pm 4$ & $43.3 \pm 0.7$ & & $114.4 \pm 7.1$ & \\ \hline & & & & $0^{\circ} / 600^{\circ} \mathrm{C}$ & $242 \pm 4.5$ & $635 \pm 2.2$ & $41.6 \pm 0.8$ & & $117.2 \pm 12.6$ & \\ \hline & & & & $90^{\circ} / 600^{\circ} \mathrm{C}$ & $228 \pm 4.6$ & $616 \pm 0.7$ & $44.6 \pm 2.8$ & & $103.9 \pm 1.4$ & \\ \hline & & & & $0^{\circ} / 700^{\circ} \mathrm{C}$ & $218 \pm 1.7$ & $482 \pm 3.2$ & $27.2 \pm 0.4$ & & $129.8 \pm 27.9$ & \\ \hline & & & & $90^{\circ} / 700^{\circ} \mathrm{C}$ & $213 \pm 4.4$ & $463 \pm 3.1$ & $32 \pm 1.8$ & & $104.4 \pm 30.9$ & \\ \hline \multirow[t]{5}{*}{LPBF} & Nimonic 263 & French Aeronautical standard & & $90^{\circ} / \mathrm{AB}$ & $818 \pm 8$ & $1085 \pm 11$ & $24 \pm 24$ & & 163 & [285] \\ \hline & & CEAT TP5 & & $0^{\circ} / \mathrm{AB}$ & $653 \pm 11$ & $860 \pm 8$ & $70 \pm 1$ & & 191 & \\ \hline & & & & $90^{\circ} / \mathrm{DA}$ & $834 \pm 13$ & $1136 \pm 26$ & $29 \pm 2$ & & 150 & \\ \hline & & & & $0^{\circ} / \mathrm{DA}$ & $697 \pm 16$ & $910 \pm 4$ & $52 \pm 12$ & & 142 & \\ \hline & & & & $90^{\circ} / \mathrm{SA}$ & $843 \pm 20$ & $1268 \pm 7$ & $29 \pm 3$ & & 199 & \\ \hline \end{tabular} \end{center} Table 9 (continued) \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample Condition & Yield strength (MPa) & \begin{tabular}{l} Tensile strength \\ (MPa) \\ \end{tabular} & Elongation (\%) & \begin{tabular}{l} Young's modulus \\ (GPa) \\ \end{tabular} & Ref \\ \hline LPBF & \begin{tabular}{l} Oxide dis \\ nickel-ba \\ \end{tabular} & & Strain rate $10^{-3}$ & \begin{tabular}{l} $0^{\circ} / \mathrm{SA}$ \\ Different process parameters \\ \end{tabular} & \begin{tabular}{l} $709 \pm 7$ \\ See Table 4 in the oris \\ \end{tabular} & \begin{tabular}{l} $981 \pm 4$ \\ nal reference paper \\ \end{tabular} & $53 \pm 2$ & 206 & [336] \\ \hline \end{tabular} \end{center} LPBF $\quad$ nickel-based superalloy $\begin{array}{ll}\text { LPBF } & \text { Steel/N } \\ \text { EBM } & \text { IN625 } \\ \end{array}$ EBM and Binder Jetting IN625 Евм IN625 Mesh Compression tes Strain rate $10^{-3} \mathrm{~s}^{-1}$ Temperatu Strain rate $\sim 10^{-3} \mathrm{~s}^{-1}$ See Fig. 4 in the original reference paper\\ $90^{\circ}$\\ $90^{\circ} / \mathrm{HIP}$\\ $\mathrm{EBM}, 0^{\circ}$\\ $\mathrm{EBM}, 90^{\circ}$\\ Binder jetting, $0^{\circ}$\\ Binder jetting, $90^{\circ}$\\ Load parallel or perpendicular to the build\\ direction/specimens with different density\\ $\mathrm{AB} / \mathrm{RT}$\\ $\mathrm{HIT} / \mathrm{RT}$\\ $\mathrm{AB} / 538^{\circ} \mathrm{C}$\\ $\mathrm{HIP} / 538^{\circ} \mathrm{C}$ 4\\ 3\\ 367\\ 369\\ 3\\ 3\\ $\mathrm{~S}$\\ 4\\ 4\\ 3\\ 3\\ 3\\ 2\\ 2 ЕBM IN625 EBM ЕвМ\\ IN718\\ IN718 IN718 \begin{center} \begin{tabular}{|c|c|c|c|c|c|} \hline \multirow{2}{*}{}\begin{tabular}{l} 410 \\ 330 \\ \end{tabular} & \multirow{2}{*}{}\begin{tabular}{l} 750 \\ 770 \\ \end{tabular} & \multicolumn{3}{|l|}{44} & \multirow[t]{2}{*}{[264]} \\ \hline & & \multicolumn{3}{|l|}{69} & \\ \hline $367 \pm 33$ & $849 \pm 37$ & \multicolumn{2}{|c|}{$44.32 \pm 4.95$} & $484 \pm 52$ & $[100]$ \\ \hline $369 \pm 7$ & $723 \pm 29$ & \multicolumn{2}{|c|}{$26.92 \pm 5.49$} & $459 \pm 36$ & \\ \hline $320 \pm 14$ & $707 \pm 12$ & \multicolumn{2}{|c|}{\multirow{2}{*}}{}\{\begin{tabular}{l} $58.74 \pm 2.14$ \\ $27.02 \pm 5.39$ \\ \end{tabular}\} & $524 \pm 47$ & \\ \hline $393 \pm 2$ & $708 \pm 22$ & & & $506 \pm 51$ & \\ \hline \multicolumn{5}{|c|}{See the original reference paper} & [318] \\ \hline 410 & 750 & \multicolumn{3}{|l|}{44} & [292] \\ \hline 330 & 770 & \multicolumn{3}{|l|}{69} & \\ \hline 300 & 590 & \multicolumn{4}{|l|}{53} \\ \hline \multirow[t]{4}{*}{230} & 610 & \multicolumn{3}{|l|}{70} & \\ \hline & & & & 0.76 & [293] \\ \hline & & & & 1.68 & \\ \hline & & & & 4.17 & \\ \hline 793 & $952 \pm 18$ & & & & [305] \\ \hline $527 \pm 19$ & $670 \pm 44.5$ & \multicolumn{2}{|l|}{$21 \pm 2.0$} & & $[310]$ \\ \hline $527 \pm 19$ & $670 \pm 44.5$ & \multicolumn{2}{|l|}{$21 \pm 2$} & & [330] \\ \hline \multirow{3}{*}{$377 \pm 39$} & $603 \pm 34$ & \multicolumn{2}{|c|}{\multirow{3}{*}}{$23 \pm 8$} & & \\ \hline & & & & $\sim 100$ & [291] \\ \hline & & & & $\sim 142$ & \\ \hline $\sim 923.91$ & 1113.05 & \multicolumn{2}{|l|}{31.51} & 98.69 & [294] \\ \hline 1128.26 & 1268.98 & & & 108.03 & \\ \hline 1132.61 & 1218.36 & 24.89 & & 106.28 & \\ \hline 1089.13 & 1180.83 & 22.08 & & 107.15 & \\ \hline 1119.57 & 1201.49 & 28.10 & & 108.32 & \\ \hline 771.74 & 1002.51 & 40.35 & & 144.23 & \\ \hline 1041.30 & 1200.62 & 31.31 & & 135.18 & \\ \hline 941.30 & 1029.27 & 14.05 & & 138.10 & \\ \hline 867.39 & 1095.02 & 38.34 & & 124.67 & \\ \hline 934.78 & 1073.49 & 35.33 & & 141.02 & \\ \hline $568 \pm 5$ & $818 \pm 43$ & $16.9 \pm 2.9$ & & & [183] \\ \hline Premature & & & & & \\ \hline 590.13 & 941.76 & 34.3 & & 151.68 & [97] \\ \hline 868.87 & 1108.37 & 22.1 & & 149.82 & \\ \hline $822 \pm 25$ & $1060 \pm 26$ & $22 \% \mathrm{EL}$ & \begin{tabular}{l} $25 \%$ \\ $\mathrm{RA}$ \\ \end{tabular} & $192 \pm 11$ & [295] \\ \hline $744 \pm 44$ & $929 \pm 20$ & $5.5 \% \mathrm{EL}$ & \begin{tabular}{l} $12 \%$ \\ $\mathrm{RA}$ \\ \end{tabular} & $180 \pm 6$ & \\ \hline $1154 \pm 46$ & $1238 \pm 22$ & $7 \% \mathrm{EL}$ & \begin{tabular}{l} $14 \%$ \\ $\mathrm{RA}$ \\ \end{tabular} & $198 \pm 12$ & \\ \hline $1187 \pm 27$ & $1232 \pm 16$ & $1.1 \% \mathrm{EL}$ & $5 \% \mathrm{RA}$ & $198 \pm 8$ & \\ \hline $590 \pm 40$ & $942 \pm 61$ & $34 \pm 2.6$ & & & [118] \\ \hline $869 \pm 32$ & $1108 \pm 50$ & $22 \pm 1.8$ & & & \\ \hline $887 \pm 16$ & $1003 \pm 21$ & $5.4 \pm 1.7$ & & & \\ \hline $822 \pm 12$ & $1082 \pm 10$ & $20 \pm 0.6$ & & & \\ \hline $957 \pm 30$ & $1142 \pm 41$ & $19 \pm 4.6$ & & & \\ \hline $974 \pm 20$ & $1186 \pm 34$ & $20 \pm 1.3$ & & & \\ \hline $967 \pm 28$ & $1186 \pm 19$ & $20 \pm 1.4$ & & & \\ \hline $632 \pm 88$ & $1069 \pm 44$ & $17 \pm 2.4$ & & & \\ \hline 1009 & 1082 & 38 & & 88 & [296] \\ \hline 752 & 834 & 17 & & 122 & \\ \hline 834 & 1055 & 20 & & 174 & \\ \hline 827 & 1048 & 8.5 & & 177 & \\ \hline $925 \pm 20$ & $1138 \pm 24$ & $15.7 \pm 4.3$ & & & [297] \\ \hline $894 \pm 24$ & $1061 \pm 83$ & $11.5 \pm 6.9$ & & & \\ \hline $1061 \pm 16$ & $1266 \pm 44$ & $21.1 \pm 1.1$ & & & \\ \hline $1035 \pm 17$ & $1240 \pm 19$ & $21.8 \pm 2.4$ & & & \\ \hline \end{tabular} \end{center} \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample Condition & Yield strength (MPa) & \begin{tabular}{l} Tensile strength \\ (MPa) \\ \end{tabular} & Elongation (\%) & \begin{tabular}{l} Young's modulus \\ (GPa) \\ \end{tabular} & Ref \\ \hline EBM & \begin{tabular}{l} IN718 \\ \end{tabular} & ASTM E8/E8M and ASTM E21 & RT and $650^{\circ} \mathrm{C}$ & \begin{tabular}{l} Various build orientations and scan \\ strategies \\ \end{tabular} & \multicolumn{3}{|c|}{See Fig. 8 in the original reference paper} & & [298] \\ \hline \multirow[t]{6}{*}{EBM} & IN718 & & Strain rate $0.0056 \mathrm{~s}^{-1}$ & $\mathrm{AB} /$ parameter set 1 & 980 & 1160 & 8.2 & & [304] \\ \hline & & & & $\mathrm{AB} /$ parameter set 2 & 980 & 1160 & 8.2 & & \\ \hline & & & & $\mathrm{AB} /$ parameter set 3 & 980 & 1160 & 8.2 & & \\ \hline & & & & \begin{tabular}{l} HT/parameter set 1 \\ \end{tabular} & 1180 & 1350 & 6.5 & & \\ \hline & & & & HT/parameter set 2 & 1290 & 1440 & 7.1 & & \\ \hline & & & & HT/parameter set 3 & 1180 & 1350 & 7.1 & & \\ \hline \multirow[t]{22}{*}{EBM} & IN718 & & & Focus offset $1 \mathrm{~mA}$ at $25^{\circ} \mathrm{C}$ & $988.8 \pm 3.7$ & $1144.0 \pm 1.5$ & $31.5 \pm 4.3$ & & [316] \\ \hline & & & & $3 \mathrm{~mA}$ & $1010.5 \pm 5.5$ & $1157.5 \pm 5.5$ & $25.3 \pm 2.6$ & & \\ \hline & & & & $6 \mathrm{~mA}$ & $1050.0 \pm 5.5$ & $1187.3 \pm 8.4$ & $25.1 \pm 5.7$ & & \\ \hline & & & & $9 \mathrm{~mA}$ & $1122.3 \pm 17.4$ & $1300.6 \pm 22.5$ & $25.3 \pm 0.9$ & & \\ \hline & & & & $12 \mathrm{~mA}$ & $1095.1 \pm 32.1$ & $1278.4 \pm 26.8$ & $25.3 \pm 0.7$ & & \\ \hline & & & & $15 \mathrm{~mA}$ & $1112.3 \pm 35.6$ & $1276.6 \pm 14.2$ & $26.4 \pm 0.4$ & & \\ \hline & & & & $18 \mathrm{~mA}$ & $978.9 \pm 8.3$ & $1053.9 \pm 11.6$ & $9.3 \pm 1.1$ & & \\ \hline & & & & $21 \mathrm{~mA}$ & $989.8 \pm 6.1$ & $1046.3 \pm 17.1$ & $8.6 \pm 0.9$ & & \\ \hline & & & & $24 \mathrm{~mA}$ & $946.3 \pm 5.8$ & $958.7 \pm 20.9$ & $6.2 \pm 0.3$ & & \\ \hline & & & & $30 \mathrm{~mA}$ & $674.3 \pm 47.8$ & $680.5 \pm 38.9$ & $6.1 \pm 0.8$ & & \\ \hline & & & & $40 \mathrm{~mA}$ & $443.5 \pm 16.4$ & $455.4 \pm 6.1$ & $4.3 \pm 0.2$ & & \\ \hline & & & & Focus offset $1 \mathrm{~mA}$ at $650^{\circ} \mathrm{C}$ & $820.8 \pm 2.5$ & $952.1 \pm 14.9$ & $19.5 \pm 6.4$ & & \\ \hline & & & & $3 \mathrm{~mA}$ & $827.8 \pm 2.6$ & $955.5 \pm 3.2$ & $18.8 \pm 1.1$ & & \\ \hline & & & & $6 \mathrm{~mA}$ & $840.0 \pm 10.8$ & $967.8 \pm 12.4$ & $17.2 \pm 0.8$ & & \\ \hline & & & & $9 \mathrm{~mA}$ & $943.3 \pm 27.2$ & $1051.8 \pm 12.4$ & $29.8 \pm 2.5$ & & \\ \hline & & & & $12 \mathrm{~mA}$ & $917.0 \pm 4.9$ & $1037.5 \pm 2.1$ & $30.3 \pm 1.1$ & & \\ \hline & & & & $15 \mathrm{~mA}$ & $922.3 \pm 12.4$ & $1041.3 \pm 7.4$ & $22.8 \pm 3.2$ & & \\ \hline & & & & $18 \mathrm{~mA}$ & $862.8 \pm 7.4$ & $988.8 \pm 2.5$ & $17.8 \pm 0.4$ & & \\ \hline & & & & $21 \mathrm{~mA}$ & $866.3 \pm 2.5$ & $994.3 \pm 4.9$ & $17.8 \pm 1.1$ & & \\ \hline & & & & $24 \mathrm{~mA}$ & $868.0 \pm 14.8$ & $967.8 \pm 51.9$ & $14.5 \pm 11.3$ & & \\ \hline & & & & $30 \mathrm{~mA}$ & $516.3 \pm 86.6$ & $537.3 \pm 57.1$ & $4.5 \pm 0.9$ & & \\ \hline & & & & $40 \mathrm{~mA}$ & $264.3 \pm 17.3$ & $320.7 \pm 7.1$ & $5.8 \pm 3.2$ & & \\ \hline \multirow[t]{4}{*}{ЕBM} & IN718 & & Strain rate $1.5 \times 10^{-4} \mathrm{~s}^{-1}$ & $0^{\circ}$ & $793 \pm 4$ & $809 \pm 14$ & $1 \pm 0.5$ & & [308] \\ \hline & & & & $45^{\circ}$ & $757 \pm 8$ & $776 \pm 12$ & $16.2 \pm 5.5$ & & \\ \hline & & & & $55^{\circ}$ & $843 \pm 13$ & $951 \pm 10$ & $11.4 \pm 2.7$ & & \\ \hline & & & & $90^{\circ}$ & $815 \pm 27$ & $879 \pm 27$ & $8.5 \pm 3.1$ & & \\ \hline \multirow[t]{3}{*}{EBM} & IN718 & & Strain rate $10^{-3} \mathrm{~s}^{-1}$ & As-build & $920 \pm 16$ & $1075 \pm 46$ & $10 \pm 3$ & $138 \pm 5$ & [321] \\ \hline & & & & STA & $1096 \pm 6$ & $1172 \pm 30$ & $6 \pm 1$ & $137 \pm 7$ & \\ \hline & & & & HIP + STA & $1100 \pm 13$ & $1190 \pm 33$ & $14 \pm 1$ & $142 \pm 4$ & \\ \hline \multirow[t]{2}{*}{ЕвМ} & IN718 & Procedure from [441] & & $0^{\circ}$ & & & & 220 & [442] \\ \hline & & & & $90^{\circ}$ & & & & 125 & \\ \hline \multirow[t]{2}{*}{EBM} & IN718 & Quasi-static tensile tests & Speed control $0.0025 \mathrm{~mm} / \mathrm{s}$ & $\mathrm{AB}$ & $406 \pm 17$ & $427 \pm 28$ & $0.7 \pm 0.2$ & & [422] \\ \hline & & & & Polished & $982 \pm 52$ & $1174 \pm 29$ & $27.8 \pm 1.4$ & & \\ \hline EBM & IN718 & \begin{tabular}{l} ASTM-E8/E8M, ASTM-E21, \\ Strain rate $0.005(\mathrm{~mm} / \mathrm{mm}) /$ \\ \end{tabular} & $\mathrm{RT}$ and $650^{\circ} \mathrm{C}$ & Various build heights & See Fig. 11 in the orig & tal reference paper & & & [96] \\ \hline \end{tabular} \end{center} Appendix C. Hardness properties of PBF Nickel-based superalloys Table 10 Summary of hardness properties of PBF Nickel-based superalloys (' ' indicates that the data was obtained from the bar chart). \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample condition & Hardness & Ref \\ \hline \multirow[t]{2}{*}{LPBF} & \multirow[t]{2}{*}{CM247LC} & \multirow[t]{2}{*}{Vickers hardness} & \multirow[t]{2}{*}{$5 \mathrm{kgf}$} & $\mathrm{AB}$ & $400 \pm 9 \mathrm{HV}$ & \multirow[t]{2}{*}{[184]} \\ \hline & & & & SA & $512 \pm 9 \mathrm{HV}$ & \\ \hline \multirow[t]{5}{*}{LPBF} & \multirow[t]{5}{*}{CM247LC} & \multirow[t]{5}{*}{Vickers hardness} & \multirow[t]{5}{*}{$5 \mathrm{kgf}$} & $\mathrm{AB}$ & $409 \pm 7 \mathrm{HV}$ & \multirow[t]{5}{*}{[311]} \\ \hline & & & & \begin{tabular}{l} Solution @ $1210^{\circ} \mathrm{C}$ \\ + ageing \\ \end{tabular} & $442 \pm 16 \mathrm{HV}$ & \\ \hline & & & & \begin{tabular}{l} Solution @ $1230^{\circ} \mathrm{C}$ \\ + ageing \\ \end{tabular} & $437 \pm 19 \mathrm{HV}$ & \\ \hline & & & & \begin{tabular}{l} Solution @ $1240{ }^{\circ} \mathrm{C}$ \\ + ageing \\ \end{tabular} & $448 \pm 23 \mathrm{HV}$ & \\ \hline & & & & \begin{tabular}{l} Solution @ $1260{ }^{\circ} \mathrm{C}$ \\ + ageing \\ \end{tabular} & $462 \pm 13 \mathrm{HV}$ & \\ \hline \multirow[t]{6}{*}{LPBF} & \multirow[t]{6}{*}{CM247LC} & \multirow[t]{6}{*}{Vickers Hardness} & \multirow[t]{6}{*}{$1 \mathrm{~kg}$} & $\mathrm{HT}$ at $450^{\circ} \mathrm{C} / 2 \mathrm{~h}$ & $430 \mathrm{HV}$ & \multirow[t]{6}{*}{[195]} \\ \hline & & & & $\mathrm{HT}$ at $600^{\circ} \mathrm{C} / 2 \mathrm{~h}$ & $425 \mathrm{HV}$ & \\ \hline & & & & $\mathrm{HT}$ at $700^{\circ} \mathrm{C} / 2 \mathrm{~h}$ & $460 \mathrm{HV}$ & \\ \hline & & & & $\mathrm{HT}$ at $750^{\circ} \mathrm{C} / 2 \mathrm{~h}$ & $490 \mathrm{HV}$ & \\ \hline & & & & $\mathrm{HT}$ at $850^{\circ} \mathrm{C} / 2 \mathrm{~h}$ & $545 \mathrm{HV}$ & \\ \hline & & & & $\mathrm{HT}$ at $975^{\circ} \mathrm{C} / 2 \mathrm{~h}$ & $500 \mathrm{HV}$ & \\ \hline \multirow[t]{4}{*}{LPBF} & FGH100L & Vickers Hardness & & $\mathrm{AB}$ & $\sim 410 \mathrm{HV}$ & [324] \\ \hline & & & & HT & $\sim 490 \mathrm{HV}$ & \\ \hline & & & & HIP & $\sim 475 \mathrm{HV}$ & \\ \hline & & & & $\mathrm{HIP}+\mathrm{HT}$ & $\sim 590 \mathrm{HV}$ & \\ \hline LPBF & FGH4096 M & Hardness & & $\mathrm{AB}$ & $\sim 280 \mathrm{HB}$ & [341] \\ \hline & & & & DA & $\sim 445 \mathrm{HB}$ & \\ \hline & & & & STA $\left(1050^{\circ} \mathrm{C}\right)$ & $\sim 390 \mathrm{HB}$ & \\ \hline & & & & STA $\left(1130^{\circ} \mathrm{C}\right)$ & $\sim 425 \mathrm{HB}$ & \\ \hline & & & & Double Aging & $\sim 385 \mathrm{HB}$ & \\ \hline LPBF & Hastelloy X & Vickers hardness & $5 \mathrm{kgf}$ & $0^{\circ} / \mathrm{AB}$ & $\sim 246.07 \mathrm{HV}$ & [179] \\ \hline & & & & $90^{\circ} / \mathrm{AB}$ & $\sim 243.37 \mathrm{HV}$ & \\ \hline & & & & $0^{\circ} / \mathrm{HT}$ & $\sim 212.36 \mathrm{HV}$ & \\ \hline & & & & $90^{\circ} / \mathrm{HT}$ & $\sim 211.69 \mathrm{HV}$ & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}$ & $\sim 205.62 \mathrm{HV}$ & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}$ & $\sim 206.29 \mathrm{HV}$ & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}+\mathrm{HT}$ & $\sim 217.08 \mathrm{HV}$ & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}+\mathrm{HT}$ & $\sim 215.73 \mathrm{HV}$ & \\ \hline LPBF & Hastelloy X & Vickers hardness & $500 \mathrm{gf}$ & Original material & $277.1 \pm 3.9 \mathrm{HV}$ & [284] \\ \hline & & & & Modified material & $280.9 \pm 4.0 \mathrm{HV}$ & \\ \hline LPBF & Hastelloy X & Vickers Hardness (HV 0.5) & $200 \mathrm{gf} / 15 \mathrm{~s}$ & $90^{\circ} / \mathrm{AB}$ & $301 \pm 8 \mathrm{HV}$ & [180] \\ \hline & & Room temperature & RT & $90^{\circ} / \mathrm{HT}$ & $\sim 195 \mathrm{HV}$ & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}$ & $\sim 231 \mathrm{HV}$ & \\ \hline & & & & $0^{\circ} / \mathrm{AB}$ & $308 \pm 12 \mathrm{HV}$ & \\ \hline & & & & $0^{\circ} / \mathrm{HT}$ & $\sim 208 \mathrm{HV}$ & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}$ & $\sim 238 \mathrm{HV}$ & \\ \hline & & & $200 \mathrm{gf} / 15 \mathrm{~s}$ & $90^{\circ} / \mathrm{AB}$ & $280 \pm 6 \mathrm{HV}$ & \\ \hline & & & $750^{\circ} \mathrm{C}$ & $90^{\circ} / \mathrm{HT}$ & $\sim 229 \mathrm{HV}$ & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}$ & $\sim 200 \mathrm{HV}$ & \\ \hline & & & & $0^{\circ} / \mathrm{AB}$ & $317 \pm 20 \mathrm{HV}$ & \\ \hline & & & & $0^{\circ} / \mathrm{HT}$ & $\sim 212 \mathrm{HV}$ & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}$ & $\sim 200 \mathrm{HV}$ & \\ \hline LPBF (SLE) & IN100 & Vickers hardness & $2000 \mathrm{gf} / 10-15 \mathrm{~s}$ & Substrate & $\sim 389.95 \mathrm{HV}$ & [351] \\ \hline & & & & Interface & $\sim 404.90 \mathrm{HV}$ & \\ \hline & & & & Deposited materials & $\sim 426.96 \mathrm{HV}$ & \\ \hline LPBF & IN625/TiB2 & Vickers hardness & $300 \mathrm{gf} / 15 \mathrm{~s}$ & \begin{tabular}{l} Linear energy \\ density (LED) 1200 \\ $\mathrm{~J} / \mathrm{m} / \mathrm{IN} 625$ \\ \end{tabular} & $\sim 299.02 \mathrm{HV}$ & [346] \\ \hline & & & & \begin{tabular}{l} LED $1200 \mathrm{~J} / \mathrm{m} /$ \\ IN625+TiB 2 \\ \end{tabular} & $\sim 626.23 \mathrm{HV}$ & \\ \hline & & & & \begin{tabular}{l} LED $800 \mathrm{~J} / \mathrm{m} /$ \\ IN625 \\ \end{tabular} & $\sim 338.24 \mathrm{HV}$ & \\ \hline & & & & \begin{tabular}{l} $\mathrm{LED} 800 \mathrm{~J} / \mathrm{m} /$ \\ $\mathrm{IN}^{25}+\mathrm{TiB}_{2}$ \\ \end{tabular} & $\sim 549.02 \mathrm{HV}$ & \\ \hline & & & & \begin{tabular}{l} LED $600 \mathrm{~J} / \mathrm{m} /$ \\ IN625 \\ \end{tabular} & $\sim 338.24 \mathrm{HV}$ & \\ \hline & & & & LED $600 \mathrm{~J} / \mathrm{m} /$ & $\sim 627.45 \mathrm{HV}$ & \\ \hline & & & & IN625+TiB & .021 .70118 & \\ \hline & & & & \begin{tabular}{l} LED $400 \mathrm{~J} / \mathrm{m} /$ \\ IN625 \\ \end{tabular} & $\sim 370.10 \mathrm{HV}$ & \\ \hline & & & & \begin{tabular}{l} $\mathrm{LED} 400 \mathrm{~J} / \mathrm{m} /$ \\ IN625+TiB 2 \\ \end{tabular} & $\sim 688.73 \mathrm{HV}$ & \\ \hline LPBF & IN625 & Vickers hardness & $500 \mathrm{gf} / 30 \mathrm{~s}$ & $\mathrm{AB}$ & $343 \mathrm{HV}$ & [348] \\ \hline & & & & Solution @ $700^{\circ} \mathrm{C}$ & $\sim 334.17 \mathrm{HV}$ & \\ \hline \end{tabular} \end{center} Table 10 (continued) \begin{center} \begin{tabular}{lll} \hline Technique & Material & Standard \\ \hline & & \\ & & \\ LPBF & IN625 & Brinell hardness \\ LPBF & IN625 & Nano-indentation \\ \end{tabular} \end{center} \begin{center} \begin{tabular}{|c|c|c|} \hline Sample condition & Hardness & Ref \\ \hline Solution @ $800^{\circ} \mathrm{C}$ & $\sim 356.20 \mathrm{HV}$ & \\ \hline Solution @ $900^{\circ} \mathrm{C}$ & $\sim 356.23 \mathrm{HV}$ & \\ \hline Solution @ $1000{ }^{\circ} \mathrm{C}$ & $\sim 276.02 \mathrm{HV}$ & \\ \hline Solution @ $1100{ }^{\circ} \mathrm{C}$ & $\sim 265.06 \mathrm{HV}$ & \\ \hline Solution @ $1200^{\circ} \mathrm{C}$ & $\sim 260.20 \mathrm{HV}$ & \\ \hline Hardness vs & Fig. $52 b$ & [353] \\ \hline Porosity & & \\ \hline 90 W/AB/0 dpa & $\sim 5.75 \mathrm{GPa}$ & [376] \\ \hline $90 \mathrm{~W} / \mathrm{AB} / 0.1 \mathrm{dpa}$ & $\sim 6.10 \mathrm{GPa}$ & \\ \hline $90 \mathrm{~W} / \mathrm{AB} / 0.5 \mathrm{dpa}$ & $\sim 6.05 \mathrm{GPa}$ & \\ \hline 90 W/AB/1 dpa & $\sim 5.6 \mathrm{GPa}$ & \\ \hline $90 \mathrm{~W} / \mathrm{AB} / 3 \mathrm{dpa}$ & $\sim 5.8 \mathrm{GPa}$ & \\ \hline $\mathrm{AB}$ & $\sim 283.25 \mathrm{HBW}$ & [303] \\ \hline $\mathrm{DA} @ 600{ }^{\circ} \mathrm{C}$ for $2 \mathrm{~h}$ & $\sim 275.04 \mathrm{HBW}$ & \\ \hline $\mathrm{DA} @ 700^{\circ} \mathrm{C}$ for $2 \mathrm{~h}$ & $\sim 311.70 \mathrm{HBW}$ & \\ \hline DA @ $800^{\circ} \mathrm{C}$ for $2 \mathrm{~h}$ & $\sim 395.91 \mathrm{HBW}$ & \\ \hline DA @ $900{ }^{\circ} \mathrm{C}$ for $2 \mathrm{~h}$ & $\sim 282.37 \mathrm{HBW}$ & \\ \hline DA @ $600^{\circ} \mathrm{C}$ for $8 \mathrm{~h}$ & $\sim 288.75 \mathrm{HBW}$ & \\ \hline DA @ $700^{\circ} \mathrm{C}$ for $8 \mathrm{~h}$ & $\sim 319.78 \mathrm{HBW}$ & \\ \hline $\mathrm{DA} @ 800^{\circ} \mathrm{C}$ for $8 \mathrm{~h}$ & $\sim 322.03 \mathrm{HBW}$ & \\ \hline DA @ $900{ }^{\circ} \mathrm{C}$ for $8 \mathrm{~h}$ & $\sim 298.91 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} DA @ $600^{\circ} \mathrm{C}$ for \\ $24 \mathrm{~h}$ \\ \end{tabular} & $\sim 306.80 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} DA @ $700{ }^{\circ} \mathrm{C}$ for \\ $24 \mathrm{~h}$ \\ \end{tabular} & $\sim 349.39 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} $\mathrm{DA} @ 800^{\circ} \mathrm{C}$ for \\ $24 \mathrm{~h}$ \\ \end{tabular} & $\sim 348.55$ HBW & \\ \hline \begin{tabular}{l} DA @ $900{ }^{\circ} \mathrm{C}$ for \\ $24 \mathrm{~h}$ \\ \end{tabular} & $\sim 316.96 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1000{ }^{\circ} \mathrm{C}$ \\ for $1 \mathrm{~h}$ \\ \end{tabular} & $\sim 212.12 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1150^{\circ} \mathrm{C}$ \\ for $1 \mathrm{~h}$ \\ \end{tabular} & $\sim 186.22 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1000^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}$ \\ \end{tabular} & $\sim 209.88 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}$ \\ \end{tabular} & $\sim 188.96 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}+$ Ageing @ \\ $600^{\circ} \mathrm{C}$ for $2 \mathrm{~h}$ \\ \end{tabular} & $\sim 198.90 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}+$ Ageing @ \\ $700^{\circ} \mathrm{C}$ for $2 \mathrm{~h}$ \\ \end{tabular} & $\sim 209.92 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}+$ Ageing $@$ \\ $800^{\circ} \mathrm{C}$ for $2 \mathrm{~h}$ \\ \end{tabular} & $\sim 194.87 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}+$ Ageing @ \\ $900^{\circ} \mathrm{C}$ for $2 \mathrm{~h}$ \\ \end{tabular} & $\sim 188.75 \mathrm{HWB}$ & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}+$ Ageing @ \\ $600^{\circ} \mathrm{C}$ for $8 \mathrm{~h}$ \\ \end{tabular} & $\sim 206.95 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}+$ Ageing @ \\ $700^{\circ} \mathrm{C}$ for $8 \mathrm{~h}$ \\ \end{tabular} & $\sim 259.97 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}+$ Ageing @ \\ $800^{\circ} \mathrm{C}$ for $8 \mathrm{~h}$ \\ \end{tabular} & $\sim 210.97 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}+$ Ageing @ \\ $900^{\circ} \mathrm{C}$ for $8 \mathrm{~h}$ \\ \end{tabular} & $\sim 188.75$ HBW & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}+$ Ageing @ \\ $600^{\circ} \mathrm{C}$ for $24 \mathrm{~h}$ \\ \end{tabular} & $\sim 227.94 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}+$ Ageing @ \\ $700^{\circ} \mathrm{C}$ for $24 \mathrm{~h}$ \\ \end{tabular} & $\sim 279.91 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}+$ Ageing @ \\ $800^{\circ} \mathrm{C}$ for $24 \mathrm{~h}$ \\ \end{tabular} & $\sim 234.94 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} Solution @ $1150{ }^{\circ} \mathrm{C}$ \\ for $2 \mathrm{~h}+$ Ageing @ \\ $900^{\circ} \mathrm{C}$ for $24 \mathrm{~h}$ \\ \end{tabular} & $\sim 192.95 \mathrm{HBW}$ & \\ \hline \begin{tabular}{l} $90^{\circ} /$ Powder $1 / 90$ \\ $\mathrm{~J} / \mathrm{mm}^{3}$ \\ \end{tabular} & $\sim 289.79 \mathrm{HV}$ & [121] \\ \hline \end{tabular} \end{center} $\sim 275.58 \mathrm{HV}$ Table 10 (continued) \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample condition & \multicolumn{2}{|l|}{Hardness} & Ref \\ \hline & & & & \begin{tabular}{l} $0^{\circ} /$ Powder 1/90 J/ \\ $\mathrm{mm}^{3}$ \\ \end{tabular} & & & \\ \hline & & & & \begin{tabular}{l} $90^{\circ} /$ Powder $3 / 90$ \\ $\mathrm{~J} / \mathrm{mm}^{3}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$\sim 303.18 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} $0^{\circ} /$ Powder $3 / 90 \mathrm{~J} /$ \\ $\mathrm{mm}^{3}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$\sim 287.03 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} $90^{\circ} /$ Powder 1/100 \\ $\mathrm{J} / \mathrm{mm}^{3}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$\sim 281.83 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} $0^{\circ} /$ Powder 1/100 \\ $\mathrm{J} / \mathrm{mm}^{3}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$\sim 275.58 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} $90^{\circ} /$ Powder $3 / 100$ \\ $\mathrm{~J} / \mathrm{mm}^{3}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$\sim 295.31 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} $0^{\circ} /$ Powder 3/100 \\ $\mathrm{J} / \mathrm{mm}^{3}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$\sim 285.73 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} $90^{\circ} /$ Powder 1/110 \\ $\mathrm{J} / \mathrm{mm}^{3}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$\sim 277.86 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} $0^{\circ} /$ Powder 1/110 \\ $\mathrm{J} / \mathrm{mm}^{3}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$\sim 271.20 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} $90^{\circ} /$ Powder 3/110 \\ $\mathrm{J} / \mathrm{mm}^{3}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$\sim 298.47 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} $0^{\circ} /$ Powder 3/110 \\ $\mathrm{J} / \mathrm{mm}^{3}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$\sim 286.22 \mathrm{HV}$} & \\ \hline LPBF & IN625 and cast iron & Micro Hardness & $500 \mathrm{gm} / 20 \mathrm{~s}$ & Cast iron susbtrate & \multicolumn{2}{|c|}{\multirow{2}{*}}{}\{\begin{tabular}{l} $\sim 200 \mathrm{HV}$ \\ $\sim 375 \mathrm{HV}$ \\ \end{tabular}\} & [371] \\ \hline & substrate & & & HAZ & & & \\ \hline & & & & \begin{tabular}{l} Sample-substrate \\ interface \\ \end{tabular} & \multicolumn{2}{|l|}{$\sim 350 \mathrm{HV}$} & \\ \hline & & & & IN625 sample & \multicolumn{2}{|l|}{$\sim 300 \mathrm{HV}$} & \\ \hline LPBF & IN625 & Microhardness & $100 \mathrm{~g} / 10 \mathrm{~s}$ & \multicolumn{3}{|c|}{See the original reference paper} & [416] \\ \hline \multirow[t]{12}{*}{LPBF} & IN718 & Vickers hardness & $1 \mathrm{kgf}$ & $250 \mathrm{~W} / \mathrm{AB}$ & \multicolumn{2}{|l|}{$320 \mathrm{HV}$} & $[67]$ \\ \hline & & & & $950 \mathrm{~W} / \mathrm{AB}$ & \multicolumn{2}{|l|}{$287 \mathrm{HV}$} & \\ \hline & & & & $250 \mathrm{~W} / \mathrm{HT}$ & \multicolumn{2}{|l|}{$360 \mathrm{HV}$} & \\ \hline & & & & $950 \mathrm{~W} / \mathrm{HT}$ & \multicolumn{2}{|l|}{$338 \mathrm{HV}$} & \\ \hline & & & & $250 \mathrm{~W} / \mathrm{HIP}$ & \multicolumn{2}{|l|}{$310 \mathrm{HV}$} & \\ \hline & & & & $950 \mathrm{~W} / \mathrm{HIP}$ & \multicolumn{2}{|l|}{$262 \mathrm{HV}$} & \\ \hline & & & & $250 \mathrm{~W} / \mathrm{HIP}+\mathrm{HT}$ & \multicolumn{2}{|l|}{$468 \mathrm{HV}$} & \\ \hline & & & & $950 \mathrm{~W} / \mathrm{HIP}+\mathrm{HT}$ & \multicolumn{2}{|l|}{$451 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} Zone $1-250 \mathrm{~W}$ \\ Matrix and zone $2-$ \\ two lines of $950 \mathrm{~W} /$ \\ $\mathrm{AB}$ \\ \end{tabular} & $330 \mathrm{HV}$ & $285 \mathrm{HV}$ & \\ \hline & & & & \begin{tabular}{l} Zone $1-250 \mathrm{~W}$ \\ Matrix and zone 2 - \\ two lines of $950 \mathrm{~W} /$ \\ HT \\ \end{tabular} & $370 \mathrm{HV}$ & $335 \mathrm{HV}$ & \\ \hline & & & & \begin{tabular}{l} Zone $1-250 \mathrm{~W}$ \\ Matrix and zone 2 - \\ two lines of $950 \mathrm{~W} /$ \\ HIP \\ \end{tabular} & $310 \mathrm{HV}$ & $260 \mathrm{HV}$ & \\ \hline & & & & \begin{tabular}{l} Zone $1-250 \mathrm{~W}$ \\ Matrix and zone 2 - \\ two lines of $950 \mathrm{~W} /$ \\ HIP + HT \\ \end{tabular} & $478 \mathrm{HV}$ & $462 \mathrm{HV}$ & \\ \hline LPBF & IN718 & Vickers hardness & $1 \mathrm{kgf}$ & \begin{tabular}{l} Zone $1-250 \mathrm{~W}$ \\ Matrix and Zone 2 - \\ two lines of $950 \mathrm{~W}$ \\ \end{tabular} & $330 \mathrm{HV}$ & $285 \mathrm{HV}$ & [281] \\ \hline & & & & \begin{tabular}{l} Zone $1-950 \mathrm{~W}$ and \\ Zone $2-$ two lines \\ of $250 \mathrm{~W}$ \\ \end{tabular} & $322 \mathrm{HV}$ & $300 \mathrm{HV}$ & \\ \hline & & & & \begin{tabular}{l} Zone $1-250 \mathrm{~W}$ and \\ Zone $2-$ four lines \\ of $950 \mathrm{~W}$ \\ \end{tabular} & $318 \mathrm{HV}$ & $285 \mathrm{HV}$ & \\ \hline & & & & \begin{tabular}{l} Zone $1-950 \mathrm{~W}$ and \\ Zone $2-$ four lines \\ of $250 \mathrm{~W}$ \\ \end{tabular} & $311 \mathrm{HV}$ & $289 \mathrm{HV}$ & \\ \hline LPBF & IN718+WC & Vickers hardness & $200 \mathrm{gf}$ & \begin{tabular}{l} Scan speed 400 \\ $\mathrm{~mm} / \mathrm{s}$ \\ \end{tabular} & $\sim 385.30$ & & $[65]$ \\ \hline & & & & \begin{tabular}{l} Scan speed 500 \\ $\mathrm{~mm} / \mathrm{s}$ \\ \end{tabular} & $\sim 402.47$ & & \\ \hline & & & & \begin{tabular}{l} Scan speed 600 \\ $\mathrm{~mm} / \mathrm{s}$ \\ \end{tabular} & $\sim 445.39$ & & \\ \hline & & & & \begin{tabular}{l} Scan speed 700 \\ $\mathrm{~mm} / \mathrm{s}$ \\ \end{tabular} & $\sim 480.26$ & & \\ \hline LPBF & \begin{tabular}{l} $\mathrm{WC}_{1-\mathrm{x}}$ reinforced \\ IN718 \\ \end{tabular} & Vickers hardness & $100 \mathrm{gf} / 15 \mathrm{~s}$ & \begin{tabular}{l} Linear energy \\ density \\ \end{tabular} & Fig. $53 a$ & & [343] \\ \hline LPBF & IN718 & Vickers hardness & $100 \mathrm{gf} / 10 \mathrm{~s}$ & \begin{tabular}{l} Linear energy \\ density \\ \end{tabular} & Fig. $53 b$ & & [344] \\ \hline \end{tabular} \end{center} Table 10 (continued) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-60} \end{center} Table 10 (continued) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-61} \end{center} Table 10 (continued) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-62} \end{center} Table 10 (continued) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-63} \end{center} Table 10 (continued) \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample condition & \multicolumn{2}{|l|}{Hardness} & Ref \\ \hline & & & & \begin{tabular}{l} Various process \\ parameters \\ \end{tabular} & \multicolumn{2}{|c|}{}\begin{tabular}{l} See Table in the original \\ reference paper \\ \end{tabular} & \\ \hline \multirow[t]{2}{*}{DMD} & 12N-01 alloy & Vickers Hardness & $100 \mathrm{gf}$ & Single track & \multicolumn{2}{|c|}{$350-400 \mathrm{HV}$} & [378] \\ \hline & & & & Double track & \multicolumn{2}{|c|}{$500-580 \mathrm{HV}$} & \\ \hline EBM & IN625 & Vickers hardness & $1 \mathrm{kgf} / 10 \mathrm{~s}$ & $\mathrm{AB}$ & $335 \mathrm{HV}$ & & [305] \\ \hline EBM & IN625 & Vickers hardness & & Mesh structure & $2.9 \mathrm{GPa}$ & 295.7 HV & [293] \\ \hline \multirow[t]{13}{*}{EBM} & IN625 & Vickers hardness and Rockwell C & $100 \mathrm{gf} / 10 \mathrm{~s} /$ Vickers & Powder & $2.6 \mathrm{GPa}$ & 265.1 HV & [292] \\ \hline & & & & $0^{\circ} / \mathrm{AB}$ & $2.8 \mathrm{GPa}$ & $285.5 \mathrm{HV}$ & \\ \hline & & & & $90^{\circ} / \mathrm{AB}$ & $2.5 \mathrm{GPa}$ & 254.9 HV & \\ \hline & & & & $0^{\circ} / \mathrm{HIP}$ & $2.2 \mathrm{GPa}$ & $224.3 \mathrm{HV}$ & \\ \hline & & & & $90^{\circ} / \mathrm{HIP}$ & $2.1 \mathrm{GPa}$ & 214.1 HV & \\ \hline & & & & $0^{\circ} / \mathrm{AB} / 538^{\circ} \mathrm{C}$ & $2.6 \mathrm{GPa}$ & 265.1 HV & \\ \hline & & & & $90^{\circ} / \mathrm{AB} / 538^{\circ} \mathrm{C}$ & $2.8 \mathrm{GPa}$ & $285.5 \mathrm{HV}$ & \\ \hline & & & & $0^{\circ} / \mathrm{HIP} / 538^{\circ} \mathrm{C}$ & $2.3 \mathrm{GPa}$ & 234.5 HV & \\ \hline & & & & $90^{\circ} / \mathrm{HIP} / 538^{\circ} \mathrm{C}$ & $2.2 \mathrm{GPa}$ & 224.3 HV & \\ \hline & & & Rockwell C $150 \mathrm{kgf}$ & $\mathrm{AB}$ & $14 \mathrm{HRC}$ & $197 \mathrm{HV}$ & \\ \hline & & & & HIP & 8 HRC & $175 \mathrm{HV}$ & \\ \hline & & & & $\mathrm{AB} / 538^{\circ} \mathrm{C}$ & $14 \mathrm{HRC}$ & $197 \mathrm{HV}$ & \\ \hline & & & & $\mathrm{HIP} / 538^{\circ} \mathrm{C}$ & 6 HRC & $172 \mathrm{HV}$ & \\ \hline EBM & In690 & Vickers hardness & $100 \mathrm{gf}$ & & \multicolumn{2}{|c|}{$233 \pm 12 \mathrm{HV}$} & [310] \\ \hline \multirow[t]{2}{*}{EBM} & In690 & Hardness & $100 \mathrm{gf}$ with $2 \mathrm{~mm}$ & Cladding & \multicolumn{2}{|c|}{\multirow{2}{*}}{}\{\begin{tabular}{l} $2.33 \pm 0.12 \mathrm{GPa}$ \\ $1.78 \pm 0.04 \mathrm{GPa}$ \\ \end{tabular}\} & $[330]$ \\ \hline & & & intervals & On build plate & & & \\ \hline \multirow[t]{8}{*}{EBM} & IN718 & Brinell hardness for macro hardness, Vickers & Brinell $187.5 \mathrm{kPa}(1839$ & $0^{\circ} /$ Macro- & $337 \pm$ & $354 \mathrm{HV}$ & [295] \\ \hline & & hardness for microhardness & & hardness/AB & & & \\ \hline & & & & $0^{\circ} /$ Macro- & $430 \pm$ & $452 \mathrm{HV}$ & \\ \hline & & & & hardness/HT & $10 \mathrm{HB}$ & & \\ \hline & & & & \begin{tabular}{l} $0^{\circ} /$ Micro-hardness/ \\ $\mathrm{AB}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$506 \pm 26 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} $90^{\circ} /$ Micro- \\ hardness/ $\mathrm{AB}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$502 \pm 29 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} $0^{\circ} /$ Micro-hardness/ \\ $\mathrm{HT}$ \\ \end{tabular} & \multicolumn{2}{|c|}{$640 \pm 15 \mathrm{HV}$} & \\ \hline & & & & \begin{tabular}{l} $90^{\circ} /$ Micro- \\ hardness/HT \\ \end{tabular} & \multicolumn{2}{|c|}{$646 \pm 22 \mathrm{HV}$} & \\ \hline \multirow[t]{2}{*}{EBM} & IN718 & Vickers hardness & & $\mathrm{AB}$ & \multicolumn{2}{|l|}{$410 \mathrm{HV}$} & [304] \\ \hline & & & & HT & \multicolumn{2}{|l|}{$470 \mathrm{HV}$} & \\ \hline EBM & IN718 & Vickers hardness & & Fast cooled & $324 \pm 18$ & & [183] \\ \hline & & & & Slow cooled & $392 \pm 15$ & & \\ \hline & & & & In-situ HT & $478 \pm 7$ & & \\ \hline EBM & IN718 & Vickers hardness & $300 \mathrm{gf} / 15 \mathrm{~s}$ & $\mathrm{AB}$ & $427.5 \mathrm{HV}$ & & [294] \\ \hline & & & & DA & $488.0 \mathrm{HV}$ & & \\ \hline & & & & STA $\left(930^{\circ} \mathrm{C}\right)$ & $479.6 \mathrm{HV}$ & & \\ \hline & & & & STA $\left(980^{\circ} \mathrm{C}\right)$ & $478.7 \mathrm{HV}$ & & \\ \hline & & & & STA $\left(1080^{\circ} \mathrm{C}\right)$ & $472.7 \mathrm{HV}$ & & \\ \hline EBM & IN718 & Vickers hardness & N.A. & $\mathrm{AB} / \mathrm{HT}$ & Fig. 56 & & [308] \\ \hline EBM & IN718 & Vickers hardness & $500 \mathrm{gf} / 15 \mathrm{~s}$ & $\mathrm{AB}$ & $438.41 \pm$ & $8.35 \mathrm{HV}$ & [160] \\ \hline & & & & HIP & $199.37 \pm$ & $10.44 \mathrm{HV}$ & \\ \hline & & & & $\mathrm{HIP}+\mathrm{HT}$ & $482.26 \pm$ & $7.83 \mathrm{HV}$ & \\ \hline EBM & IN718 & Vickers hardness & $1 \mathrm{kgf} / 15 \mathrm{~s}$ & $\mathrm{AB}$ & $421.80 \mathrm{H}$ & & [321] \\ \hline & & & & STA & $468.72 \mathrm{H}$ & & \\ \hline & & & & $\mathrm{HIP}+\mathrm{STA}$ & $474.38 \mathrm{H}$ & & \\ \hline EBM & IN718 & Vickers hardness & $500 \mathrm{gf} / 15 \mathrm{~s}$ & Set A & $387.92 \mathrm{H}$ & & [367] \\ \hline & & & & Set B & $386.95 \mathrm{H}$ & & \\ \hline & & & & Set C & $391.08 \mathrm{H}$ & & \\ \hline & & & & Set D & $394.96 \mathrm{H}$ & & \\ \hline EBM & \begin{tabular}{l} nickel-based \\ superalloy \\ \end{tabular} & Vickers hardness & $1 \mathrm{kgf}$ & $90^{\circ}$ & Fig. $52 a$ & & [146] \\ \hline EBM & Rene 142 & Vickers hardness \& Rockwell C & Vickers 100 gf/Rockwell & Powder & $3.4 \mathrm{GPa}$ & $346.7 \mathrm{HV}$ & [354] \\ \hline & & & $1.5 \mathrm{~N}$ & $0^{\circ} / \mathrm{AB}$ & $4.1 \mathrm{GPa}$ & $418.1 \mathrm{HV}$ & \\ \hline & & & & $90^{\circ} / \mathrm{AB}$ & $4.2 \mathrm{GPa}$ & $428.3 \mathrm{HV}$ & \\ \hline & & & & $0^{\circ} / \mathrm{AB}$ & 39 HRC & $379 \mathrm{HV}$ & \\ \hline & & & & $90^{\circ} / \mathrm{AB}$ & 42 HRC & $406 \mathrm{HV}$ & \\ \hline \end{tabular} \end{center} Appendix D. Fatigue properties of PBF Nickel-based superalloys Table 11 Summary of fatigue properties of PBF Nickel-based superalloys (' $\sim$ ' indicates that the data was obtained from the bar chart). \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & \begin{tabular}{l} Load ratio \\ (R) \\ \end{tabular} & Frequency & Sample condition & Ref \\ \hline LPBF & EP708 & & \begin{tabular}{l} Fatigue endurance study at room \\ temperature, load from 340 to \\ $380 \mathrm{MPa}$, cycles from $2 \cdot 10^{6}$ to \\ sample failure \\ \end{tabular} & & & \begin{tabular}{l} Hot rolling + HT \\ HT \\ HIP + HT \\ \end{tabular} & [395] \\ \hline LPBF & Hastelloy X & & \begin{tabular}{l} Low cycle fatigue test at RT in \\ strain control with triangular \\ wave shape, Strain ratio $=1$ \\ Thermomechanical fatigue test \\ elevated temperature using a \\ trapezoid waveform, Strain ratio \\ $=0$ \\ \end{tabular} & & & \begin{tabular}{l} Specimens were built in the $0^{\circ}, 45^{\circ}$ and $90^{\circ}$ \\ direction \\ \end{tabular} & [384] \\ \hline LPBF & \begin{tabular}{l} Hastelloy ${ }^{\circledR}$ \\ X \\ \end{tabular} & & \begin{tabular}{l} Four-point bend fatigue test, stress \\ amplitude range $450 \mathrm{MPa}-900$ \\ $\mathrm{MPa}$ \\ S-N tension-tension fatigue test, \\ stress amplitude range 500 \\ $\mathrm{MPa}-800 \mathrm{MPa}$ \\ \end{tabular} & 0.1 & $117 \mathrm{~Hz}$ & \begin{tabular}{l} Specimens were built in the $0^{\circ}$ and $90^{\circ}$ \\ direction. $\mathrm{AB} / \mathrm{HIP}$ \\ \end{tabular} & [255] \\ \hline LPBF & IN625 & ASTM E647 & Fatigue test at RT & 0.1 & $20 \mathrm{~Hz}$ & \begin{tabular}{l} $0^{\circ} / \mathrm{SR}$ \\ $45^{\circ} / \mathrm{SR}$ \\ $90^{\circ} / \mathrm{SR}$ \\ $90^{\circ} \perp / \mathrm{SR}$ \\ $0^{\circ} / \mathrm{HIP}$ \\ $45^{\circ} / \mathrm{HIP}$ \\ $90^{\circ} / \mathrm{HIP}$ \\ $90^{\circ} \perp / \mathrm{HIP}$ \\ \end{tabular} & [312] \\ \hline LPBF & IN625 & & Fatigue Crack Growth at RT & Varied & $20 \mathrm{~Hz}$ & Different $\mathrm{R}$ ratios & [399] \\ \hline LPBF & IN625 & ASTM E466 & \begin{tabular}{l} Fatigue Limit $\left(10^{7}\right.$ cycles $)$ test at \\ RT and at $650^{\circ} \mathrm{C}$ \\ \end{tabular} & 0.1 & 20 & HIP & [438] \\ \hline LPBF & IN626 & & \begin{tabular}{l} Stress amplitude range 200 \\ $\mathrm{MPa}-300 \mathrm{MPa}$ \\ \end{tabular} & -1 & $20 \mathrm{~Hz}$ & $\mathrm{AB}$, polished specimens & [252] \\ \hline LPBF (DMLS) & IN718 & ASTM E606 & \begin{tabular}{l} Low cycle fatigue at RT with strain \\ amplitudes at $0.6 \%$, or $0.8 \%$, or \\ $1.0 \%$, or $1.2 \%$, or $1.4 \%$ and a \\ mean strain of $0.5 \%$, strain rate 4 \\ $\times 10^{-3} \mathrm{~s}^{-1}$ \\ \end{tabular} & & $0.15 \mathrm{~Hz}$ & \begin{tabular}{l} Specimens were built in the $0^{\circ}$ and $45^{\circ}$ \\ direction, SA/HIP \\ \end{tabular} & $[176]$ \\ \hline LPBF & IN718 & & The crack growth rate study & 0.1 & $25 \mathrm{~Hz}$ & \begin{tabular}{l} Modified with Re. Specimens were built in \\ the $0^{\circ}, 45^{\circ}$ and $90^{\circ}$ direction, SA \\ \end{tabular} & [385] \\ \hline LPBF & IN718 & ASTM E647 & \begin{tabular}{l} High cycle fatigue, fatigue test at \\ RT, $800^{\circ} \mathrm{F}$ and $1200^{\circ} \mathrm{F}$ \\ \end{tabular} & 0.01 & $40 \mathrm{~Hz}$ & \begin{tabular}{l} SR + HIP + HSA. AB surface condition, LSG \\ surface condition \\ \end{tabular} & [386] \\ \hline LPBF & IN718 & & \begin{tabular}{l} High cycle fatigue plane bending \\ testing \\ \end{tabular} & 0 & $20 \mathrm{~Hz}$ & \begin{tabular}{l} Specimens were built in the $0^{\circ}$ and $90^{\circ}$ \\ direction, SR + ageing \\ \end{tabular} & [387] \\ \hline LPBF & IN718 & & \begin{tabular}{l} Thermomechanical Fatigue \\ Testing, temperature cycling \\ between 350 and $650{ }^{\circ} \mathrm{C}$ with a \\ strain amplitude of $\pm 0.45 \%$. \\ Strain rate $=-1$ \\ \end{tabular} & & & \begin{tabular}{l} AB - 250,950 W or Functionally Graded \\ (FGM) \\ Heat treated - HA \\ \end{tabular} & [388] \\ \hline LPBF & IN718 & & \begin{tabular}{l} Low-cycle fatigue tests at ambient \\ temperature. Strain amplitudes of \\ $\Delta \varepsilon / 2= \pm 0.35 \%, \Delta \varepsilon / 2= \pm 0.5 \%$ \\ and $\Delta \varepsilon / 2= \pm 0.8 \%$. Strain rate = \\ $6 \times 10^{-3} \mathrm{~s}^{-1}$ \\ \end{tabular} & & & \begin{tabular}{l} AB/combinations of solution, ageing, HIP \\ and Arc-PVD \\ \end{tabular} & [172] \\ \hline LPBF & IN718 & ASTM E466 & Fatigue test at RT & 0.1 & $15 \mathrm{~Hz}$ & \begin{tabular}{l} HT $1-1200{ }^{\circ} \mathrm{C}$ for $1 \mathrm{~h}$ under argon \\ atmosphere \\ HT 2 - Max temperature of $980^{\circ} \mathrm{C}$ following \\ the heat/hold/cool cycle for $24 \mathrm{~h}$ \\ \end{tabular} & [256] \\ \hline LPBF (DMLS) & IN718 & & \begin{tabular}{l} High cycle fatigue test, vibration \\ bending testing, chord-wise \\ bending (or two-stripe) mode \\ \end{tabular} & & $1600 \mathrm{~Hz}$ & $\mathrm{AB}$ & [257] \\ \hline LPBF & IN718 & & \begin{tabular}{l} Fatigued samples were pre- \\ strained to $1 \%$ then tested \\ \end{tabular} & 0.05 & $1 \mathrm{~Hz}$ & $\mathrm{SR} / \mathrm{SR}+\mathrm{HSA}$ & [254] \\ \hline LPBF (DMLS) & IN718 & ASTM E466 & Fatigue test at RT & 0.1 & $0.5 \mathrm{~Hz}$ & \begin{tabular}{l} Dog bone structure, lattice structure. SR + \\ SA \\ \end{tabular} & [220] \\ \hline LPBF & IN718 & ASTM E647 & The crack growth rate study & 0.1 & \begin{tabular}{l} 80 to 60 \\ $\mathrm{~Hz}$ \\ \end{tabular} & \begin{tabular}{l} $\mathrm{AB}$, compact tension specimen with crack \\ growth plane parallel to the build direction \\ \end{tabular} & [389] \\ \hline LPBF & IN718 & ASTM E647 & The crack growth rate study & 0.1 & \begin{tabular}{l} 80 to 60 \\ $\mathrm{~Hz}$ \\ \end{tabular} & \begin{tabular}{l} $\mathrm{AB}$, compact tension specimen with crack \\ growth plane parallel to the build direction \\ \end{tabular} & [390] \\ \hline LPBF & IN718 & ASTM E649 & \begin{tabular}{l} Fatigue Crack Growth at RT and \\ $650^{\circ} \mathrm{C}$ \\ \end{tabular} & 0.1 & $10 \mathrm{~Hz}$ & $\mathrm{AB}$ & [424] \\ \hline LPBF & IN718 & ASTM E466 & High cycle fatigue at RT & -1 & & & [391] \\ \hline \end{tabular} \end{center} Table 11 (continued) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-66} \end{center} Table 11 (continued) \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & \begin{tabular}{l} Load ratio \\ (R) \\ \end{tabular} & Frequency & Sample condition & Ref \\ \hline & & & \begin{tabular}{l} Stress controlled fatigue test at \\ $650^{\circ} \mathrm{C}$ \\ \end{tabular} & & & & \\ \hline LPBF & IN939 & & \begin{tabular}{l} Low-cycle fatigue at RT and 750 \\ ${ }^{\circ} \mathrm{C}$, total strain amplitude $0.5 \%$, \\ strain rate $6 \times 10^{-3}$ \\ \end{tabular} & -1 & & $\mathrm{AB} / \mathrm{SA}$ & [253] \\ \hline \multirow[t]{2}{*}{LPBF} & K536 & ASTM E466 & Tested at $400^{\circ} \mathrm{C}$ & 0.1 & $100 \mathrm{~Hz}$ & \begin{tabular}{l} SR/built in $0^{\circ}$ \\ SR/built in $90^{\circ}$ \\ \end{tabular} & [319] \\ \hline & & & Tested at $600{ }^{\circ} \mathrm{C}$ & & & \begin{tabular}{l} SR/built in $0^{\circ}$ \\ SR/built in $90^{\circ}$ \\ \end{tabular} & \\ \hline LPBF & K536 & & \begin{tabular}{l} Stress controlled at $400^{\circ} \mathrm{C}$ and \\ $600{ }^{\circ} \mathrm{C}$ \\ \end{tabular} & 0.1 & $100 \mathrm{~Hz}$ & Specimens built at $0^{\circ}$ and $90^{\circ}$ & [319] \\ \hline EBM & IN718 & & \begin{tabular}{l} Constant amplitude Fatigue Limit \\ ( $10^{6}$ cycles) test at RT \\ \end{tabular} & -1 & 10 & $\mathrm{AB} /$ Polished & $[422]$ \\ \hline EBM & IN718 & & \begin{tabular}{l} Four-point bending fatigue test at \\ RT \\ \end{tabular} & 0.1 & $20 \mathrm{~Hz}$ & \begin{tabular}{l} HIP + HT/Machined surface \\ HIP + HT/AB surface \\ HT/machined surface \\ HT/AB surface \\ \end{tabular} & [397] \\ \hline EBM & IN718 & ASTM E647 & \begin{tabular}{l} Fatigue test at $550{ }^{\circ} \mathrm{C}$ and $2160 \mathrm{~s}$ \\ dwell time \\ \end{tabular} & 0.05 & $10 \mathrm{~Hz}$ & 1h Homogenisation/48h Homogenisation & [426] \\ \hline EBM & IN718 & ASTM E606 & Low cycle fatigue at $650^{\circ} \mathrm{C}$ & -1 & $0.5 \mathrm{~Hz}$ & \begin{tabular}{l} Specimens were built in the $0^{\circ}$ and $90^{\circ}$ \\ direction. $\mathrm{AB} / \mathrm{HIP}+\mathrm{SA}$ \\ \end{tabular} & [296] \\ \hline \multirow[t]{9}{*}{EBM} & IN718 & & \begin{tabular}{l} Four-point bending fatigue test at \\ room temperature \\ \end{tabular} & 0.1 & $20 \mathrm{~Hz}$ & \begin{tabular}{l} STA + HIP/AB surface/cross-section $10 \times$ \\ $10 \mathrm{~mm}^{2}$ (with contour) \\ \end{tabular} & [321] \\ \hline & & & & & & STA + HIP/machined surface/cross-section & \\ \hline & & & & & & $10 \times 10 \mathrm{~mm}^{2}$ (with contour) & \\ \hline & & & & & & STA + HIP/machined surface/cross-section & \\ \hline & & & & & & $6 \times 6 \mathrm{~mm}^{2}$ (without contour) & \\ \hline & & & & & & \begin{tabular}{l} STA/AB surface/cross-section $10 \times 10 \mathrm{~mm}^{2}$ \\ (with contour) \\ \end{tabular} & \\ \hline & & & & & & STA/machined surface/cross-section $10 \times$ & \\ \hline & & & & & & $10 \mathrm{~mm}^{2}$ (with contour) & \\ \hline & & & & & & \begin{tabular}{l} STA/machined surface/cross-section $6 \times 6$ \\ $\mathrm{~mm}^{2}$ (without contour) \\ \end{tabular} & \\ \hline \end{tabular} \end{center} \section*{Appendix E. Creep properties of PBF Nickel-based superalloys} Table 12 Summary of creep properties of PBF Nickel-based superalloys (' ' ' indicates that the data was obtained from the bar chart). \begin{center} \begin{tabular}{|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample condition & Ref \\ \hline LPBF & CM247LC & & Small Punch creep test at $950{ }^{\circ} \mathrm{C}$ and $150 \mathrm{~N}$ & $\mathrm{HIP}+\mathrm{STA}$ & $[412]$ \\ \hline LPBF & C263 & & Small punch testing, tested at $780^{\circ} \mathrm{C}$ & \begin{tabular}{l} Specimens are built in $0^{\circ}$ and $90^{\circ}$. Solution $(1150$ \\ $\left.{ }^{\circ} \mathrm{C}\right)+$ ageing, Solution $\left(1275^{\circ} \mathrm{C}\right)+$ ageing \\ \end{tabular} & [410] \\ \hline LPBF & Hastelloy X & & Creep test at $815^{\circ} \mathrm{C}$ & Specimens are built in $0^{\circ}, 45^{\circ}$ and $90^{\circ}$ & $[384]$ \\ \hline LPBF & IN718 & & \begin{tabular}{l} Creep test under the constant compressive stress of $725 \mathrm{MPa}$ \\ at $630^{\circ} \mathrm{C}$ \\ \end{tabular} & DA \& SA & $[275]$ \\ \hline LPBF & IN718 & & \begin{tabular}{l} Creep test under the constant compressive stress of $900 \mathrm{MPa}$ \\ at $630^{\circ} \mathrm{C}$. Stress-change tests were performed at stresses \\ between 900 and $1100 \mathrm{MPa}$. The load was step-wise \\ increased by $50 \mathrm{MPa}$. \\ \end{tabular} & \begin{tabular}{l} DA, solution at $930^{\circ} \mathrm{C}+$ ageing, solution at 1000 \\ ${ }^{\circ} \mathrm{C}+$ ageing \\ \end{tabular} & [190] \\ \hline LPBF & IN718 & ISO 204 & Creep test at $700^{\circ} \mathrm{C}$, stress range from 250 to $375 \mathrm{MPa}$ & \begin{tabular}{l} Specimens are built in $0^{\circ}$ and $90^{\circ} . \mathrm{SR}, \mathrm{SR}+$ \\ solution $\left(980^{\circ} \mathrm{C}\right)+$ ageing, SR + solution $(1065$ \\ $\left.{ }^{\circ} \mathrm{C}\right)+$ ageing \\ \end{tabular} & $[405]$ \\ \hline LPBF & IN718 & & Creep test at $650^{\circ} \mathrm{C}$ and $550 \mathrm{MPa}$ & $\mathrm{AB}, \mathrm{SA}$ & [250] \\ \hline LPBF & IN718 & & Creep test at $650^{\circ} \mathrm{C}$ and $550 \mathrm{MPa}$ & Specimens are built in $0^{\circ}$ and $90^{\circ} . \mathrm{AB}, \mathrm{SA}, \mathrm{DA}$ & $[406]$ \\ \hline LPBF & IN718 & & Creep test at $650{ }^{\circ} \mathrm{C}$ and $650 \mathrm{MPa}$ & $\mathrm{AB}, 2$ bar specimens & $[75]$ \\ \hline LPBF & IN718 & \begin{tabular}{l} ASTM \\ E139 \\ \end{tabular} & Creep test at $650^{\circ} \mathrm{C}$ and $690 \mathrm{MPa}$ & $\mathrm{AB}, \mathrm{SA}$, functionally graded built & [388] \\ \hline LPBF & IN718 & & Creep test at $650^{\circ} \mathrm{C}$ and $650 \mathrm{MPa}$ & \begin{tabular}{l} AB, HSA, CNC/WEDM machined, 2 bar \\ specimens \\ \end{tabular} & $[407]$ \\ \hline LPBF & IN718 & & Creep test at $650^{\circ} \mathrm{C}$ and $550 \mathrm{MPa}$ & \begin{tabular}{l} AB \\ Solution @ $980{ }^{\circ} \mathrm{C}$ for $1 \mathrm{~h}+$ ageing \\ Solution @ $1045^{\circ} \mathrm{C}$ for $1 \mathrm{~h}+$ ageing \\ Solution @ $1065^{\circ} \mathrm{C}$ for $1 \mathrm{~h}+$ ageing \\ Solution @ $1120^{\circ} \mathrm{C}$ for $1 \mathrm{~h}+$ ageing \\ Solution @ $1180^{\circ} \mathrm{C}$ for $1 \mathrm{~h}+$ ageing \\ Solution @ $1180^{\circ} \mathrm{C}$ for $4 \mathrm{~h}+$ ageing \\ HIP \\ HIP + ageing \\ \end{tabular} & [408] \\ \hline LPBF & IN718 & & Small punch testing, creep test at $650{ }^{\circ} \mathrm{C}$ and $600 \mathrm{~N}$ & & [411] \\ \hline \end{tabular} \end{center} Table 12 (continued) \begin{center} \begin{tabular}{|c|c|c|c|c|c|} \hline Technique & Material & Standard & Test condition & Sample condition & Ref \\ \hline & & & & \begin{tabular}{l} Forged-N: loading direction parallel to the \\ forging direction \\ Forged-P: loading direction perpendicular to the \\ forging direction \\ Casted \\ LPBF specimen: Loading direction parallel to the \\ build direction \\ \end{tabular} & \\ \hline LPBF & IN718 & \begin{tabular}{l} ASTM \\ E139 \\ \end{tabular} & $650^{\circ} \mathrm{C}$ and $620 \mathrm{MPa}$ & 1 different STA & $[326]$ \\ \hline LPBF & IN718 & \begin{tabular}{l} CEN \\ 15627 \\ \end{tabular} & Small Punch creep test at $650^{\circ} \mathrm{C}$ and $400 \mathrm{~N}$ & \begin{tabular}{l} Specimens built at $0^{\circ}$ or $90^{\circ}$ and STA or \\ Homogenisation + Ageing \\ \end{tabular} & [430] \\ \hline LPBF & IN718 & \begin{tabular}{l} ASTM \\ E139 \\ \end{tabular} & $650^{\circ} \mathrm{C}$ and $600 \mathrm{MPa}$ & \begin{tabular}{l} Specimens built at $0^{\circ}, 90^{\circ}$ and $45^{\circ}$ \\ AB/STA \\ Meander/Stripe strategy \\ \end{tabular} & $[414]$ \\ \hline LPBF & IN718 & \begin{tabular}{l} ASTM \\ E139 \\ \end{tabular} & $650{ }^{\circ} \mathrm{C}$ and $600 \mathrm{MPa}$ & \begin{tabular}{l} Specimens built at $0^{\circ}, 90^{\circ}$ and $45^{\circ}+$ STA \\ Single or Muli-laser \\ \end{tabular} & $[415]$ \\ \hline LPBF & IN738LC & ISO 204 & \begin{tabular}{l} The creep machine used was a Unisteel multipoint machine \\ with a load capacity of $30 \mathrm{kN}$ and a $20: 1$ lever ratio. Creep test \\ at $850^{\circ} \mathrm{C}$. \\ \end{tabular} & Specimens are built in $0^{\circ}$ and $90^{\circ}$ & $[249]$ \\ \hline LPBF & IN738LC & ISO204 & Creep test at $850{ }^{\circ} \mathrm{C}$. A similar procedure respect to [249] & Specimens are built in $0^{\circ}$ and $90^{\circ}$ & $[251]$ \\ \hline LPBF & \begin{tabular}{l} nickel-based \\ superalloy \\ \end{tabular} & & Creep test at $650^{\circ} \mathrm{C}$ and $550 \mathrm{MPa}$ & \begin{tabular}{l} Specimens built in $90^{\circ} /$ solution and aging \\ treatment $\left(980^{\circ} \mathrm{C} / 1 \mathrm{~h} / \mathrm{AC}+718{ }^{\circ} \mathrm{C} / 8 \mathrm{~h} / \mathrm{FC}+621\right.$ \\ $\left.{ }^{\circ} \mathrm{C} / 10 \mathrm{~h} / \mathrm{AC}\right)$ \\ Specimens built in $90^{\circ} / \mathrm{Yttrium}$ addition/ \\ solution and aging treatment $\left(980^{\circ} \mathrm{C} / 1 \mathrm{~h} /\right.$ \\ $\left.\mathrm{AC}+718^{\circ} \mathrm{C} / 8 \mathrm{~h} / \mathrm{FC}+621^{\circ} \mathrm{C} / 10 \mathrm{~h} / \mathrm{AC}\right)$ \\ Specimens built in $90^{\circ} / \mathrm{Yttrium}$ addition/direct \\ aging treatment $\left(718^{\circ} \mathrm{C} / 8 \mathrm{~h} / \mathrm{FC}+621^{\circ} \mathrm{C} / 10 \mathrm{~h} /\right.$ \\ AC) \\ Specimens built in $0^{\circ} /$ direct aging treatment \\ $\left(718{ }^{\circ} \mathrm{C} / 8 \mathrm{~h} / \mathrm{FC}+621{ }^{\circ} \mathrm{C} / 10 \mathrm{~h} / \mathrm{AC}\right)$ \\ Specimens built in $0^{\circ} / \mathrm{Yttrium} \mathrm{addition} /$ direct \\ aging treatment $\left(718{ }^{\circ} \mathrm{C} / 8 \mathrm{~h} / \mathrm{FC}+621{ }^{\circ} \mathrm{C} / 10 \mathrm{~h} /\right.$ \\ AC) \\ \end{tabular} & $[413]$ \\ \hline EBM & IN718 & & \begin{tabular}{l} Compression + Tension creep test at $800^{\circ} \mathrm{C}$. Load increased \\ stepwise in a compression test \\ \end{tabular} & SA, Specimens are built in $0^{\circ}$ and $90^{\circ}$. & $[295]$ \\ \hline EBM & IN718 & \begin{tabular}{l} ASTM \\ E319 \\ \end{tabular} & Creep test at $580 \mathrm{MPa}$ and $600 \mathrm{MPa}$, at $650^{\circ} \mathrm{C}$ & \begin{tabular}{l} Specimens are built in $0^{\circ}$ and $90^{\circ}$ with point net \\ fill scan strategy and standard melt fill scan \\ strategy, post-treated with HIP + STA \\ \end{tabular} & [409] \\ \hline \end{tabular} \end{center} \section*{Appendix F. Sample preparation} A summary of commonly used etchants for PBF manufactured nickel-based superalloys in literature is given in Table 13. \section*{Table 13} Commonly used etchants for PBF manufactured nickel-based superalloys in literature. \begin{center} \begin{tabular}{|c|c|c|} \hline Alloys & Chemical etching & Electrolytic etching \\ \hline IN718 & \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-68} & \begin{tabular}{l} - Etch in $10 \%$ oxalic acid for 5-10 s [73-75]. \\ - Etch in a mixture of $70 \mathrm{ml}$ of phosphoric acid and $30 \mathrm{ml}$ of water at RT, using $5 \mathrm{~V}$ and \\ last for 5-120 s at RT [76,77]. \\ - Etch in a solution of $12 \mathrm{ml}$ H3PO4 $+40 \mathrm{ml} \mathrm{HNO} 3+48 \mathrm{ml}$ H2SO4 at $6 \mathrm{~V}$ for $5 \mathrm{~s}$ [78]. \\ - Etch in a solution of $50 \mathrm{ml}$ lactic acid, $3 \mathrm{~g}$ oxalic acid and $150 \mathrm{ml}$ hydrochloric acid for \\ 10-20 s at a voltage of $2 \mathrm{~V}$ (DC) [79]. \\ - Etch in a solution of $50 \mathrm{ml}$ hydroxypropionic acid, $150 \mathrm{ml} \mathrm{HCl}$ acid, and $3 \mathrm{~g}$ oxalic acid \\ with a constant current of $2 \mathrm{~V}$ for a few seconds [80]. \\ \end{tabular} \\ \hline IN625 & \begin{tabular}{l} - Etch with aqua regia solution (HNO3 and $\mathrm{HCl}$ in a proportion of \\ 1:3) for $10-60 \mathrm{~s}$ [449,450]. \\ - Lactic acid-HCl mixture [451]. \\ - Etch with Kalling's No.2 etchant for $1-2 \mathrm{~s}$ [452]. \\ - Etch in $10 \mathrm{ml} \mathrm{HNO} 3,10 \mathrm{ml} \mathrm{HCl}$ and $15 \mathrm{ml} \mathrm{CH3COOH}$ for $30 \mathrm{~s}$ \\ [348] . \\ [348] \end{tabular} & \begin{tabular}{l} - Etch in $7.5 \%$ oxalic acid for approx. $10 \mathrm{~s}$ [454]. \\ - Etch in a solution of $70 \mathrm{ml}$ phosphoric acid and $30 \mathrm{ml} \mathrm{H} 2 \mathrm{O}$ using $1-5 \mathrm{~V}$ for $5 \mathrm{~s}-2 \mathrm{~min}$ at \\ RT [264,266,269,292]. \\ \end{tabular} \\ \hline \end{tabular} \end{center} Table 13 (continued) \begin{center} \begin{tabular}{|c|c|c|} \hline Alloys & Chemical etching & Electrolytic etching \\ \hline Hastelloy X & \includegraphics[max width=\textwidth]{2024_04_13_99ab48aa81fd824b9ddbg-69} & - Etch in a solution of $10 \mathrm{~g}$ oxalic acid in $100 \mathrm{ml}$ of water at $6 \mathrm{~V}$ over $8 \mathrm{~s}$ [284,455]. \\ \hline IN939 & & - Etch in $\mathrm{HF}$ at $3 \mathrm{~V}$ for $3-5 \mathrm{~s}$ [456]. \\ \hline IN738LC & \begin{tabular}{l} - Molybdic acid ( $0.5 \mathrm{~g}$ MoO3, $200 \mathrm{ml} \mathrm{H} 2 \mathrm{O}, 50 \mathrm{ml} \mathrm{HCl}$ and $50 \mathrm{ml}$ \\ HNO3) at $40^{\circ} \mathrm{C}[457]$. \\ - Adler reagent ( $50 \mathrm{ml} \mathrm{H} 2 \mathrm{O}, 100 \mathrm{ml} \mathrm{HCl}, 30 \mathrm{~g} \mathrm{FeCl} 3,6 \mathrm{~g}(\mathrm{NH} 4) 2$ \\ [CuCl4] 2) for few seconds 249,251$].$ \\ [CuCl4] \end{tabular} & \begin{tabular}{l} - Etch with $10 \%$ H2PO4 [119]. \\ - Etch in a solution of $10 \mathrm{vol} \%$ oxalic acid with $6 \mathrm{~V}$ and $0.4 \mathrm{~A}$ for $26 \mathrm{~s}$ [58]. \\ \end{tabular} \\ \hline CMSX-4 & \begin{tabular}{l} - V2A etchant at $338 \mathrm{~K}-343 \mathrm{~K}\left(65-70{ }^{\circ} \mathrm{C}\right)$ [458]. \\ - Marble's reagent $(50 \mathrm{ml} \mathrm{HCl}, 50 \mathrm{ml} \mathrm{H} 2 \mathrm{O}$ and $10.0 \mathrm{~g}$ CuSO4) \\ [459] . \\ [459] \end{tabular} & \\ \hline CM247LC & \begin{tabular}{l} - Kalling's No. 2 reagent ( $5 \mathrm{~g} \mathrm{CuCl} 2,100 \mathrm{ml} \mathrm{HCl}$, and $100 \mathrm{ml}$ \\ ethanol) [38]. \\ - Etch for $3-5 \mathrm{~s}$ in Kalling's reagent $(5 \mathrm{~g} \mathrm{CuCl} 2+100 \mathrm{ml} \mathrm{HCl}+100$ \\ ml distilled water) [184,460]. \\ \end{tabular} & \\ \hline \end{tabular} \end{center} \section*{Author statement} AC devised the underlying project alongside IA and $\mathrm{CH}$, the main conceptual ideas and proof outline. $\mathrm{ZX}$ initiated the collection of resources and first draft which were developed to completion and substantially augmented by SS, HC GG, WW, AC, CH all contributed sections to the work and provided insight to the literature and contributions to date. 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Alloys Compd. 615 (2014) 338-347. \begin{itemize} \item \end{itemize} \end{document} \documentclass[10pt]{article} \usepackage[utf8]{inputenc} \usepackage[T1]{fontenc} \usepackage{graphicx} \usepackage[export]{adjustbox} \graphicspath{ {./images/} } \usepackage{amsmath} \usepackage{amsfonts} \usepackage{amssymb} \usepackage[version=4]{mhchem} \usepackage{stmaryrd} \usepackage{hyperref} \hypersetup{colorlinks=true, linkcolor=blue, filecolor=magenta, urlcolor=cyan,} \urlstyle{same} \title{Advances in additively manufactured titanium alloys by powder bed fusion and directed energy deposition: Microstructure, defects, and mechanical behavior } \author{} \date{} %New command to display footnote whose markers will always be hidden \let\svthefootnote\thefootnote \newcommand\blfootnotetext[1]{% \let\thefootnote\relax\footnote{#1}% \addtocounter{footnote}{-1}% \let\thefootnote\svthefootnote% } %Overriding the \footnotetext command to hide the marker if its value is `0` \let\svfootnotetext\footnotetext \renewcommand\footnotetext[2][?]{% \if\relax#1\relax% \ifnum\value{footnote}=0\blfootnotetext{#2}\else\svfootnotetext{#2}\fi% \else% \if?#1\ifnum\value{footnote}=0\blfootnotetext{#2}\else\svfootnotetext{#2}\fi% \else\svfootnotetext[#1]{#2}\fi% \fi } \begin{document} \maketitle Review Article \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-01} \\ ${ }^{a}$ Centre for Advanced Materials and Manufacturing, School of Engineering, Edith Cowan University, 270 Joondalup Drive, Joondalup, Perth, WA 6027, SA \\ ${ }^{\mathrm{b}}$ School of Engineering, M050, The University of Western Australia, 35 Stirling Highway, Crawley, Perth, WA 6009, SA \\ ${ }^{\mathrm{C}}$ Institute of Metals, College of Material Science and Engineering, Changsha University of Science \& Technology, Changsha 410004, China \\ ${ }^{\mathrm{d}}$ School of Material Science and Engineering, Jiangsu University of Science and Technology, Zhenjiang 212100, China \\ e State Key Laboratory of Metal Matrix Composites, School of Material Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China\\ \} \section*{A R T I C L E I N F O} \section*{Article history:} Received 5 August 2023 Revised 5 November 2023 Accepted 8 November 2023 Available online 25 November 2023 \section*{Keywords:} Powder bed fusion Directed energy deposition Titanium alloys Phase transformation Defects Mechanical property \begin{abstract} A B S T R A C T $\mathrm{Ti}$ and its alloys have been broadly adopted across various industries owing to their outstanding properties, such as high strength-to-weight ratio, excellent fatigue performance, exceptional corrosion resistance and so on. Additive manufacturing (AM) is a complement to, rather than a replacement for, traditional manufacturing processes. It enhances flexibility in fabricating complex components and resolves machining challenges, resulting in reduced lead times for custom designs. However, owing to distinctions among various AM technologies, Ti alloys fabricated by different AM methods usually present differences in microstructure and defects, which can significantly influence the mechanical performance of built parts. Therefore, having an in-depth knowledge of the scientific aspects of fabrication and material properties is crucial to achieving high-performance Ti alloys through different AM methods. This article reviews the mechanical properties of Ti alloys fabricated by two mainstream powder-type AM techniques: powder bed fusion (PBF) and directed energy deposition (DED). The review examines several key aspects, encompassing phase formation, grain size and morphology, and defects, and provides an in-depth analysis of their influence on the mechanical behaviors of Ti alloys. This review can aid researchers and engineers in selecting appropriate PBF or DED methods and optimizing their process parameters to fabricate high-performance Ti alloys for a wide range of industrial applications. \end{abstract} (C) 2024 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science \& Technology. This is an open access article under the CC BY license (\href{http://creativecommons.org/licenses/by/4.0/}{http://creativecommons.org/licenses/by/4.0/}) \section*{1. Introduction} Titanium (Ti) and its alloys are widely utilized in various industrial applications due to their exceptional attributes, including excellent fatigue performance, outstanding strength-to-weight ratio, exceptional corrosion resistance, relatively low elastic modulus, and superior biocompatibility [1-8]. Traditional manufacturing techniques, such as casting [9-13], wrought [14], space holder technique [15], powder metallurgy [16-18], foaming [19], and rolling [20] are commonly used for producing Ti alloy parts. However, these methods still require improvement for produc- \footnotetext{\begin{itemize} \item Corresponding authors \end{itemize} E-mail addresses: \href{mailto:yjliu@csust.edu.cn}{yjliu@csust.edu.cn} (Y.J. Liu), \href{mailto:lychen@just.edu.cn}{lychen@just.edu.cn} (L.Y. Chen), \href{mailto:wang_liqiang@sjtu.edu.cn}{wang\_liqiang@sjtu.edu.cn} (L.Q. Wang), \href{mailto:l.zhang@ecu.edu.au}{l.zhang@ecu.edu.au}, \href{mailto:lczhangimr@gmail.com}{lczhangimr@gmail.com} (L.C. Zhang). } ing Ti alloys due to challenges related to shape complexity, machining difficulties, and susceptibility to oxidation, etc. [21-23]. As such, emerging additive manufacturing (AM) offers a potential supplementary approach to address these challenges in Ti alloy fabrication. AM techniques apply a layer-by-layer strategy to fabricate products from computer-aided design (CAD) models by selectively melting and solidifying raw materials [21,24], offering advantages in Ti alloy production [21,25-29]. Although AM is a well-established advanced manufacturing technology that has attracted significant attention over the past two decades and has been widely applied in various industrial sectors, its industrial adoption has been slower than initially anticipated [28,30,31]. Consequently, it is increasingly regarded as a complementary approach rather than a complete replacement for traditional manufacturing processes. Recently, considerable efforts have concentrated on AM of Ti alloys. In particular, the utilization of multi-laser AM systems, \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-02} \end{center} Fig. 1. Typical types of metal AM techniques (Reproduced with permission from Ref. [47]. Copyright (2022), Elsevier). ranging from two to eight lasers, has demonstrated the potential to substantially reduce the production time for large-sized $\mathrm{Ti}$ components, such as Ti-6Al-4V [32] and Ti-6.5Al-2Zr-Mo-V alloys [33]. Moreover, multi-material AM has exhibited remarkable versatility in meeting the demands for multifunctional parts. Nevertheless, the disparity in thermal and physical properties has given rise to challenges related to material inhomogeneity and the occurrence of interfacial defects [34-36]. For instance, Wang et al. [37] have utilized the benefits of AM technique by employing elemental mixed powder to investigate $\beta$-type Ti-35Nb alloys. Their study demonstrated a lower cost and greater availability of multipowder feedstock, resulting in a relatively low Young's modulus and moderate corrosion resistance. However, the chemical inhomogeneity in the alloy prepared using powder mixture remains a cause for concern. A recent study of Wang et al. [38] has examined the influence of microstructural and chemical inhomogeneity on the mechanical behaviors of Ti-35Nb alloys prepared using powder mixture and found that the presence of undissolved $\mathrm{Nb}$ induced inhomogeneity, which had a negative impact on the tensile ductility $(3.9 \% \pm 1.1 \%)$, while maintaining a high YS $(648 \pm 13 \mathrm{MPa})$. Using prealloyed Ti-35Nb powder can effectively address the issue of inhomogeneity, due to its stable melting process, leading to improved tensile ductility of $23.5 \% \pm 2.2 \%$, although with a slightly lower YS of $485 \pm 28 \mathrm{MPa}$ [39]. Some researchers have also investigated the corrosion behavior of AM-built Ti alloys. Dai et al. [40] have explored the corrosion resistance of AM-fabricated Ti-6Al-4V alloys in $\mathrm{NaCl}$ solution, where they have reported that AM-fabricated Ti-6Al-4V alloys exhibit poorer corrosion behavior than commercial Ti-6Al-4V (Grade 5) alloys because of the dominant $\alpha^{\prime}$ and less $\beta$ phases. Lu et al. [41] also reported that the AM-fabricated Ti-6Al-4V alloy displayed reduced film corrosion resistance compared to its conventional as-cast counterpart. However, it demonstrated enhanced passivation film stability in dynamic potential polarization curves than the as-cast alloy. Further surface laser remelting significantly improved the surface corrosion resistance of both Ti-6Al-4 V materials. Moreover, Qin et al. [42] found that the corrosion performance of the AM-fabricated $\beta$-type Ti-24Nb-4Zr-8Sn (Ti-2448) alloy was comparable to that of the wrought counterpart. Accordingly, many challenges remain in achieving a well-balanced combination of mechanical and corrosion performance in AM-fabricated Ti alloys.\\ Recently, review papers have highlighted the growing interest in the utilization of Ti alloys for the biomedical industry via AM, such as powder bed fusion (PBF) [43-46]. For instance, Zhang and Chen [44] have conducted a review of $\mathrm{L}$ (laser)-PBF-prepared Ti alloys, which have demonstrated progress in microstructure design, mechanical and corrosion behaviors. Similarly, EB (electron beam)PBF-fabricated Ti alloys with porous and solid structures have also been reviewed, focusing on their microstructure, fatigue behavior and corrosion properties [45]. Moreover, Wang et al. [21] have provided a review of the characteristics of single and multi-melt pool tracks using L-PBF, emphasizing the importance of process parameter optimization and Ti alloy powder qualification. Their review also highlights the significance of precise in-situ measurements and reliable modeling for optimizing parameters. However, there is a notable absence of detailed comparisons of the mechanical behaviors of Ti alloys produced through L-PBF, EB-PBF and directed energy deposition (DED) in the existing literature. As such, this review aims to fill this research gap by thoroughly examining the mechanical behavior of Ti alloys fabricated by two mainstream AM techniques: PBF and DED. By exploring various aspects, such as phase transformation, grain size/morphology and defects, this article provides an in-depth analysis of current advances in the mechanical behaviors of AM-fabricated Ti alloys. \section*{2. Additive manufacturing} As shown in Fig. 1, AM techniques are categorized into seven types based on ASTM F2792-12a [48]. Among those AM techniques, sheet lamination (SHL), binder jetting (BJT), directed energy deposition (DED) and powder bed fusion (PBF) are widely employed for manufacturing parts using metallic materials as the feedstock. However, compared with BJT and SHL, PBF and DED are more commonly used for metal fabrication due to two main reasons. Firstly, PBF, DED, and BJT can use metallic powder feedstock to fabricate products, while SHL only uses metal sheets as feedstock. In general, AM techniques using powder feedstock generally result in parts with higher accuracy and surface finish [49], while utilizing wires and metallic sheets as printing feedstock material usually leads to more defects, lower geometry precision, high surface roughness and limitations for the production of complex shapes [50]. Secondly, PBF and DED can directly produce a net-\\ shaped part from a computer model without additional intermediate processing steps to achieve designed shapes in contrast to BJT and SHL [28]. Hence, PBF and DED are more suitable for manufacturing metallic products with high mechanical performance, whereby they have received extensive investigation. Accordingly, this section will provide a further overview of PBF and DED. \subsection*{2.1. Powder bed fusion} PBF techniques apply electron or laser beam power to melt a selected area on each thin layer of a pre-deposited powder bed, in order to fabricate products [51-53]. Based on the types of heat sources used, PBF can be categorized into electron beam PBF (EBPBF) and laser PBF (L-PBF). As shown in Fig. 2(a), an L-PBF machine consists of the laser power section, galvanometer-driven mirrors, recoating arm, building platform, shield gas chamber, and other parts [28]. In contrast, a typical EB-PBF system (Fig. 2(b)) is similar to the L-PBF system, except for its special electron beam power section, electromagnetic coils, powder hopper and a vacuum chamber. As illustrated in Fig. 2(a), a 3D computer-aided design (CAD) model needs to be created first in STL format, which will then be sliced into printable 2D layers with a preset layer thickness $[52,54]$. In the EB-PBF process, an extra preheating step is required to lightly sinter the metallic powder feedstock, in order to avoid repulsion and electrostatic charging before the manufacturing process [28]. In contrast, preheating powder feedstock is optional in the L-PBF production process. The major processing parameters of the EB-PBF and L-PBF techniques include heating source power, scanning rate, layer thickness and hatch spacing [57]. These processing parameters directly affect the energy density, which is of vital guidance and can significantly influence the final performance of products produced by EBPBF and L-PBF. The volume energy density can be determined by Eq. (1) [58]: \begin{equation*} E_{\mathrm{V}}=\frac{P}{v \cdot t \cdot s} \tag{1} \end{equation*} where $E_{\mathrm{V}}$ is the laser or electron volumetric energy density (J $\left.\mathrm{mm}^{-3}\right), P$ is the input power $(\mathrm{W}), v$ is the scanning rate $\left(\mathrm{mm} \mathrm{s}^{-1}\right)$, $t$ is the layer thickness $(\mathrm{mm})$ and $s$ is the hatching space $(\mathrm{mm})$. Due to differences in heating sources, there are also some distinctions between the L-PBF and EB-PBF techniques. Firstly, due to the smaller laser spot in size (therefore higher manufacturing accuracy), L-PBF-fabricated products can usually achieve a higher asbuilt surface finish and precision than the products fabricated by EB-PBF [59-62]. Secondly, L-PBF technology can adopt a wider variety of engineering materials as feedstock, while only electrically conductive powder can be used as the feedstock of EB-PBF, due to the heating mode of the electron beam [28,63]. Thirdly, because of the higher input power of the electron beam and energy absorption of metallic powder, EB-PBF has a higher production rate than L-PBF [61]. \subsection*{2.2. Directed energy deposition} In addition to PBF technology, a great deal of attention has also been paid to the DED technique over the past few years. The DED system generally includes a laser heating source, a multi-axis control system and a feeding system. A typical DED system is shown in Fig. 2(c). DED uses the laser to create an active molten pool, whereby the feeding system utilizes inert gas to transfer metallic powder to the melt pool to layer-by-layer manufacture products from CAD models [28,64]. During the feedstock-adding process, inert gas is also transferred to create a shield gas atmosphere for the molten pool to prevent oxidation influences [65]. Because of differences in manufacturing principles, there are notable distinctions between DED and PBF. Firstly, DED can fabricate larger scaled parts compared to PBF techniques, where there are no limitations on the build chamber size due to the extrusionbased deposition of material [66]. Furthermore, due to its special feeding system, DED can efficiently fabricate functionally graded materials by adjusting feedstock material during the production process [67]. In addition, multi-axis nozzles used in the DED technique offer greater flexibility for material deposition from various angles, where they have become an effective tool for repairing damaged parts [65]. However, unlike the PBF system, DED lacks recoaters to level the powder feedstock before manufacturing each layer. This can lead to the accumulation of build errors in each layer, resulting in lower fabrication accuracy of DED-fabricated products compared to those produced through PBF [66,68]. A comparison of common information for these $3 \mathrm{AM}$ techniques is shown in Table 1. \section*{3. Phase transformation in AM-fabricated Ti alloys} \subsection*{3.1. Characterizing phase transformation} The crystal structure of pure Ti remains hexagonal close-packed (hcp, $\alpha$ phase) at low temperatures, but transitions to the bodycentered cubic (bcc, $\beta$ phase) structure when the temperature exceeds the allotropic transition temperature $\left(882.5^{\circ} \mathrm{C}\right)$ [76]. It has been shown that the addition of various alloying elements can significantly affect the phase transformation temperature of Ti [77]. $\alpha$-stabilizers, such as $\mathrm{Al}, \mathrm{C}$, and $\mathrm{O}$, usually can increase the $\beta / \alpha$ transit temperature, while $\beta$-stabilizers, such as Ta, Mo, and $\mathrm{Nb}$, can decrease the $\beta / \alpha$ transit temperature [78]. Ti alloys can generally be classified into three types based on the remaining phases, namely $\alpha$-type, $(\alpha+\beta)$-type and $\beta$-type Ti alloys $[21,79]$. $\alpha$-type Ti alloys consist of a single solid solution of $\alpha$ phase and usually contain various grades of commercially pure Ti (CPTi) alloys [80]. $\alpha$-type Ti alloys usually exhibit good weldability, excellent creep resistance, and outstanding corrosion resistance, making them suitable for applications in chemical engineering and high-temperature fields $[81,82]$. However, because of the brittleness and stability of the hcp structure, $\alpha$-type Ti alloys usually present low strength at room temperature, and commonly cannot be enhanced by microstructural modifications from heat treatment $[81,83]$. In contrast, $(\alpha+\beta)$-type Ti alloys contain more $\beta$ stabilizers $(\sim 4 \%-16 \%)$ and can maintain about $5-30$ vol.\% of $\beta$ phase $[78,79]$. This leads to $(\alpha+\beta)$-type Ti alloys exhibiting excellent corrosion resistance and superior strength at room temperature [84-86]. The Ti-6Al-4V alloy is one of the most widely used $\alpha+\beta$ dual-phase Ti alloys, with over half of Ti alloy usage being attributed to this alloy. As a result, it has been extensively applied in biomedical and energy industries [80,87]. Nevertheless, despite its popularity, the alloy's $\mathrm{Al}$ and $\mathrm{V}$ elements can have negative impacts on human health when used as a medical implant. Excessive $\mathrm{V}$ intake can cause dehydration, diarrhea and reduced weight gain [88], while an excess of Al can lead to Alzheimer's disease [89]. In addition, Young's modulus of Ti-6Al-4V ( 110-120 GPa) is still significantly higher than that of human bones, where the different Young's modulus between human bone and implant can induce the stress-shielding effect [90]. Therefore, novel $\beta$-Ti alloys consisting of nontoxic compositions have been considered as alternative materials of Ti-6Al-4V for medical implants. $\beta$-type Ti alloys contain the highest proportion of $\beta$-stabilizers in Ti alloys, and mainly include the $\beta$ phase at room temperature [91]. Due to the higher proportion of $\beta$ phase in $\beta$-type Ti alloys, properties such as higher toughness, plasticity and heat treatment capability can be achieved, along with a lower elastic modulus [81]. Thus, the stress shield effect can effectively be decreased by using $\beta$-type Ti alloys. Moreover, $\beta$-type Ti alloys also present better corrosion resistance, because the micro-galvanic effect between various phases can be decreased in $\beta$-type Ti alloys [92].\\ (a) 3D designed model \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-04(3)} \end{center} Computer-aided design software \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-04(1)}\\ Scan direction\\ Sliced 2D layer \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-04} \end{center} Laser Powder Bed Fusion\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-04(2)}\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-04(4)} Fig. 2. Schematic diagram of additive manufacturing: (a) L-PBF (Reproduced with permission from Ref. [21]. Copyright (2023), Elsevier), (b) EB-PBF systems (Reproduced with permission from Ref. [55]. Copyright (2021), Elsevier), and (c) DED powder feed system (Reproduced with permission from Ref. [56]. Copyright (2014), Springer Nature). Table 1 \begin{center} \begin{tabular}{llll} Comparison of L-PBF, EB-PBF and DED techniques. & & & \\ \hline AM techniques & L-PBF & EB-PBF & DED (powder based) \\ \hline Heating source & Laser beam & Electron beam & Laser beam \\ Typical Power $(\mathrm{W})$ & $\sim 100-1000$ & $\sim 3500$ & $\sim 500$ \\ Beam spot size $(\mu \mathrm{m})$ & $40-100$ & $100-200$ & $380-900$ \\ Scan speed $\left(\mathrm{mm} \mathrm{s}^{-1}\right)$ & $\sim 100-2000$ & $>1000$ & $3-15$ \\ Typical layer thickness $(\mu \mathrm{m})$ & $10-50$ & 50 & $100-250$ \\ Typical cooling rate $\left(\mathrm{K} \mathrm{s}^{-1}\right)$ & $10^{3}-10^{8}$ & $10^{3}-10^{5}$ & $10^{4}-10^{6}$ \\ Build plate temperature $\left({ }^{\circ} \mathrm{C}\right)$ & $\sim 80-200$ & $\sim 500-750$ & - \\ Building environment & Argon/Nitrogen & Vacuum & Argon \\ Preheating of powder $\left({ }^{\circ} \mathrm{C}\right)$ & - & $\sim 600-750$ & - \\ Refs. & $[69-72]$ & $[69,72,73]$ & $[70,74,75]$ \\ \end{tabular} \end{center} \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-05} \end{center} Fig. 3. Microstructure of L-PBF-produced CP-Ti alloys observed by (a) optical microscopy (OM) (Reproduced with permission from Ref. [94]. Copyright (2017), Elsevier) and (b) scanning electron microscopy (SEM) (Reproduced with permission from Ref. [95]. Copyright (2016), Elsevier), (c) transmission electron microscopy (TEM) micrograph of CP-Ti fabricated by EB-PBF (Reproduced with permission from Ref. [96]. Copyright (2015), Elsevier), and (d) OM micrograph of DED-fabricated CP-Ti alloys (Reproduced with permission from Ref. [94]. Copyright (2017), Elsevier). As mentioned above, the presence of different phases plays a critical role in determining the mechanical properties of Ti alloys [37]. In addition to alloy composition, process conditions during fabrication, such as cooling rate and build plate temperature, also have a significant impact on the phase formation and transformation of Ti alloys [21]. As a result, the phase formation of Ti alloys produced by different AM techniques can vary due to differences in the production process. In particular, the amount of nonequilibrium martensite phases is distinct in Ti alloys fabricated by different AM techniques. The formation of the martensite usually requires two essential conditions: (i) the build plate temperature is lower than the martensite start temperature; (ii) an appropriate low cooling temperature and high cooling rate must be achieved during the AM process. The required cooling rate and martensite start temperature vary depending on the chemical composition of the Ti alloy. For CP-Ti alloys, the martensite start temperature is around $850{ }^{\circ} \mathrm{C}$, where the required cooling rate should be more than 300-1000 $\mathrm{K} \mathrm{s}^{-1}$ for martensite formation [93]. As shown in Table 1, the cooling rate in all $3 \mathrm{AM}$ techniques satisfies the formation conditions of martensite in the CP-Ti alloy. However, the final phase constituents vary in CP-Ti alloys fabricated using these 3 AM techniques. As shown in Fig. 3(a), Attar et al. [94] have reported that L-PBF-fabricated CP-Ti alloy contains acicular and lathtype martensitic $\alpha^{\prime}$ phase. Sing et al. [95] have also shown that\\ L-PBF-fabricated CP-Ti alloys exhibit a microstructure consisting of mixed platelet $\alpha$ and acicular $\alpha^{\prime}$ phases, as shown in Fig. 3(b). By contrast, as shown in Fig. 3(c), only $\alpha$-lath phase is formed in CP-Ti alloy fabricated by EB-PBF [96]. This is because the manufacturing process of EB-PBF usually takes several hours at high build plate temperatures, which is almost the same as the annealing/aging process [62]. This type of phenomenon in the EB-PBF process facilitates the $\alpha^{\prime}$ phase to transform into the $\alpha$ phase [96]. Fig. 3(d) shows the DED-fabricated CP-Ti alloy, which also only contains the plate-like $\alpha$ phase [94,97]. DED requires a higher energy density than that of L-PBF in order to achieve high-density fabrication, resulting in a higher penetration (melt pool) depth during the producing process of DED [94]. Thus, compared to L-PBF, DED involves more repeated heat cycles in the melt layers, resulting in the transfer of more heat to the deposited layers, and causing the $\alpha^{\prime}$ phase to dissolve into the $\alpha$ phase. Compared with the CP-Ti alloy, the martensite start temperature of the Ti-6Al-4V is around $575^{\circ} \mathrm{C}$ [98]. Additionally, it should be noted here that for the Ti-6Al- $\mathrm{V}$, the formation of $\alpha^{\prime}$ martensite is also dependent on the cooling rate during the fabrication process. More specifically, cooling rates exceeding $683 \mathrm{~K} \mathrm{~s}^{-1}$ would result in complete $\alpha^{\prime}$ martensite formation, while cooling rates ranging from $683 \mathrm{~K} \mathrm{~s}^{-1}$ to $293 \mathrm{~K} \mathrm{~s}^{-1}$ result in incomplete $\alpha^{\prime}$ martensite formation. Cooling rates lower than $293 \mathrm{~K} \mathrm{~s}^{-1}$ would not result\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-06} Fig. 4. Microstructures of Ti-6Al-4V alloys fabricated by (a) L-PBF (Reproduced with permission from Ref. [113]. Copyright (2017), Elsevier), (b) EB-PBF (Reproduced with permission from Ref. [114]. Copyright (2015), Springer Nature), and (c, d) DED techniques (Reproduced with permission from Ref. [112]. Copyright (2019), John Wiley and Sons). in the formation of $\alpha^{\prime}$ martensite [98]. Although the cooling rate in the 3 AM techniques met the above conditions for the complete martensite formation of the Ti-6Al-4V alloy, the results were not as expected. As seen from Fig. 4(a), due to the relatively highest cooling rate, the Ti-6Al-4V alloy fabricated by L-PBF is mainly composed of fine $\alpha^{\prime}$ phase [72,95,99-103]. Moreover, in addition to the $\alpha^{\prime}$ phase, Ren et al. [104] have reported 0.1 vol.\% $\beta$ phase in L-PBF-fabricated Ti-6Al-4V samples. By contrast, as shown in Fig. 4(b), the EB-PBF-fabricated Ti-6Al-4V alloy contained dominant $\alpha$ phase and trace $\beta$ phase, where the $\alpha^{\prime}$ phase has not been noted in EB-PBF-fabricated Ti-6Al-4V samples in the literature, e.g. Refs. [72,99,101,105,106]. Neikter et al. [107] have reported on the basket-weave microstructure of mixed $(\alpha+\beta)$ phases in EBPBF-fabricated Ti-6Al-4V samples. Further works have also shown mixed $\alpha+\alpha^{\prime}$ phases occurring in the top surface of the EB-PBFfabricated Ti-6Al-4V sample, due to this area having a higher cooling rate [108]. The distinction in the phase constituent in Ti-6Al$4 \mathrm{~V}$ alloys fabricated by L-PBF and EB-PBF technologies can be attributed to different building conditions in L-PBF and EB-PBF techniques. As shown in Table 1, the elevated build plate temperature in the EB-PBF fabrication chamber, typically exceeding the martensite start temperature of the Ti-6Al- $4 \mathrm{~V}$ alloy, encourages the formation of $\alpha+\beta$ phases instead of the martensitic $\alpha^{\prime}$ phase transformation [109], and can even trigger the transformation of $\alpha^{\prime}$ phase into $\alpha$ phases, especially during prolonged printing [110]. For the DED-fabricated Ti-6Al-4V samples, the presence of the dominant fine needle-shaped $\alpha^{\prime}$ phase microstructure has been observed by Hao et al. [111]. In contrast to this, Razavi et al. [112] have reported on the coexistence of both $\alpha^{\prime}$ phase and $\alpha+\beta$ phases in Ti-6Al-4V samples fabricated by the DED method, as shown in Fig. 4(c) and (d). In addition, Zhai et al. [105] have indicated that the Ti-6Al-4V alloy fabricated by DED with low power\\ ( $330 \mathrm{~W})$ mainly contained the $\alpha^{\prime}$ phase, while a mixed microstructure of $\alpha^{\prime}$ and $\alpha+\beta$ phases occurred in the Ti-6Al-4V alloy fabricated with high power DED ( $780 \mathrm{~W})$. Moreover, Rashid et al. [106] also noted that DED-fabricated Ti-6Al-4V alloys mainly consisted of $\alpha$ and $\beta$ phases. As indicated in Table 1, DED exhibits an intermediate cooling rate, resulting in higher $\alpha^{\prime}$ martensite content in Ti-6Al-4V alloys compared to EB-PBF and lower content compared to L-PBF. Moreover, substantial thermal accumulation during DED fabrication can induce partial $\alpha^{\prime}$ phase decomposition. Due to the properties of nontoxic elements and lower elastic modulus, $\beta$-type Ti alloys have been considered as potential alternative materials for medical implants in place of Ti-6Al-4V alloys. The Ti-2448 [62], Ti-25Nb-3Zr-3Mo-2Sn (TLM) [115], Ti-35Nb-7Zr5Ta [116], Ti-35Nb-2Ta-3Zr [117-119], and Ti-35Nb [39] alloys are examples of typical $\beta$-type Ti alloys that are known for their low elastic modulus, whereby the formation of distinct phases within these alloys is also influenced by the specific AM techniques employed. For instance, it has been reported that Ti-2448 alloy fabricated using L-PBF exhibits a single $\beta$ phase $[42,62,120]$. In addition, Hernandez et al. [121] have shown that the Ti-2448 alloy fabricated by EB-PBF using a powder preheating of $250{ }^{\circ} \mathrm{C}$ contained both $\beta$ and $\alpha^{\prime \prime}$ phases, as shown in Fig. 5(a). Furthermore, Fig. 5(b) shows that the EB-PBF-fabricated Ti-2448 alloys exhibited both $\alpha$ and $\beta$ phases when the build plate heating temperature was set at $500{ }^{\circ} \mathrm{C}[62,122]$. With regard to the Ti-35Nb-7Zr-5Ta alloy, some studies have reported that the L-PBF-fabricated Ti-35Nb-7Zr-5Ta alloy only exhibits a single $\beta$ phase in the X-ray diffraction (XRD) pattern [123125]. However, as shown in Fig. 5(c) and (d), some studies have also found that the $\omega$ phase occurred in the lamellar mechanical twins in the L-PBF-fabricated Ti-35Nb-7Zr-5Ta alloy, as evident via further observation of the corresponding selected area electron\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-07} Fig. 5. Microstructures of Ti-2448 alloy produced by EB-PBF with (a) powder preheating of $250{ }^{\circ} \mathrm{C}$ (Reproduced with permission from Ref. [121]. Copyright (2013), Elsevier) and (b) build plate temperature of $500^{\circ} \mathrm{C}$ (Reproduced with permission from Ref. [122]. Copyright (2016), Elsevier); (c) lamellar mechanical twins and (d) its SAED patterns of the L-PBF-built Ti-35Nb-7Zr-5Ta alloy (Reproduced with permission from Ref. [116]. Copyright (2022), Elsevier); (e) bright field TEM image and (f) relevant SAED patterns in the DED-produced Ti-35Nb-7Zr-5Ta alloy (Reproduced with permission from Ref. [127]. Copyright (2006), John Wiley and Sons). diffraction (SAED) pattern and transmission electron microscopy (TEM) [116,126]. Compared with the Ti-2448 alloy, the Ti-35Nb7Zr-5Ta alloy has a higher amount of $\beta$ stabilizers (Nb), which hinders the $\alpha^{\prime \prime}$ phase formation due to the high percentage of $\beta$ stabilizers. Accordingly, only the $\omega$ phase formed in L-PBF-fabricated Ti-35Nb-7Zr-5Ta alloys. Similarly, Wang et al. [39] have reported that Ti-35Nb alloys fabricated using L-PBF with prealloyed powder primarily exhibit metastable $\beta$ phase, $\alpha^{\prime \prime}$ phase and trace nano $\omega$ phase, while the alloys fabricated with mixed powder showed a combination of $\beta+\alpha$ phases and trace amounts of $\alpha^{\prime \prime}$ phase. This difference in microstructure can be attributed to the presence of many undissolved $\mathrm{Nb}$ particles in the mixed powder sample, resulting in their distinct microstructure during the fabrication process. In addition, as shown in Fig. 5(e) and (f), Banerjee et al. [127] have claimed that the DED-fabricated Ti-35Nb-7Zr-5Ta alloy contains the $\beta$ phase and nanometer-scale $\omega+\alpha$ phases. Similarly, Nartu et al. [128] have investigated the Ti-35Nb-7Zr-5Ta alloy fabricated by DED with different laser powers, finding that only a single $\beta$ phase occurred in the XRD pattern of Ti-35Nb-7Zr-5Ta alloy fabricated by DED under $400 \mathrm{~W}$ laser power, while both $\alpha$ and $\beta$ phases presented in the XRD pattern of DED-fabricated Ti-35Nb7Zr-5Ta alloy with 500 and $600 \mathrm{~W}$ laser power [128]. According to Gu et al. [129], linear energy density (LED), which is calculated as the laser power $(P)$ divided by the scanning speed $(v)$, is a critical factor in determining heat input during the melting process. In a study conducted by Nartu et al. [128], the DED process with different laser powers of 400, 500 and $600 \mathrm{~W}$ resulted in LED values of $31.5,39.4$ and $47.2 \mathrm{~J} / \mathrm{mm}$ respectively, which were much higher than the LED of L-PBF $(0.16 \mathrm{~J} / \mathrm{mm})$ reported in the study by Luo et al. [116]. This suggests that there is a larger heat accumulation during the process of DED. Moreover, the higher heat accumulation of DED at 500 and $600 \mathrm{~W}$ laser powers, compared to DED at $400 \mathrm{~W}$, prolongs the cooling time of each deposited layer, providing favorable conditions for the diffusion transition from the $\beta$ phase to the $\alpha$ phase. \subsection*{3.2. Influence of phase transformation on mechanical properties} Table 2 shows the mechanical properties of Ti alloys fabricated by 3 AM techniques. L-PBF-fabricated Ti alloys typically exhibit relatively higher hardness and strength, but lower ductility compared to Ti alloys fabricated by EB-PBF and DED. This can be ascribed to the formation of more non-equilibrium phases in L-PBF-fabricated Ti alloys. According to Attar et al. [130], compared with $\alpha$ phase, there is a higher density of dislocations, increased amount of closely spaced interfaces and separation of adjacent laths and plates in the martensite $\alpha^{\prime}$ phase. In addition, the martensitic $\alpha^{\prime}$ phase also can result in the lattice structure deformation in Ti alloys [59]. Therefore, the dislocation slip can be effectively hindered by the martensitic structure during deformation, Table 2 Mechanical and tensile properties of various Ti alloys fabricated by different AM techniques. \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|} \hline Material & Method & $\sigma_{0.2}(\mathrm{MPa})$ & $\sigma_{\text {UTS }}(\mathrm{MPa})$ & $H_{\mathrm{v}}$ & $E(\mathrm{GPa})$ & $\varepsilon(\%)$ & Refs. \\ \hline CP-Ti & L-PBF & $555 \pm 3$ & $757 \pm 12.5$ & 261 & - & $19.5 \pm 1.8$ & [130] \\ \hline CP-Ti & L-PBF & $620 \pm 20$ & $703 \pm 16$ & $213 \pm 10$ & $112 \pm 3$ & $5.2 \pm 0.3$ & [95] \\ \hline CP-Ti & EB-PBF & $377 \pm 10$ & $475 \pm 15$ & - & - & $28.5 \pm 0.5$ & [96] \\ \hline CP-Ti & DED & $518 \pm 5$ & $640 \pm 6$ & 241 & $126 \pm 3$ & 29 & [97] \\ \hline Ti-6Al-4V & L-PBF & $1056 \pm 64$ & $1166 \pm 107$ & $383 \pm 11$ & $132 \pm 16$ & $6.1 \pm 2.6$ & [95] \\ \hline Ti-6Al-4V & L-PBF & 1036 & 1086 & - & - & 10.4 & [104] \\ \hline Ti-6Al-4V & L-PBF & 1110 & 1267 & 409 & 109 & 7.3 & $[100]$ \\ \hline Ti-6Al-4V & L-PBF & $1143 \pm 30$ & $1219 \pm 20$ & - & - & $4.9 \pm 0.6$ & [101] \\ \hline Ti-6Al-4V & L-PBF & $1150 \pm 67$ & $1246 \pm 134$ & - & - & $1.4 \pm 0.5$ & $[72]$ \\ \hline Ti-6Al-4V & L-PBF & $962 \pm 47$ & $1166 \pm 25$ & - & - & $1.7 \pm 0.3$ & [102] \\ \hline Ti-6Al-4V & L-PBF & 736 & 1051 & 360 & 110 & 11.9 & [103] \\ \hline Ti-6Al-4V & L-PBF & $960 \pm 12$ & $1042 \pm 16$ & - & - & $9.3 \pm 1.04$ & [134] \\ \hline Ti-6Al-4V & L-PBF & 1234 & 1286 & - & - & 5.22 & [135] \\ \hline Ti-6Al-4V & EB-PBF & 645 & 778 & - & - & 12.2 & [104] \\ \hline Ti-6Al-4V & EB-PBF & $869 \pm 7.2$ & $928 \pm 9.8$ & 311 & - & $9.9 \pm 1.7$ & [101] \\ \hline Ti-6Al-4V & EB-PBF & 1001 & 1073 & - & - & 11 & [105] \\ \hline Ti-6Al-4V & EB-PBF & $890 \pm 15$ & $930 \pm 22$ & 347 & 125 & 11 & [106] \\ \hline Ti-6Al-4V & EB-PBF & $846 \pm 7$ & $976 \pm 11$ & - & - & $15 \pm 2$ & $[72]$ \\ \hline Ti-6Al-4V & EB-PBF & $830 \pm 5$ & $910 \pm 10$ & 328 & $118 \pm 5$ & - & [136] \\ \hline Ti-6Al-4V & EB-PBF & 735 & 775 & 377 & 93 & 2.3 & [137] \\ \hline Ti-6Al-4V & EB-PBF & $13.2-16.3$ & $944.5-964.5$ & 330 & - & 14.6 & [138] \\ \hline Ti-6Al-4V & DED (low power) & 1005 & 1103 & - & - & 4 & [105] \\ \hline Ti-6Al-4V & DED (high power) & 990 & 1042 & - & - & 7 & [105] \\ \hline Ti-6Al-4V & DED & $1020 \pm 13$ & $1100 \pm 9$ & 395 & 125 & 8 & [106] \\ \hline Ti-6Al-4V & DED & $976 \pm 24$ & $1099 \pm 2$ & $360 \pm 10$ & - & $4.9 \pm 0.1$ & [139] \\ \hline Ti-6Al-4V & DED & $916 \pm 26$ & $1032 \pm 31$ & - & $113 \pm 5$ & $19 \pm 4$ & [140] \\ \hline Ti-6Al-4V & DED & $950 \pm 2$ & $1025 \pm 2$ & - & - & $5 \pm 1$ & [141] \\ \hline Ti-6Al-4V & DED & $945 \pm 13$ & $1041 \pm 12$ & - & - & $14.5 \pm 1.2$ & [142] \\ \hline Ti-2448 & L-PBF & $563 \pm 38$ & $665 \pm 18$ & $220 \pm 6$ & $53 \pm 1$ & $13.8 \pm 4.1$ & [120] \\ \hline Ti-2448 & L-PBF & $490 \pm 16$ & $700 \pm 6$ & $219 \pm 8$ & $49 \pm 1$ & $22 \pm 1$ & [143] \\ \hline Ti-2448 & EB-PBF & - & - & 255.1 & - & - & [121] \\ \hline Ti-25Nb-3Zr-3Mo-2Sn & L-PBF & $592 \pm 21$ & $716 \pm 14$ & - & - & $37 \pm 5$ & [115] \\ \hline Ti-35Nb-7Zr-5Ta & L-PBF & $569.5-668.5$ & $575.1-675.0$ & - & $76.7-85.26$ & $26.2-31.7$ & [125] \\ \hline Ti-35Nb-7Zr-5Ta & L-PBF & $816 \pm 26$ & $830 \pm 26$ & - & $66.5 \pm 1.5$ & $16.5 \pm 1.8$ & [116] \\ \hline Ti-35Nb-7Zr-5Ta & L-PBF & 309 & 631 & $205 \pm 10$ & 81 & $14-15$ & $[124]$ \\ \hline Ti-35Nb-7Zr-5Ta & DED & 813.6 & 834.2 & - & 55 & 19 & [127] \\ \hline Ti-35Nb-2Ta-3Zr & L-PBF & $\sim 430$ & 552 & - & - & 21 & [117] \\ \hline $\mathrm{Ti}-35 \mathrm{Nb}$ & L-PBF & $648 \pm 13$ & $803 \pm 33$ & $274 \pm 7$ & $84 \pm 2$ & $3.9 \pm 1.1$ & $[38]$ \\ \hline $\mathrm{Ti}-35 \mathrm{Nb}$ & L-PBF & $636 \pm 80$ & $750 \pm 51$ & $241 \pm 14$ & $85 \pm 2$ & $2.2 \pm 1.4$ & $[39]$ \\ \hline $\mathrm{Ti}-35 \mathrm{Nb}$ & L-PBF & $485 \pm 28$ & $645 \pm 9$ & $174 \pm 7$ & $72 \pm 1$ & $23.5 \pm 2.2$ & $[39]$ \\ \hline \end{tabular} \end{center} $\sigma_{0.2}$ - tensile yield strength; UTS - ultimate tensile strength; $H_{\mathrm{V}}$ - Vickers hardness; $E$ - elastic modulus and $\varepsilon$ - tensile elongation. which results in dislocation strengthening and improved strength of Ti alloys [131-133]. However, the delayed dislocation movement can negatively impact the ductility of the material. Zhao et al. [99] have reported that the martensitic $\alpha^{\prime}$ phase has high strength and low ductility, which is associated with Ti alloys containing the $\alpha^{\prime}$ phase. Therefore, it is reasonable that L-PBF-fabricated Ti alloys present relatively higher strength and hardness, but lower ductility in comparison to those produced by EB-PBF and DED methods. \section*{4. Grain size and morphology in AM-fabricated Ti alloys} \subsection*{4.1. Understanding grain size} In addition to variations in phase distinctions, the grain size in Ti alloys fabricated by different AM techniques also varies. In general, grain size is mainly influenced by growth time (i.e., solidification rate) [21], where longer growth time leads to the formation of coarser grain, while shorter growth time typically results in fine grains. However, accurately predicting the growth time of grains in Ti alloys is a complex task due to several reasons. Firstly, different types of grains in Ti alloys nucleate and grow within different temperature ranges. For example, $\beta$ grains usually form when the temperature is below the melting point and above the allotropic transition temperature, so the growth time of $\beta$ grains in Ti alloys depends on the dwell time between the allotropic transition temperature and the melting point [144]. In contrast, $\alpha$ grains form when the temperature is below the allotropic transition tempera- ture, whereby the growth time of $\alpha$ grains is mainly impacted by the cooling rate [107]. Moreover, the final grain size in Ti alloys varies due to differences in the manufacturing processes employed by different AM techniques. Currently, the literature on Ti-6Al-4V alloy fabricated by AM techniques is most abundant, which provides adequate evidence for comparing the grain sizes of Ti alloys. Accordingly, the following sections will present the Ti-6Al-4V alloy as an example to discuss size distinctions in grain sizes among $\mathrm{Ti}$ alloys fabricated by different AM techniques. \subsection*{4.1.1. Prior $\beta$ grain size} As discussed above, when the temperature drops below the allotropic transition temperature, $\beta$ grains will partially decompose into $\alpha / \alpha^{\prime}$ grains in Ti-6Al-4V alloys, based on different producing conditions. Thus, the $\beta$ grain is commonly referred to as the "prior $\beta$ grain". Usually, the Ti-6Al-4V alloy fabricated by the DED process tends to exhibit larger prior $\beta$ grains compared to those produced by L-PBF and EB-PBF techniques. Fig. 6(a) presents electron backscatter diffraction (EBSD) images of the horizontal layer (scan plane) in the L-PBF-fabricated Ti-6Al-4V alloy, showcasing a considerably homogeneous microstructure characterized by a basketweave and needle-like morphology of $\alpha / \alpha^{\prime}$ phases, with an average width of less than $10 \mu \mathrm{m}$. These phases undergo a phase transformation from the prior $\beta$ grains. The corresponding reconstructed EBSD microstructure of the prior $\beta$ grains is depicted in Fig. 6(b), revealing extensive variation in the size of $\beta$ grains with different orientations, averaging around $40-150 \mu \mathrm{m}$. Simonelli et al.\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-09(1)}\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-09} Fig. 6. EBSD images for horizontal layers (scan plane) of Ti-6Al-4V alloys fabricated using different methods: (a) by L-PBF and (b) corresponding reconstructed parent prior $\beta$ grains (Reproduced with permission from Ref. [147]. Copyright (2023), Elsevier), and (c) by EB-PBF, small circles in the IPF at the lower right corner depict the parent $\beta$ grains in the scan plane, with the labeled regions (1)-(7) representing these parent $\beta$ grains (Reproduced with permission from Ref. [148]. Copyright (2020), Elsevier); (d) three dimensional optical images of DED-processed Ti-6Al-4V showing prior $\beta$ grains and $\alpha$ laths (Reproduced with permission from Ref. [149]. Copyright (2021), Elsevier), and (e) results of the $\beta$ grain width and its evolution trend with different laser powers (Reproduced with permission from Ref. [150]. Copyright (2021), Elsevier). [145] have also reported a prior $\beta$ grain thickness of $103 \mu \mathrm{m}$ in L-PBF-fabricated Ti-6Al-4V alloys. Fig. 6(c) displays the EBSD orientation map of the scan plane for the Ti-6Al-4V alloy produced using EB-PBF. The microstructure reveals basket-weave $\alpha$ laths/plates distributed through the microstructure. The labeled regions (1)-(7) represent the prior parent $\beta$ grains, which exhibit size variations ranging from approximately $50 \mu \mathrm{m}$ to $200 \mu \mathrm{m}$. Similarly, Lancaster et al. [146] have found the thickness size of prior $\beta$ grains to be around $246 \mu \mathrm{m}$ in EB-PBF-fabricated Ti-6Al-4V alloys. Fig. 6(d) presents the three-dimensional (3D) optical images of the Ti-6Al$4 \mathrm{~V}$ counterpart fabricated using DED, where the images reveal the presence of columnar prior $\beta$ grains along the build direction, with an average length ranging from $1.1 \pm 0.4 \mathrm{~mm}$ to $1.7 \pm 0.8 \mathrm{~mm}$. When viewed from the top of the scan plane, equiaxed $\beta$ grains of similar size can be observed, with an average size ranging from $280 \mu \mathrm{m}$ to $510 \mu \mathrm{m}$, with increased laser heat input density from $141 \mathrm{~J} / \mathrm{mm}$ to $283 \mathrm{~J} / \mathrm{mm}$. At higher magnification, the images showcase the presence of needle-like $\alpha$-laths inside the prior $\beta$ grains, with widths ranging from $0.7 \mu \mathrm{m}$ to $1.7 \mu \mathrm{m}$. Fig. 6(e) provides additional insights into the effect of laser power on the width of $\beta$ grains. The results indicate that as the laser power increases, the width of the $\beta$ grains also increases. More specifically, the width rises from $140 \mu \mathrm{m}$ at $700 \mathrm{~W}$ to $160 \mu \mathrm{m}$ at $850 \mathrm{~W}$, and then jumps to $200 \mu \mathrm{m}$ at $1000 \mathrm{~W}$. Moreover, this variation in $\beta$ grain width\\ Table 3 Prior $\beta$ grain size of Ti-6Al-4V alloys fabricated by various AM technologies. \begin{center} \begin{tabular}{lllll} \hline Material & Method & Grain types & Grain width $($ thickness) $(\mu \mathrm{m})$ & Refs. \\ \hline Ti-6Al-4V & L-PBF & prior $\beta$ & 87 & $[107]$ \\ Ti-6Al-4V & L-PBF & prior $\beta$ & $40-150$ & $[147]$ \\ Ti-6Al-4V & L-PBF & prior $\beta$ & 103 & $[145]$ \\ Ti-6Al-4V & EB-PBF & prior $\beta$ & 93 & $[107]$ \\ Ti-6Al-4V & EB-PBF & prior $\beta$ & $50-200$ & $[148]$ \\ Ti-6Al-4V & EB-PBF & prior $\beta$ & 246 & $[146]$ \\ Ti-6Al-4V & DED & prior $\beta$ & 202 & $[107]$ \\ Ti-6Al-4V & DED & prior $\beta$ & $280-510$ & $[149]$ \\ Ti-6Al-4V & DED & prior $\beta$ & $140-200$ & $[150]$ \\ Ti-6Al-4V & DED & prior $\beta$ & $200-400$ & $[151]$ \\ Ti-6Al-4V & DED & prior $\beta$ & $300 \pm 60$ & $[152]$ \\ \end{tabular} \end{center} has a slight impact on the volume fraction of the basket-weave $\alpha$ phase. Table 3 provides a summary of the prior $\beta$ grain sizes in Ti-6Al$4 \mathrm{~V}$ alloys produced through different AM methods. As indicated in the table, DED-manufactured Ti-6Al-4V alloys typically exhibit larger prior $\beta$ grains, while those produced via L-PBF have smaller prior $\beta$ grains. EB-PBF-fabricated Ti-6Al-4V alloys fall in between these two categories.\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-10} Fig. 7. Thermal history of a single layer of Ti-6Al-4V during the production process of (a) L-PBF (Reproduced with permission from Ref. [153]. Copyright (2020), John Wiley and Sons), (b) EB-PBF (Reproduced with permission from Ref. [154]. Copyright (2021), Elsevier), and (c) DED (Reproduced with permission from Ref. [155]. Copyright (2021), Elsevier) techniques. In the case of AM techniques, which employ a unique layerby-layer production process, the thermal history becomes significantly more complex. This is ascribed to the transfer and accumulation of heat within the deposited layers from newly added layers. Consequently, this phenomenon can significantly affect dwell time within the temperature range required for the nucleation and growth of $\beta$ grains. Fig. 7 shows the thermal history of a single layer of Ti-6Al-4V during different AM techniques. As depicted in Fig. 7(c), DED provides an extended dwell time in the temperature range, conducive to $\beta$ grain growth. On the one hand, the temperature peak reaches the $\beta$ grain growth temperature zone over five times within a single layer. On the other hand, as shown in Fig. 7(a) and (b), the temperature peaks in L-PBF and EB-PBF only reach the $\beta$-grain growth temperature region about three times. The discrepancy can be attributed to the higher penetration depth of DED, which has been discussed in detail in Section 3.1. Therefore, the longer dwell time achieved by DED technology allows for adequate growth time, resulting in larger sizes of prior $\beta$ grains in Ti-6Al-4V alloy compared to counterparts fabricated by L-PBF and EB-PBF.\\ Table 4 $\alpha / \alpha^{\prime}$ lath thickness of Ti-6Al-4V alloys fabricated by different AM technologies. \begin{center} \begin{tabular}{lllll} \hline Material & Method & Grain types & Grain width (thickness) $(\mu \mathrm{m})$ & Refs. \\ \hline Ti-6Al-4V & L-PBF & $\alpha / \alpha^{\prime}$ & $0.62-1.32$ & $[160]$ \\ Ti-6Al-4V & L-PBF & $\alpha / \alpha^{\prime}$ & 0.5 & $[161]$ \\ Ti-6Al-4V & L-PBF & $\alpha / \alpha^{\prime}$ & $0.37 \pm 0.10$ & $[156]$ \\ Ti-6Al-4V & L-PBF & $\alpha / \alpha^{\prime}$ & $0.2-1$ & $[72]$ \\ Ti-6Al-4V & L-PBF & $\alpha / \alpha^{\prime}$ & $0.2-1.2$ & $[158]$ \\ Ti-6Al-4V & EB-PBF & $\alpha / \alpha^{\prime}$ & $3.61 \pm 0.92$ & $[156]$ \\ Ti-6Al-4V & EB-PBF & $\alpha / \alpha^{\prime}$ & 1.7 & $[107]$ \\ Ti-6Al-4V & EB-PBF & $\alpha / \alpha^{\prime}$ & $1-1.2$ & $[162]$ \\ Ti-6Al-4V & EB-PBF & $\alpha / \alpha^{\prime}$ & 1.9 & $[163]$ \\ Ti-6Al-4V & EB-PBF & $\alpha / \alpha^{\prime}$ & 1.4 & $[99]$ \\ Ti-6Al-4V & EB-PBF & $\alpha / \alpha^{\prime}$ & $2.09-2.14$ & $[164]$ \\ Ti-6Al-4V & DED & $\alpha / \alpha^{\prime}$ & $0.75-0.8$ & $[165]$ \\ Ti-6Al-4V & DED & $\alpha / \alpha^{\prime}$ & $0.7-1.7$ & $[149]$ \\ Ti-6Al-4V & DED & $\alpha / \alpha^{\prime}$ & 1.81 & $[166]$ \\ Ti-6Al-4V & DED & $\alpha / \alpha^{\prime}$ & 0.7 & $[158]$ \\ Ti-6Al-4V & DED & $\alpha / \alpha^{\prime}$ & 0.8 & \\ \hline \end{tabular} \end{center} \subsection*{4.1.2. $\alpha / \alpha$ ' lath thickness} The purpose of this section is to offer a macroscopic understanding of grain size variations in Ti alloys produced by different AM methods by comparing the sizes of $\alpha$ and $\alpha^{\prime}$ grains, collectively referred to as $\alpha / \alpha^{\prime}$ grains. Typically, the unique principles inherent to each AM technique can lead to various sizes of $\alpha / \alpha^{\prime}$ grains in Ti-6Al-4V alloys fabricated by different AM methods. Some researchers have conducted a comparison of the grain size and morphology between Ti-6Al-4V processed through L-PBF and EB-PBF $[156,157]$. Fig. 8(a) presents the EBSD inverse pole figure (IPF) of L-PBF fabricated Ti-6Al-4V, revealing a basketweave needleshaped $\alpha$ ' morphology with a width (thickness) of $0.37 \pm 0.10 \mu \mathrm{m}$ (Fig. 8(c)). In contrast, Fig. 8(b) demonstrates that the EB-PBF fabricated counterpart exhibited a Widmanstätten $\alpha$ morphology with an average thickness of $3.61 \pm 0.92 \mu \mathrm{m}$ (Fig. 8(c)). Fig. 8(d) and (e) shows the $\alpha$ lath displayed an average size of $1.81 \mu \mathrm{m}$ in DED-fabricated Ti-6Al-4V. Lee et al. [158] further conducted a comparison of the needle-like $\alpha / \alpha^{\prime}$ lath between Ti-6Al-4V processed through L-PBF and DED, where in L-PBF, the $\alpha^{\prime}$ laths exhibited a width ranging from $0.2 \mu \mathrm{m}$ to $1.2 \mu \mathrm{m}$ and a length ranging from $2 \mu \mathrm{m}$ to $12 \mu \mathrm{m}$. Contrasted to this, the DED counterpart displayed $\alpha$ laths with a width of $0.8 \mu \mathrm{m}$ and a length of $20 \mu \mathrm{m}$. Table 4 lists the $\alpha / \alpha^{\prime}$ lath thicknesses reported in recent studies for Ti-6Al-4V alloys fabricated using 3 AM techniques. Based on the results of these studies, it is apparent that the $\alpha / \alpha^{\prime}$ grains generally attain a larger size in EB-PBF-fabricated Ti-6Al-4V alloys than those counterparts fabricated by L-PBF and DED. In addition, the size of $\alpha / \alpha^{\prime}$ grains of DED-fabricated Ti-6Al-4V alloys is slightly bigger than that of L-PBF-fabricated Ti-6Al-4V alloys. \subsection*{4.1.3. Influence of grain size on mechanical properties} There is also a significant connection between grain size and the mechanical properties of Ti alloys fabricated by AM techniques. As evident in Table 2, Ti-6Al-4V alloy samples fabricated by LPBF and DED present relatively higher YS and ultimate tensile strength (UTS) compared to those fabricated by EB-PBF. In addition, the highest hardness is observed in Ti-6Al-4V parts fabricated by L-PBF, whereas EB-PBF-fabricated forms present relatively lower hardness when compared to counterparts fabricated by the other 2 AM methods. The improved strength and hardness in Ti alloys fabricated by L-PBF and DED can be attributed to their finer $\alpha / \alpha^{\prime}$ grain size. According to Whang [167], both the strength and hardness of metallic materials generally increase with decreasing grain size, with the maximum strengthening effect observed when grains are refined to 20-30 nm (through the Hall-Petch relationship). The relation between grain size and mechanical properties can be repre- \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-11} \end{center} Fig. 8. EBSD IPF images of Ti-6Al-4V (Ti-64) along the build direction for (a) L-PBF, (b) EB-PBF, and (c) corresponding lath size distribution (Reproduced with permission from Ref. [156]. Copyright (2022), Elsevier); (d) DED and (e) corresponding lath size distribution histograms (Reproduced with permission from Ref. [159]. Copyright (2022), Elsevier). sented by a Hall-Petch equation [168]: \begin{equation*} H=H_{0}+k_{\mathrm{H}} \delta_{\alpha_{\text {lath }}}^{-\frac{1}{2}} \tag{2} \end{equation*} where $H$ is the hardness or YS, $\delta_{\alpha \text { lath }}$ is the $\alpha$ lath thickness, $H_{0}$ and $k_{\mathrm{H}}$ are constants. It can be seen from Eq. (2) that as grain size decreases, both the YS and hardness of the material are enhanced. This relationship has been supported by the findings of Xiao et al. [169], examining the relationship between grain size and YS of Ti$6 \mathrm{Al}-4 \mathrm{~V}$ using EB-PBF and L-PBF, and concluding that most of the data align with the Hall-Petch theory. Moreover, Galarraga et al. [73] have evaluated the influence of $\alpha$ lath thickness on the mechanical properties of Ti-6Al-4V alloys, showing that increasing the lath thickness of the $\alpha$ grain leads to a slight reduction in UTS, while the YS decreases significantly. Furthermore, the hardness of the Ti-6Al-4V alloy was decreased with an increase in $\alpha$ lath thickness. The relationship between $\alpha$ lath thickness and mechanical properties of Ti-6Al-4V is shown in Fig. 9. As can be seen from Fig. 9(a), the influence of the average $\alpha$ lath width on the tensile YS of L-PBF fabricated Ti-6Al-4V is depicted. The curve represent- ing the Hall-Petch relationship demonstrates a good fit between the width and strength (red symbols), consistent with other findings in the literature. Fig. 9(b) and (c) displays the Hall-Petch relationship between the average width of $\alpha$ lath and the average width of prior $\beta$ grain on the Vickers hardness of EB-PBF fabricated Ti-6Al-4V alloys. Fig. 9(b) demonstrates a significant association between the increase in the average width of $\alpha$ lath and corresponding decrease in Vickers hardness. The fitted curve coefficient $\left(R^{2}=0.91\right)$ applies to both the fabricated rectangular plate (RP) and round bar (RB), indicating a robust relationship between the mean $\alpha$ lath width and hardness in a single fitting line. On the one hand, the significant correlation observed between the average width of $\alpha$ lath and hardness highlights the potential of the average width of $\alpha$ lath as a promising microstructural factor for fine-tuning the mechanical properties of EB-PBF-built Ti-6Al-4V. On the other hand, Fig. 9(c) examines the effect of the average width of prior $\beta$ grain on the hardness of the fabricated RP and RB Ti-6Al-4V samples. The data reveal two separate fitting curves for the fabricated RB and RP samples, with an $R^{2}$ value of $\sim 0.86$.\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-12} Fig. 9. (a) Influence of $\alpha$ lath thickness (width) on the YS of L-PBF Ti-6Al-4V (Reproduced with permission from Ref. [170]. Copyright (2019), Elsevier), influence of $\alpha$ lath width (b) and prior $\beta$ grain width (c) on the hardness of EB-PBF Ti-6Al-4V (Reproduced with permission from Ref. [171]. Copyright (2019), Elsevier), and (d) effect of $\alpha$ lath size on YS of DED Ti-6Al-4V (Reproduced with permission from Ref. [150]. Copyright (2021), Elsevier). This $R^{2}$ value is noticeably lower than the $R^{2}$ value of 0.91 for the mean $\alpha$ lath width. Consequently, the results suggest that prior $\beta$ grain width cannot be considered as an important microstructural indicator, and that controlling the $\alpha$ lath width is the primary factor for determining the mechanical properties of EB-PBF-built Ti-6Al-4V. Fig. 9(d) presents a comparison of the influence of $\alpha$ lath width on the tensile YS of DED-fabricated Ti-6Al-4V. The data points were fitted to a Hall-Petch relationship, whereby the results demonstrate that the variation in $\alpha$ lath width correlates well with the YS, with an $R^{2}=0.97937$. In contrast, the fitting of the HallPetch YS curve for $\beta$ grain width (Figure not shown here) only yielded an $R^{2}=0.34831$. These findings suggest that the width of the $\alpha$ lath plays a more important role in affecting the mechanical properties of DED-fabricated Ti-6Al-4V compared to the width of the $\beta$ grains. Moreover, the fitting results in Fig. 9 affirm the reliability of the Hall-Petch relationship in predicting YS and hardness, which can be directly applied to Ti-6Al-4V alloys produced using $\mathrm{L}-\mathrm{PBF}, \mathrm{EB}-\mathrm{PBF}$, and DED, despite the variations in grain sizes and morphologies resulting from these different methods. This mechanism can be explained by dislocation theory. In crystalline materials, smaller grains create a larger amount of grain boundaries, leading to grain-boundary strengthening. The mechanism of grain-boundary strengthening can be explained with re- gard to two aspects. Firstly, because of the different lattice orientations in neighboring grains, the grain boundaries act as pinning points that prevent the further propagation of dislocations. Secondly, grain boundaries exhibit a high degree of disorder and consist of a large number of dislocations, which also obstruct the continuous movement of dislocations within slip planes. Owing to the limited movement of dislocations, plastic deformation is delayed, thereby enhancing YS and hardness. In addition, fine grains promote a more uniform stress distribution and relieve stress concentration. Therefore, smaller grain size contributes to increased fracture toughness [172], which can result in higher UTS. \subsection*{4.2. Grain morphology in Ti alloys} Apart from grain size, the mechanical properties of AMfabricated Ti alloys also can be influenced by grain morphology. The two most common grain morphologies observed in Ti alloys are columnar grains and equiaxed grains, which can occur under different conditions during the production process $[21,173]$. The formation of columnar grains involves the epitaxial growth from parent grains in the substrate across multiple layers and melt pools. In contrast, equiaxed grains form through nucleation and growth within each individual melting pool. Among the mi-\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-13} Fig. 10. Columnar prior $\beta$ grains formed in Ti-6Al-4V alloys produced by (a) L-PBF (Reproduced with permission from Ref. [40]. Copyright (2016), Elsevier), (b) EB-PBF (Reproduced with permission from Ref. [177]. Copyright (2019), Elsevier), and (c) DED (Reproduced with permission from Ref. [178]. Copyright (2020), Nature Portfolio). Table 5 Mechanical tensile properties of Ti-6Al-4V alloys fabricated by $3 \mathrm{AM}$ techniques with different building orientations at room temperature. \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|} \hline Material & Method & Condition & Tensile axis orientation & $\sigma_{0.2}(\mathrm{MPa})$ & $\sigma_{\text {UTS }}(\mathrm{MPa})$ & $\varepsilon(\%)$ & Refs. \\ \hline Ti-6Al-4V & L-PBF & As-built & Horizontal & $1075 \pm 25$ & $1199 \pm 49$ & $7.6 \pm 0.5$ & $[184]$ \\ \hline Ti-6Al-4V & L-PBF & As-built & Vertical & $967 \pm 10$ & $1117 \pm 3$ & $8.9 \pm 0.4$ & [184] \\ \hline Ti-6Al-4V & L-PBF & As-built & Horizontal & $1070 \pm 50$ & $1250 \pm 50$ & $5.5 \pm 1$ & [185] \\ \hline Ti-6Al-4V & L-PBF & As-built & Vertical & $1050 \pm 40$ & $1180 \pm 30$ & $8.5 \pm 1.5$ & [185] \\ \hline Ti-6Al-4V & EB-PBF & Machined & Horizontal & 1063 & - & 7.1 & [163] \\ \hline Ti-6Al-4V & EB-PBF & Machined & Vertical & 997 & - & 8.8 & [163] \\ \hline Ti-6Al-4V & EB-PBF & Machined & Horizontal & $817 \pm 6$ & $916 \pm 13$ & $9.3 \pm 1.6$ & [186] \\ \hline Ti-6Al-4V & DED & Machined & Horizontal & $960 \pm 26$ & $1063 \pm 20$ & $10.9 \pm 1.4$ & [142] \\ \hline Ti-6Al-4V & DED & Machined & Vertical & $958 \pm 19$ & $1064 \pm 26$ & $14 \pm 1$ & [142] \\ \hline Ti-6Al-4V & DED & Machined & Vertical & $941 \pm 4$ & $1027 \pm 1$ & $8.6 \pm 0.1$ & [159] \\ \hline Ti-6Al-4V & DED & Machined & Horizontal & $1027 \pm 6$ & $1077 \pm 14$ & $2.9 \pm 0.6$ & [187] \\ \hline Ti-6Al-4V & DED & Machined & Vertical & $1031 \pm 68$ & $1106 \pm 52$ & $6.8 \pm 1.2$ & [187] \\ \hline \end{tabular} \end{center} $\sigma_{0.2}$ - yield strength; UTS - ultimate tensile strength; $\varepsilon$ - elongation. crostructures of Ti alloys fabricated by most AM techniques, columnar prior $\beta$ grains are more prevalent compared to equiaxed grains. This is ascribed to the complicated thermal history during the AM process, which can result in heat accumulation between newly deposited layers and promote the epitaxial growth of columnar grains. Furthermore, in most AM processes, the chemical components of Ti-6Al-4V between different deposited layers remain nearly identical. This similarity in composition provides essential conditions for grain epitaxial growth because adjacent deposited layers have the same crystal structure. Consequently, the epitaxial growth of metallic grains does not require the nucleation of a new phase [174,175]. As seen in Fig. 10, the columnar prior $\beta$ grains formed in Ti-6Al-4V alloys fabricated by different AM methods. In general, during the solidification process of each layer, columnar grains usually will grow along with the orientation of the temperature gradient $(G)[21,176]$, causing them to extend from the boundary towards the center of the melt pool. Therefore, the growth direction of grains can be significantly influenced by the energy input and the shape of the molten pool that is induced during the process. The formation of grain boundary $\alpha$ grains (GB- $\alpha$ ) along the prior- $\beta$ grain boundaries in Ti-6Al-4V alloys produced via different $\mathrm{AM}$ processes is a consequence of the diffusion transition from $\beta$ to $\alpha$ phases. As depicted in Fig. 11, elongated GB- $\alpha$ grains are evident along the prior- $\beta$ grain boundaries in Ti-6Al-4V alloys fabricated through L-PBF, EB-PBF, and DED techniques. Notably, the relatively higher cooling rates in DED and L-PBF result in smaller GB- $\alpha$ grain sizes in Ti alloys compared to EB-PBF-fabricated Ti-6Al$4 \mathrm{~V}$ alloys. The existence of GB- $\alpha$ grains and columnar prior- $\beta$ grains is generally considered as the main contributor to anisotropy observed in the mechanical properties of Ti alloys [182,183]. Thus, Ti alloys fabricated by AM techniques usually exhibit anisotropy in mechanical properties. Table 5 highlights the anisotropy in the mechanical behaviors of AM-built Ti-6Al-4V. Two phenomena are particularly noteworthy here. Firstly, the tensile strength and ductility present a converse trend with respect to the anisotropy property [184]. Secondly, horizontally fabricated Ti-6Al-4V typically exhibit higher strengths than vertically fabricated samples across all 3 AM techniques. Conversely, vertically manufactured parts exhibit better ductility than horizontally manufactured parts for Ti-6Al-4V alloys built by the three studied AM techniques. Anisotropy in ductility can be ascribed to the difference in loading conditions experienced by GB- $\alpha$ grains and columnar prior $\beta$ grains when subjected to stress in different directions [142]. As shown in Fig. 12(a-c), when a tensile load is applied vertically on the boundaries of prior $\beta$ grains, the short axes of the prior$\beta$ grains and GB- $\alpha$ grains undertake the loads, resulting in the separation of adjacent prior- $\beta$ grains and intergranular fracture. In contrast, as shown in Fig. 12(d), when the applied load is parallel to the boundaries of prior $\beta$ grains, the entire prior- $\beta$ grains and GB- $\alpha$ grains share the load, resulting in an increased effective slip distance between neighboring grains. Therefore, compared to horizontally fabricated samples, vertically fabricated Ti-6Al-4V presents better ductility. In addition to ductility, strength anisotropy can be ascribed to varying numbers of prior- $\beta$ grains and GB- $\alpha$ grains formed in Ti-6Al-4V samples fabricated by AM techniques with different build orientations. Compared to vertically fabricated samples, the horizontally fabricated Ti-6Al-4V parts involve more prior$\beta$ grains. This disparity arises because the column $\beta$ grains usually tend to continuously grow along the build orientation, resulting in elongation of the $\beta$ grains but not a significant change in $\beta$ grains quantity. Thus, according to the Hall-Petch relationship, horizontally fabricated Ti-6Al-4V samples have more prior- $\beta$ grain boundaries which obstruct dislocation slip and enhance the strength of samples [140,188].\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-14} Fig. 11. GB- $\alpha$ grains in Ti-6Al-4V alloys manufactured by (a, b) L-PBF (Reproduced with permission from Ref. [179]. Copyright (2021), Elsevier), (c, d) EB-PBF (Reproduced with permission from Ref. [180]. Copyright (2020), Elsevier) and (e, f) DED (Reproduced with permission from Ref. [181]. Copyright (2021), Elsevier). \subsection*{4.3. Columnar to equiaxed transition (CET)} Apart from columnar grains, researchers have also attempted to realize the columnar to equiaxed transition (CET) in AM-built Ti alloys. One approach to achieve CET is by adjusting printing parameters to change the ratio of temperature gradient $(G)$ and solidification rate $(R)$, denoted as $G / R[21,189]$. A low $G / R$ is favorable for promoting CET [21]. Bontha et al. [190] have explored the influences of processing parameters on the solidification mode of\\ DED-fabricated Ti-6Al-4V alloys by conducting analytical (Rosenthal) and numerical (FEM) modeling. They have indicated that by tuning process variables, such as scan speed and laser power, a graded microstructure along the build deposition depth (height) can be formed, particularly transitioning from columnar to mixed or equiaxed morphology near the surface region with higher laser powers, while keeping the scan speed constant. In a study by Wu et al. [151], these authors found that the length of columnar grains became shorter and was substituted by very large equiaxed grains \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-15(1)} \end{center} Fig. 12. (a) Morphology of DED-fabricated Ti-6Al-4V alloys, (b) highlighted rectangular section used for schematic diagram of GB- $\alpha$ and prior- $\beta$ grains subjected to loads in different directions: (c) represent the load orientation perpendicular to the long axes of prior- $\beta$ grains and (d) represent the load orientation parallel to the long axes of prior- $\beta$ grains (Reproduced with permission from Ref. [142]. Copyright (2015), Elsevier).\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-15} Fig. 13. (a) Equiaxed grains partially form in the DED-fabricated Ti-6Al-4V alloys when adopting the laser power of $1 \mathrm{~kW}$ (Reproduced with permission from Ref. [191]. Copyright (2022), Springer Nature) and (b) equiaxed grains of the L-PBF-built Ti-6Al-4V alloy (Reproduced with permission from Ref. [113]. Copyright (2017), Elsevier). in the DED-built Ti-6Al-4V alloys near the sample's bottom location when an extremely high laser power was used, while other parameters were kept constant. Similarity, as shown in Fig. 13(a), Marshall et al. [191] also found whilst the high laser power (1 kW) was applied, a partially equiaxed microstructure can occur in the DED-fabricated Ti-6Al-4V alloys because of the reduced $G$. In addition, as seen in Fig. 13(b), Xu et al. [113] demonstrated that equiaxed grains also occurred in L-PBF-built Ti-6Al-4V alloys when adopting a small layer thickness and inter-layer time. This approach promoted heat accumulation, resulting in reduced $G$ and facilitated CET. Although the number of cases is limited, these previous cases have shown that appropriately increasing energy density is an effective strategy to decrease $G$ to achieve CET in Ti alloys manufactured by PBF and DED. In addition to adjusting $G$ and $R$, providing extra favorable conditions for heterogeneous nucleation is another effective way to realize CET. Zhang et al. [192] have reported that utilizing a combination of high powder feed rate and lower energy density promotes the formation of equiaxed grains in DED-fabricated Ti-6Al-2Sn-2Zr3Mo-1.5Cr-2Nb alloy, as shown in Fig. 14(b). The increased amount of unmelted powder in the molten pool created by the high powder feed rate favors heterogeneous nucleation and promotes CET. In addition, the use of lower laser energy density can also decrease the rate of powder melting, resulting in an increased nucleation \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-15(2)} \end{center} Fig. 14. Equiaxed grains formed in the (a) bottom part of EB-PBF-fabricated Ti6Al-4V alloy (Reproduced with permission from Ref. [194]. Copyright (2013), Elsevier), (b) DED-fabricated Ti-6Al-2Sn-2Zr-3Mo-1.5Cr-2Nb (Reproduced with permission from Ref. [192]. Copyright (2016), Elsevier) and (c) Ti-6.5Al-3.5Mo-1.5Zr-0.3Si alloys (Reproduced with permission from Ref. [173]. Copyright (2015), Elsevier). rate. Similarly, as shown in Fig. 14(c), Wang et al. [173] indicated that equiaxed grains can nearly fully form in DED-fabricated Ti6.5Al-3.5Mo-1.5Zr-0.3Si alloys under conditions of relatively higher powder feed rate and lower energy density. A similar phenomenon was reported in DED-fabricated Ti-25V-15Cr-2Al-0.2C alloys in the research of Ref. [193]. Moreover, Antonysamy et al. [194] noted equiaxed grains forming in the bottom part of EB-PBF-fabricated Ti-6Al-4V alloys (Fig. 14(a)), which could be ascribed to heterogeneous nucleation induced by partly unmelted powder. During the AM production process, the build plate serves as a thermal conductor, often referred to as a "heat sink" [28]. As the build height increases, the efficiency of thermal conductivity from the newly deposited layer to the build plate decreases. Consequently, powder in the upper layers retains more heat energy, enabling better melting. However, this thermal limitation can lead to incomplete powder melting, especially in the lower layers, due to rapid heat dissipation from the build plate. This phenomenon somehow can impact overall print quality and part integrity. Although CET can be achieved by increasing powder feed rate and decreasing energy density, these changed printing parameters may result in undesired side effects, such as increased porosity, cracks, delamination, and others. Therefore, it is ideal to achieve CET while still utilizing optimized printing parameters. In recent research by Ref. [178], high-intensity ultrasound was verified as an effective method for achieving the CET in Ti-6Al-4V alloy during the fabrication process of DED (Fig. 15(a)). This is due to the induction of acoustic cavitation in the liquid by ultrasonic irradiation, which agitates the metallic liquid and promotes nucleation during the solidification process [195,196]. In addition, Todaro et al. [178] reported that ultrasonic irradiation does not affect the melted rate, as both ultrasound-treated and untreated DEDfabricated Ti-6Al-4V alloys exhibited similar porosity levels. Some researchers have also attempted to achieve CET from alloy constitution modification. It has been observed that larger values of the growth restriction factor $(Q)$ facilitate more nucleation and influence the development of equiaxed grains. According to Kozlov and\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-16} Fig. 15. (a) OM image of the DED-built Ti-6Al-4V alloy with ultrasound (scale bars, $1 \mathrm{~mm}$ ) (Reproduced with permission from Ref. [178]. Copyright (2020), Nature Portfolio); (b) OM image of the EB-PBF-fabricated Ti6Al4V-7Cu alloy (Reproduced with permission from Ref. [201]. Copyright (2022), Elsevier), and (c) OM micrograph of the DEDfabricated Ti-8.5Cu alloy (Reproduced with permission from Ref. [199]. Copyright (2019), Springer Nature). Schmid-Fetzer [197], a higher $Q$ not only effectively refines grains, but also affects the formation of equiaxed grains. Based on this principle, some research groups have investigated the addition of alloying elements such as beryllium (Be), silicon (Si) or boron (B) to control the grain microstructure of AM-fabricated Ti alloys [198]. However, the addition of those alloying elements usually results in a decrease in the size of the columnar grains or only achieves partial CET in AM-fabricated Ti alloys. Recent research by Zhang et al. [199] has overcome this challenge. As shown in Fig. 15(c), they found that by using the same printing parameters, the addition of copper $(\mathrm{Cu})$ enables DED-fabricated Ti-8.5Cu to fully consist of equiaxed grains, while DED-fabricated Ti-6Al-4V alloys still comprised of columnar grains. This can be ascribed to the significantly higher $Q(62 \mathrm{~K})$ of the Ti-8.5Cu alloy compared to the $Q$ value of only $8 \mathrm{~K}$ in Ti-6Al-4V alloys. Similarly, Choi et al. [200] employed mixed powders of Ti-6Al-4V and Co-Cr-Mo alloys as feedstock, where they reported that equiaxed grains can fully form in DED-fabricated samples containing $5 \mathrm{wt} \%$ and $10 \mathrm{wt} \%$ Co due to the high $Q$ value of Co. Furthermore, as seen in Fig. 15(b), Mosallanejad et al. [201] also proved that an equiaxed microstructure can be obtained in the EB-PBF-fabricated Ti-6Al-4V-7Cu alloy by using a mixed powder of Ti-6Al-4V and Cu. Given the existing challenges in achieving CET in Ti alloys, there are only a limited number of studies focused on the me- chanical properties of AM-fabricated Ti alloys containing equiaxed grains. As shown in Fig. 16(a), Xu et al. [113] reported that L-PBFfabricated Ti-6Al-4V alloys with equiaxed grain structures exhibit enhanced strength compared to the mill-annealed Ti-6Al-4V alloy. This improvement is attributed to the finer grain size resulting from the rapid cooling rate during L-PBF. In contrast, when comparing these equiaxed-grain Ti-6Al-4V alloys to typical L-PBFfabricated ones, a slight reduction in strength is observed, but with an increase in ductility. This variance is attributed to the higher energy density utilized in the process, promoting the formation of lamellar $(\alpha+\beta)$ phases within the prior $\beta$ equiaxed grains instead of $\alpha^{\prime}$ martensite, which is less favorable for ductility. Moreover, the unique morphology and the absence of continuous GB- $\alpha$ grains in equiaxed grain structures significantly reduce stress concentrations, effectively enhancing material crack resistance and ductility. In another study by Todaro et al. [178], as shown in Fig. 16(b), both YS and UTS can be improved by $\sim 12 \%$ for DED-fabricated Ti-6Al-4V alloys with ultrasound treatment; however, the ductility of DED-built Ti-6Al-4V alloys remains similar, regardless of ultrasound treatment. Compared with columnar grains, the size of equiaxed grains is smaller in DED-built Ti-6Al-4V alloys with ultrasound treatment. As such, achieving CET through various methods in the AM of Ti alloys can provide more flexibility to adjust mechanical properties.\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-17(1)} Fig. 16. (a) Engineering stress-strain curves of (a) mill-annealed and L-PBF-built Ti-6Al-4V alloys (Reproduced with permission from Ref. [113]. Copyright (2017), Elsevier) and (b) DED-built Ti-6Al-4V alloys with and without ultrasound treatment (Reproduced with permission from Ref. [178]. Copyright (2020), Nature Portfolio). \section*{5. Defects analysis and mitigation} \subsection*{5.1. Porosity} Porosity is a normal defect in metallic materials fabricated through AM techniques, where its complete elimination remains a challenge [202,203]. Porosity represents the void volume in a material, quantified as the pore volume to the total volume ratio of the material. During AM, the formation of pores can be induced by a series of complex phenomena, which are still subject to ongoing debate regarding their origins [204]. In general, the incomplete closure of a keyhole, entrapped gas and lack of fusion (LOF) are deemed the primary factors for the formation of pores in AMfabricated Ti alloys [66]. The keyhole formation is commonly induced by excessively high-power density during the AM production process. These keyholes, which can be unstable and prone to collapse, trap internal vapor and surrounding gas, leading to the creation of pores in the melt pool [21]. Keyhole-induced pores are usually irregular and depend on the size and shape of keyholes, as shown in Fig. 17(a) and (b). Entrapped gas also can contribute to pores during the AM process. The entrapped gas can originate from two main sources. Firstly, during the powder production stage, the atomization process could cause gas to remain within the powder particles [205]. Subsequently, the gas can be released and retained within the material during the AM process. Tammas-Williams et al. [206] have reported that some small spherical pores observed in EB-PBF-fabricated Ti-6Al-4V originated from the argon gas carried by the powder feedstock during atomization. Secondly, shielded gas used during the AM process can also become entrapped in the melt pool, resulting in pore formation. Generally, pores induced by entrapped gas tend to be nearly spherical and tiny, as shown in Fig. 17(c). Improper printing parameters, such as the excessively high scan speed, low power from heat source or large hatching space, can result in insufficient energy density, which prevents complete melting and results in LOF. When a new layer is deposited over an area with LOF, pores form between the boundaries of the two adjacent layers [207]. LOF-induced pores, which are usually bigger and irregular compared with pores induced by entrapped gas [208], are shown in Fig. 17(d). In general, the amount, shapes and sizes of pores vary significantly in Ti alloys fabricated by different AM techniques. As shown in Table 6, L-PBF-fabricated Ti alloys exhibit more pores and the highest porosity, while Ti alloys fabricated by EB-PBF and DED tend to have fewer pores and relatively lower porosity. According to Liu\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-17} Fig. 17. Different types of pores induced by (a) (Reproduced with permission from Ref. [62]. Copyright (2016), Elsevier) and (b) keyhole (Reproduced with permission from Ref. [209]. Copyright (2020), Elsevier), (c) entrapped gas (Reproduced with permission from Ref. [210]. Copyright (2020), Elsevier), and (d) LOF in Ti alloys manufactured by AM techniques (Reproduced with permission from Ref. [211]. Copyright (2015), Elsevier). et al. [62], L-PBF-fabricated Ti-2448 alloys contain ten times more defects (pores) compared to those fabricated by EB-PBF. On the one hand, the majority of pores that occur in L-PBF-fabricated Ti alloys are often bigger and irregular in shape, with some categorized as pores induced by keyhole and LOF. On the other hand, Ti alloys fabricated by EB-PBF and DED methods typically exhibit smaller, nearly spherical pores, with most pores being gas-induced. The distinctions in pore characteristics among the 3 AM techniques can be attributed to their different manufacturing processes. The typical melting process of EB-PBF and L-PBF is presented in Fig. 18. According to Liu et al. [62], L-PBF, with its smaller size of laser beam spot, effects of ray reflection and discontinuous scanning mode, can easily form deep key holes, resulting in the creation of more conical keyhole-induced pores. In contrast, EB-PBF, characterized by a larger electron beam spot and continuous scanning track, typically forms a wider melt pool, resulting in the presence of spherical pores in Ti alloys [62]. In addition, the L-PBF process is carried out in a shield gas-filled chamber, which can potentially cause the entrapment of gas in the melt pool [21]. Conversely, EB-PBF employs a vacuum building chamber, effectively Table 6 Relative porosity and defect size and shape of AM-fabricated Ti alloys. \begin{center} \begin{tabular}{|c|c|c|c|c|c|} \hline Material & Method & Porosity (\%) & Equivalent diameter $(\mu \mathrm{m})$ & Pore shapes & Refs. \\ \hline CP-Ti & L-PBF & $0.27-0.25$ & - & - & $[212]$ \\ \hline CP-Ti & L-PBF & $3.6-0.5$ & - & - & $[130]$ \\ \hline CP-Ti & L-PBF & $1.1-0.4$ & - & Irregular & $[94]$ \\ \hline CP-Ti & L-PBF & $1.9 \pm 0.51$ & - & - & $[213]$ \\ \hline CP-Ti & EB-PBF & 0.3 & - & - & $[96]$ \\ \hline CP-Ti & DED & $2.5-0.5$ & - & Irregular & $[94]$ \\ \hline Ti-6Al-4V & L-PBF & 0.01 & $8-63$ & Spherical; flat & $[214]$ \\ \hline Ti-6Al-4V & L-PBF & 0.29 & $3-108$ & Irregular & [99] \\ \hline Ti-6Al-4V & L-PBF & 0.5 & - & - & $[215]$ \\ \hline Ti-6Al-4V & EB-PBF & 0.005 & $3-28$ & Spherical & [99] \\ \hline Ti-6Al-4V & EB-PBF & 0.0397 & 13.3 & Spherical & $[206]$ \\ \hline Ti-6Al-4V & EB-PBF & 0.25 & - & & $[215]$ \\ \hline Ti-6Al-4V & EB-PBF & 0.072 & 81.7 & Spherical & $[216]$ \\ \hline Ti-6Al-4 V ELI & EB-PBF & 0.15 & - & Spherical; flat & $[217]$ \\ \hline Ti-6Al-4V & DED & 0.01 & $26.12-128.98$ & Spherical & $[218]$ \\ \hline Ti-6Al-4V & DED & 0.1 & $1-3$ & Spherical & $[75]$ \\ \hline Ti-6Al-4V & DED & 0.00066 & - & Spherical & $[219]$ \\ \hline Ti-6Al-4V & DED & 0.09 & - & - & $[215]$ \\ \hline Ti-48Al-2Nb-0.7Cr-0.3Si & EB-PBF & $0.17 \pm 0.24$ & 30 & Spherical & $[220]$ \\ \hline $\mathrm{Ti}-34 \mathrm{Nb}$ & L-PBF & $0.12 \pm 0.1$ & - & Spherical, irregular & $[221]$ \\ \hline Ti-35Nb & L-PBF & $\pm 0.6$ & - & \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-18} & $[38]$ \\ \hline Ti-35Nb-7Zr-5Ta & L-PBF & $1-7$ & - & - & $[123]$ \\ \hline Ti-35Nb-7Zr-5Ta & L-PBF & $1.5-0.5$ & - & Spherical, irregular & $[125]$ \\ \hline Ti-2448 & L-PBF & 0.37 & $20-140$ & Irregular; conical & $[62]$ \\ \hline Ti-2448 & EB-PBF & 0.06 & $20-120$ & Spherical & $[62]$ \\ \hline TI-6Al-2Sn-4Zr-2Mo & EB-PBF & $0.6-0.8$ & - & Spherical & $[222]$ \\ \hline \end{tabular} \end{center} \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-18(1)} \end{center} \section*{Electron beam melting process} Fig. 18. Simplified diagram of the melting process of L-PBF and EB-PBF (Reproduced with permission from Ref. [62]. Copyright (2016), Elsevier). avoiding shield gas-induced pores in Ti alloys. Further, Gaytan et al. [223] have noted that the pores in EB-PBF-fabricated Ti alloys are primarily caused by gas contained in the powder feedstock and the vaporization of low-boiling elements [223]. Moreover, due to the relatively lower energy density and the higher likelihood of LOF during the L-PBF process, elongated and flat pores are more commonly observed in L-PBF-fabricated Ti alloys, resulting in a higher pore content compared to EB-PBF-fabricated Ti alloys. In contrast to PBF systems, the DED process directly adds metallic powder through a nozzle into the active melt pool without a pre-deposited powder bed. This can result in insufficient release of gas from the bottom of the powder bed, leading to the formation of gas-induced pores. In addition, the size of the laser beam spot in DED is similar to that of EB-PBF and bigger than that of L-PBF [72], resulting in fewer keyhole-induced pores in Ti alloys fabricated by DED. Moreover, due to heat accumulation and higher energy density in the DED process, LOF occurs less frequently in DED-fabricated Ti alloys. Therefore, DED-fabricated Ti alloys generally exhibit relatively lower porosity compared to Ti alloys fabricated by PBF. \subsection*{5.2. Surface roughness} Surface roughness is also a common defect in Ti alloys manufactured by AM techniques. This is generally related to two factors: Table 7 Surface roughness of Ti-6Al-4V alloys manufactured by different AM techniques. \begin{center} \begin{tabular}{lllll} \hline Method & L-PBF & EB-PBF & DED & Refs. \\ \hline Actual surface roughness (ASR, $\mu \mathrm{m})$ & 8.485 & $28.803-45.7$ & $7.867-63.9$ & $[105,225,226]$ \\ Edge surface roughness (ESR, $\mu \mathrm{m})$ & $5-40$ & $25-131$ & $0.24-13.3$ & $[227-230]$ \\ \hline \end{tabular} \end{center} \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-19} \end{center} Fig. 19. Surface morphology of Ti-6Al-4V manufactured by (a) L-PBF and (b) EB-PBF (Reproduced with permission from Ref. [231]. Copyright (2018), MDPI). the actual surface roughness (ASR) and the edge surface roughness (ESR). The ASR presents the surface roughness of the flat layer and can be calculated using Eq. (3) [224]: \begin{equation*} R_{\mathrm{as}}=\frac{1}{N} \sum_{i=1}^{N}\left|f_{\mathrm{n}}\right| \tag{3} \end{equation*} where $R_{\mathrm{as}}$ is the ASR, $f_{n}$ is the height of a peak or the depth of a valley in one measured position, and $N$ represents the number of measured positions. The ASR can vary among Ti alloys manufactured by different AM techniques. A comparison in Table 7 reveals that the Ti-6Al-4V alloy fabricated by both L-PBF and DED methods can achieve relatively lower ASR in comparison with the counterparts manufactured by EB-PBF. Two reasons can explain this phenomenon. Firstly, the size of powder feedstock used in L-PBF is smaller than the powder adopted in EB-PBF. As reflected in Fig. 19, many larger unmelted metallic powders adhere to the surface of Ti-6Al-4V fabricated by the EB-PBF, while the size of powders is much smaller on the surface of L-PBF-fabricated Ti-6Al-4V alloy. In addition, the DED technique has a special feeding system that allows unused metallic powders to be blown away, resulting in fewer powder particles adhering to the surface of the fabricated products. Therefore, Ti alloys fabricated by L-PBF and DED usually can achieve lower ASR than those of EB-PBF-fabricated Ti alloys. In addition to ASR, the ESR also can effectively influence the roughness of products fabricated by AM techniques. The "stair step effect" is the primary factor contributing to ESR, as shown in Fig. 20(a). This effect occurs when a thicker deposited layer and more inclined build direction are present, leading to an increase in the severity of the "stair step effect". The average roughness $\left(R_{\mathrm{a}}\right)$ caused by "stair step effect" can be calculated using Eq. (4) [232]: \begin{equation*} R_{\mathrm{a}}=1000 t_{\mathrm{l}} \sin \left(\frac{90-\theta}{4}\right) \tan (90-\theta) \tag{4} \end{equation*} where $t_{1}$ is the layer thickness and $\theta$ represents the build angle. Table 7 demonstrates that both L-PBF and DED techniques yield Ti alloys with lower ESR compared to the EB-PBF technique. Fig. 20(b) shows that the thicker deposited layer $(70 \mu \mathrm{m})$ during the EB-PBF contributes to higher ESR in fabricated Ti-2448 alloys compared to deposited layers $(50 \mu \mathrm{m})$ in the L-PBF production process. This disparity in layer thickness leads to the higher ESR observed in EBPBF-fabricated Ti-2448 alloys. Furthermore, in the case of DED, although the deposited layer is also thick, the inter-layer fusion is more pronounced because of the higher energy density and heat accumulation effect during the DED process. Therefore, Ti alloys fabricated by both L-PBF and DED present higher surface accuracy than those fabricated by EB-PBF. \subsection*{5.3. Residual stress} Residual stress (RS) is stress that exists in solid material when products are under the conditions of thermal equilibrium and not subjected to external stress [233]. The RS in metals is usually originated from the large-scale temperature variation experienced during the production process. According to DebRoy et al. [28], when a metallic material suffers substantial temperature changes during production, the thermal strain of the material can exceed the limitation of its elastic strain, leading to the conversion of strain from elastic to plastic and resulting in stress accumulation. The relationship between the thermal strain of metals and temperature change can be expressed as Eq. (5) [28]: $\varepsilon_{\mathrm{T}}=\alpha\left(T-T_{0}\right)$ where $\varepsilon_{\mathrm{T}}$ is the thermal strain of metals, $\alpha$ is the coefficient of thermal expansion (CTE) of material, $T$ and $T_{0}$ are the local temperature and defined initial temperature, respectively. The CTE of Ti alloys is usually above $1 \times 10^{-5} \mathrm{~K}^{-1}$, which means even a temperature change of only a few hundred $\mathrm{K}$ can cause stress accumulation [28]. Therefore, due to the large thermal gradient $(G)$ in the melt pool, RS can more easily occur in Ti alloys manufactured by AM techniques compared to those produced with traditional methods [234]. The L-PBF method, characterized by large G and rapid cooling rates, tends to result in relatively higher levels of RS [235]. XRD patterns are commonly used to analyze the magnitude and type of RS in material. A right shift in the diffraction peak represents the compressive RS occurring in the fabricated material, while a left shift indicates tensile RS. Moreover, the larger the displacement distance (shift) of the diffraction peak, the larger the $\mathrm{RS}$ remains near the material's tested surface region. As shown in the inset of Fig. 21, compared to EB-PBF-fabricated Ti-6Al-4V alloys, the diffraction peaks of $\alpha / \alpha^{\prime}$ phase in L-PBF-fabricated Ti-6Al$4 \mathrm{~V}$ exhibit a greater right shift, which indicates the presence of higher compressive RS near the region. In addition, Sharma et al. [236] have also reported a high RS of $295 \pm 12.5 \mathrm{MPa}$ in L-PBFfabricated Ti-6Al-4V alloy. Szost et al. [237] have also indicated that Ti-6Al-4V fabricated by DED has a maximum tensile RS of $280 \mathrm{MPa}$ and a maximum compressive RS of $250 \mathrm{MPa}$. Contrastingly, compared to Ti alloys fabricated by L-PBF and DED methods, the RS is very small and can often be ignored in Ti alloys fabricated by the EB-PBF $[105,238,239]$. This is because the production process of EB-PBF typically involves prolonged exposure to high build platform temperatures $\left(\sim 650-750{ }^{\circ} \mathrm{C}\right)$ [73], which is similar to the annealing process. Therefore, the process of producing EB-PBF can effectively mitigate RS in final fabricated Ti alloys. The distribution of RS can vary in Ti alloys manufactured by AM techniques. The RS distribution of Ti-6Al-4V fabricated by L$\mathrm{PBF}$ is shown in Fig. 22, where it can be seen that high tensile RS occurs near the surface edge surrounding area, which gradually transitions into compressive RS towards the central area of surface layer. Furthermore, Ahmad et al. [240] have shown that the maximum tensile RS ranges from 805 to $837 \mathrm{MPa}$ in the area that is \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-20(3)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-20} \end{center} EBM sample \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-20(2)} \end{center} SLM sample \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-20(1)} \end{center} EBM melting process \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-20(4)} \end{center} (b) Fig. 20. (a) Schematic diagram of the "stair step effect" (Reproduced with permission from Ref. [232]. Copyright (2015), Springer Nature) and (b) surface morphologies and melting schematic diagram of Ti-2448 alloys manufactured by L-PBF and EB-PBF (Reproduced with permission from Ref. [62]. Copyright (2016), Elsevier). \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-20(5)} \end{center} Fig. 21. XRD pattern of Ti-6Al-4V alloys fabricated by L-PBF and EB-PBF (Reproduced with permission from Ref. [104]. Copyright (2021), Elsevier). near the surface edge surrounding area, while the maximum compressive RS ranges from 434 to $459 \mathrm{MPa}$ in the area close to the center of surface layer in the Ti-6Al-4V alloy manufactured by L-\\ PBF. Moreover, Yakout et al. [241] have reported tensile RS in the range of 445-528 MPa at a depth of $0.4 \mathrm{~mm}$ below the top surface, which decreases to around 20-264 MPa at a depth of $1 \mathrm{~mm}$ below the top surface. \subsection*{5.4. Inhomogeneity and chemical composition alteration} The design types and volume of constituent elements are crucial in determining the properties of the final fabricated products. Even minor changes in the chemical composition can have a substantial influence on the performance of products. However, it is currently challenging to completely avoid chemical composition changes in final products fabricated by AM techniques. Therefore, it is necessary to gain a deep understanding of the mechanism behind the chemical composition changes in Ti alloys during different AM techniques to improve the stability of the properties of final products. Interstitial element pickup and alloying elements evaporation are commonly two main sources for chemical composition changes. The interstitial elements pickup typically leads to an increase in the total element content in AM-fabricated Ti alloys. Conversely, the evaporation or loss of alloying elements results in a decrease in their concentration in the alloys. Accordingly, understanding and controlling these factors are crucial to guaran-\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-20(6)} Fig. 22. Stress distribution at the surface layer (scan plane) of Ti-6Al-4V alloys manufactured by L-PBF: (a) contour stress maps and (b) stress distribution along the diagonal (Reproduced with permission from Ref. [240]. Copyright (2018), Elsevier).\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-21} Fig. 23. (a) Relationship between the interstitial element concentration in Ti-6Al-4V alloys manufactured by L-PBF and the 0 concentration of building chamber (Reproduced with permission from Ref. [242]. Copyright (2020), Elsevier) and (b) trend of O concentration and hardness of L-PBF-built Ti-6Al-4V alloys with the increase of laser beam exposure time from increasing amount of laser passes (Reproduced with permission from Ref. [243]. Copyright (2019), Elsevier). tee the stability of desired chemical constitution and properties of AM-fabricated Ti alloys. \subsection*{5.4.1. Interstitial element pickup} During the AM process, additional impurities can be introduced into the fabricated Ti alloys from the surrounding environment. Researchers have investigated the effects of oxygen $(0)$ concentration and laser parameters on chemical composition changes in Ti alloys fabricated using the L-PBF. Dietrich et al. [242] have investigated chemical composition changes in L-PBF-fabricated Ti-6Al-4V alloys under different $O$ concentrations in the building chamber. As shown in Fig. 23(a), increasing the 0 concentration in the build chamber results in an increase in the $\mathrm{O}$ and nitrogen $(\mathrm{N})$ content in the Ti-6Al-4V alloy compared to the powder raw material. The 0 content increased by up to $\sim 41 \%$ (from $\sim 0.13 \mathrm{wt} . \%$ to $\sim 0.18 \mathrm{wt} . \%$ ), and the $\mathrm{N}$ content increased by $\sim 222 \%$ (from $\sim 0.0062 \mathrm{wt} \%$ to $\sim 0.02 \mathrm{wt} . \%$ ). Additionally, the content of extra interstitial elements can also change in AM-fabricated Ti alloys, even at a constant gas concentration in the building chamber. Velasco-Castro et al. [243] have applied L-PBF to produce Ti-6Al-4V samples with different numbers $(1,3,5,9$ beam passes) of repeated passes of the beam on each layer, whereby there was more oxygen in LPBF-built Ti-6Al-4V samples with the greater number of beams passes in each layer. This is because, as shown in Fig. 23(b), more beams passing increases the total energy input, which results in a longer melting time during the production process. As a result, the molten pool will repeatedly be exposed to the building chamber, promoting the absorption of interstitial elements. Similarly, $\mathrm{Na}$ et al. [244] have reported $\mathrm{O}$ and $\mathrm{N}$ elements to increase in L-PBFbuilt CP-Ti alloys with increased laser power. However, Wang et al. [39] have stated that the L-PBF-built Ti-35Nb alloy ( 0.12 wt.\% 0) using prealloyed powder exhibited a very low $\mathrm{O}$ pick-up of approximately $0.02 \mathrm{wt} . \%$ compared to the initial $0.10 \mathrm{wt} \% \mathrm{O}$ in the prealloyed powder under optimized processing parameters. Conversely, the Ti-35Nb produced using a powder mixture showed a higher $\mathrm{O}$ content of $0.30 \mathrm{wt} . \%$ as a result of incorporating elemen- \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-22} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-22(2)} \end{center} (b) Fig. 24. (a) Ti-6Al-4V alloy manufactured by DED (the color of the O-contaminated area is darker) (Reproduced with permission from Ref. [142]. Copyright (2015), Elsevier) and (b) concentration of $\mathrm{N}$ and $\mathrm{O}$ elements in the powder feedstock and Ti-6Al-4V samples built by DED with different processing conditions: ArS (Argon Shielding, general condition), SE-HB (Sealed Environment, Hot Base), SE-CB (Sealed Environment, Cold Base) (Reproduced with permission from Ref. [245]. Copyright (2020), MDPI). tal Ti powder (with 0.10 wt.\% O) and Nb powder (with 0.32 wt.\% 0). Additionally, both built prealloyed and powder mixture Ti-35Nb showed very low $\mathrm{N}$ content, with values of $0.024 \mathrm{wt} \% \mathrm{~N}$ and 0.030 wt.\% N, respectively. With regard to Ti alloys fabricated using the DED process, studies have also reported the occurrence of additional interstitial elements, particularly O. Carroll et al. [142] have found that the 0 content of DED-fabricated Ti-6Al-4V alloy was about 0.2046 wt.\%, approximately $0.0316 \mathrm{wt}$.\% higher than the 0 content of the initial powder. In addition, due to water-cooling system leaking, they reported that in areas affected by leaked water vapor, the $O$ content was measured to be 0.2170 wt.\%. As shown in Fig. 24(a), the intrusion of $\mathrm{O}$ caused oxidation and darkening of the lower parts of the fabricated products. Moreover, Carrozza et al. [245] have also found both $\mathrm{O}$ and $\mathrm{N}$ contents to increase in DED-fabricated Ti-6Al$4 \mathrm{~V}$ alloys, with the $\mathrm{O}$ content showing a greater increase compared to the $\mathrm{N}$ content, as indicated in Fig. 24(b). Although L-PBF and DED processes typically take place in a building chamber with an inert gas atmosphere provided by feeding systems, the presence of mixed interstitial elements in fabricated Ti alloys can still be observed. Firstly, it is challenging for industrial inert gas to completely eliminate $\mathrm{O}$ during its production process, whereby a portion of the $\mathrm{O}$ can enter the chamber along with the inert gas via the gas supply systems. Additionally, the sensitivity of $\mathrm{O}$ sensors varies, whereby the detection limit for $\mathrm{O}$ content is typically around 1000 ppm [243]. Moreover, apart from 0 , the inert gas itself is also a source of interstitial elements. According to Kornilov et al. [246], both $\mathrm{O}$ and $\mathrm{N}$ atoms have small enough sizes (van der Waals radii of 152 and $155 \mathrm{pm}$, severally) to fit into the voids of the Ti crystal structure. Therefore, it is challenging to fully avoid the additional interstitial elements occurring in Ti alloys fabricated by the L-PBF and DED methods. Whilst it is widely known that the production process of the EB-PBF system takes place in a vacuum chamber, $\mathrm{O}$ intrusion can still occur in $\mathrm{Ti}$ alloys fabricated by the EB-PBF method. Formanoir et al. [163] have reported $\mathrm{O}$ contamination in EB-PBF-fabricated Ti-6Al-4V, noting that some water moisture from the air can freeze on the walls of the building chamber during the evacuation of air to achieve a vacuum environment. \subsection*{5.4.2. Evaporation of alloying elements} During the AM process, the high temperature of melting and deposition processes can result in the evaporation of certain alloying elements, leading to the alloy composition deviation between the feedstock and produced parts [247,248]. This phenomenon occurs because different alloying elements have varying melting \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-22(1)} \end{center} Fig. 25. Relationship between the saturated vapor pressure and temperature for different elements (Reproduced with permission from Ref. [252]. Copyright (2020), MDPI). points, with some elements being more volatile and susceptible to selective vaporization during AM processes [249]. In general, metallic elements with lower melting points, such as aluminium (Al), are more prone to evaporation and loss during the hightemperature melting process. Keaveney et al. [250] analyzed the condensate from the L-PBF building chamber after the production of Ti-6Al-4V alloy and indicated a higher concentration of Al compared to vanadium $(\mathrm{V})$ in the collected condensate. In addition, the amount of the alloying elements loss varies among different AM techniques. In general, compared to L-PBF, EB-PBF-fabricated Ti alloys experience a more significant loss of alloying elements. Brice et al. [251] have reported a decrease of $\sim 0.9 \mathrm{wt} . \% \mathrm{Al}$ element in Ti-6Al-4V alloy manufactured by the EB-PBF method, While Gaytan et al. [223] have shown a $\sim 0.6-1.0$ wt.\% reduction of the $\mathrm{Al}$ element in EB-PBF-fabricated Ti-Al-4V alloys. In contrast, the reduction of the Al element was $\sim 0.55$ wt.\% in L-PBF-built Ti-6Al-4V [99]. As illustrated in Fig. 25, the evaporation of alloying elements is affected by temperature, with higher temperatures leading to increased loss. A higher saturated vapor pressure can result in a larger loss of alloying element [252]. Moreover, the evaporation of alloying elements usually occurs on the molten pool surface during the AM process, whereby the ratio of the surface area to vol- \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-23} \end{center} Fig. 26. (a) EBSD maps for DED-fabricated Ti6Al4V-Mo alloys with different Ti6Al4V and Mo mixed ratio (wt.\%) at four interfaces: (b) measurements taken pre and post $\beta$-phase reconstruction at the $100 \% \mathrm{Ti6Al4V} / 75 \% \mathrm{Ti6Al4V}-25 \%$ Mo interface, while (c-e) $\beta$-phase with Mo increasing at three interfaces (Reproduced with permission from Ref. [257]. Copyright (2017), Elsevier); (f) and (g) BSEM graph and EDS elemental maps for Ti-35Nb fabricated via L-PBF adopting mixed and prealloyed powder (Reproduced with permission from Ref. [39]. Copyright (2022), Elsevier). ume can greatly influence the level of alloying loss [28,253]. According to [62], the molten pool in EB-PBF is wider and shallower (width: $\sim 280 \pm 23 \mu \mathrm{m}$; depth: $\sim 152 \pm 15 \mu \mathrm{m}$ ) than the molten pool in L-PBF (width: $\sim 146 \pm 17 \mu \mathrm{m}$; depth: $\sim 172 \pm 21 \mu \mathrm{m}$ ). This indicates that the surface area-to-volume ratio of the molten pool is smaller during the production process of L-PBF. According to Klassen et al. [247], the vacuum environment of the EB-PBF building chamber can further promote the evaporation of alloying elements. \subsection*{5.4.3. Inhomogeneity challenges} Inhomogeneity challenges in the microstructure of AM-built Ti alloys have also been investigated. For instance, Wang et al. [39] have contrasted the melting behavior of the Ti-35Nb alloy manufactured by L-PBF, adopting a powder mixture and prealloyed powder. The backscattered SEM (BSEM) image and energy dispersive X-ray spectroscopy (EDS) mapping in Fig. 26(f) show that the Ti-35Nb sample produced from a powder mixture exhibited a heterogeneous microstructure with undissolved $\mathrm{Nb}$ particles of vary-\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-24} Fig. 27. Influence of surface roughness (a) and porosity (b) on the fatigue properties of Ti-6Al-4V samples manufactured by L-PBF and EB-PBF (Reproduced with permission from Ref. [238]. Copyright (2018), Elsevier). ing sizes. In contrast, the Ti-35Nb specimen from prealloyed powder (Fig. 26(g)) displayed a uniform and homogeneous microstructure with complete melting of Nb. Previous research has also highlighted the challenge of inhomogeneity caused by unmelted $\mathrm{Nb}$ particles in Ti alloys produced via L-PBF, resulting in both microstructural and chemical inhomogeneity [37,38,254,255]. Similarly, addressing the issue of inhomogeneity through parameter optimization in EB-PBF poses a significant challenge. The utilization of EB-PBF with blended elemental powders in the production of a Ti-10at.\%Nb alloy has revealed the formation of a heterogeneous microstructure [256]. Additionally, DED offers more advantages in producing functionally graded Ti alloys, including the ability to make in-situ adjustments to powder or mixture. However, it also poses challenges related to inhomogeneity, both in terms of chemical composition and microstructure. When examining DEDfabricated Ti-6Al-4V alloys with the addition of varying percentages of Mo content, two significant sources of inhomogeneity become apparent. Firstly, chemical inhomogeneity exists within each functionally graded layer due to insufficient energy to fully melt the high melting point of Mo (Fig. 26(a)). Secondly, microstructural inhomogeneity becomes more pronounced, particularly near the interfaces between gradient transitions along the build direction. This is evident in Fig. 26(b-e), where varying Mo contents result in different grain morphologies. For example, $100 \mathrm{wt} . \%$ Ti-6Al-4V forms coarse columnar $\beta$-grains, while the addition of $25 \mathrm{wt} \%$ Mo reduces the size of columnar $\beta$-grains without a specific crystallographic orientation. Further increasing the Mo content to $50 \mathrm{wt} . \%$ leads to the formation of fine equiaxed $\beta$-grains, and a Mo content of 75 wt.\% further decreases the equiaxed grain size to approximately 5-60 $\mu \mathrm{m}$. Thus, the challenge remains to optimize the melting process for powder mixtures, taking advantage of their cost-effectiveness and fast production while achieving better melting and homogeneity in the microstructure. Thus, continuous and future efforts will be necessary to enhance process control and mitigate challenges associated with interstitial element pickup, alloying element evaporation and inhomogeneity. \subsection*{5.5. Mechanical properties affected by defects} Understanding the impact of both macro and micro defects on the mechanical performance of AM-fabricated Ti alloys is crucial for optimizing their performance. As noted above, in AM-fabricated Ti alloys various defects can arise, including macro defects such as surface roughness, porosity and residual stress (RS), as well as micro defects resulting from changes in chemical composition and inhomogeneity issues. These defects, regardless of their scale, have the potential to impact the material's structural integrity and overall mechanical behavior.\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-25} Fig. 28. Relationship between fatigue crack growth ( $\mathrm{d} a / \mathrm{d} n$ ) and stress intensity factor range (dK) in Ti-6Al-4V fabricated by (a) L-PBF with different conditions and (b) EB-PBF (Reproduced with permission from Ref. [262]. Copyright (2019), Elsevier).\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-25(1)} Fig. 29. (a) Trend of hardness with increasing interstitial elements ( $O$ and N) in CP-Ti fabricated by L-PBF (Reproduced with permission from Ref. [244]. Copyright (2018), Elsevier) and (b) relationship between mechanical properties and the quantity of 0 element in L-PBF built Ti-6Al-4V (Reproduced with permission from Ref. [242]. Copyright (2020), Elsevier). \subsection*{5.5.1. Mechanical properties related to macro defects} It has been reported that macro defects can greatly affect the fatigue property of AM-fabricated Ti alloys [258]. Moreover, variations in production processes among different AM techniques can result in distinct fatigue properties of Ti alloys. Table 8 provides the fatigue properties of Ti alloys manufactured by different AM techniques, indicating that Ti-6Al-4V alloys manufactured by L-PBF and DED methods exhibit higher fatigue strength compared to EBPBF-fabricated samples. In contrast, the EB-PBF-built Ti-6Al-4V alloy presents higher fatigue toughness than those of Ti-6Al-4V alloys manufactured by L-PBF and DED. The fatigue strength of Ti alloys manufactured by different AM techniques can be attributed to two key aspects related to crack initiation behavior. Firstly, Ti-6Al-4V alloys manufactured by L-PBF Table 8 Fatigue properties of Ti-6Al-4V alloys fabricated by different AM techniques. \begin{center} \begin{tabular}{llllll} \hline Producing methods & Condition & $R$ & $\Delta \sigma_{\mathrm{w}}$ & $\Delta K_{\text {th }}$ & Refs. \\ \hline L-PBF & As-built & 0.1 & 550 & 1.7 & $[101,259]$ \\ L-PBF & As-built & 0.1 & $220 \pm 24$ & - & $[231]$ \\ EB-PBF & As-built & 0.1 & $200-250$ & 3.8 & $[162,239]$ \\ EB-PBF & As-built & 0.1 & $115 \pm 13$ & - & $[231]$ \\ DED & As-built, machined & 0.1 & $482 *$ & $2.8-3.5$ & $[105,112]$ \\ \hline \end{tabular} \end{center} $R$ - fatigue stress ratio; $\Delta \sigma_{\mathrm{w}}$ - threshold stress; $\Delta K_{\mathrm{th}}$ - crack propagation threshold; “*"- at the condition of $1 \times 10^{6}$ cycles. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-26(3)} \end{center} (c) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-26} \end{center} \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-26(1)} \end{center} (d) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_0b318dfc6a83a894290cg-26(2)} \end{center} Fig. 30. (a) Schematic diagram of interstitial $O$ elements inside the octahedral and tetrahedral lattices of Ti-35Ta alloy with a bcc lattice structure and (b) the change trend of lattice constant with the $O$ increase (Reproduced with permission from Ref. [272]. Copyright (2021), Elsevier); (c) HAADF-STEM graph of a dislocation impeded by the interstitial $\mathrm{O}$ array (top-left inset presents the geographic phase of the strain for the selected area; the $\mathrm{O}$ interstitial array is showed by iDPC-STEM image at the bottom-right inset) and (d) enlarged HAADF-STEM image illustrates the dislocation movement is impeded by a large amount of interstitial O element (Reproduced with permission from Ref. [265]. Copyright (2023), Nature Portfolio). and DED methods contain the presence of martensite $\alpha$ ' phases and finer $\alpha / \alpha^{\prime}$ grains. During fatigue tests, cyclic load leads to the formation of plastic slip localization on the surface of Ti-6Al-4V samples, which serves as the starting point for fatigue crack initiation. The presence of martensite $\alpha^{\prime}$ phases and finer $\alpha / \alpha^{\prime}$ grains can effectively hinder the dislocation slip and delay the formation of the plastic slip localization on the sample surface, resulting in improved fatigue strength of Ti-6Al-4V alloy fabricated by L-PBF and DED methods [101,105]. In addition, surface roughness and porosity are also important factors for the fatigue strength. These defects create local stress concentrations that serve as sites for fatigue crack initiation sites during the fatigue test. Chastand et al. [238] have explored the influences of surface roughness and porosity on fatigue strength. Their findings, as shown in Fig. 27(a) and (b), suggest that a better surface finish (after polishing) can increase fatigue strength by $\sim 100 \%$ for $10^{7}$ cycles compared to asbuilt ones. Additionally, when internal porosity is further reduced through hot isostatic pressing (HIP) and followed by polishing, a significant $\sim 80 \%$ increase in fatigue strength is observed compared to polished-only counterparts. This highlights the substantial impact of both surface roughness and internal porosity on the fatigue strength of Ti alloys produced through AM. Therefore, achieving a high-quality surface finish and minimizing internal porosity are essential through process optimization or post-processing techniques to boost fatigue strength and ensure the reliable performance of AM-fabricated Ti alloys. Contrasted to fatigue strength, fatigue toughness is almost not influenced at all by surface roughness. This can be attributed to the close relationship between fatigue toughness and fatigue crack growth (FCG) behavior, which is significantly affected by RS. According to Leuders et al. [259], RS is the main factor for the distinctions of FCG in L-PBF-fabricated Ti alloys with and without heat treatment. This is because the effective stress intensity factor ratio can be raised by tensile residual stress in the materials to increase\\ \includegraphics[max width=\textwidth, center]{2024_04_13_0b318dfc6a83a894290cg-27} Fig. 31. SEM images of the tensile fracture surfaces at the cross section for L-PBF fabricated Ti-35Nb: (a, b) specimen using prealloyed powder and (c, d) specimen using an elemental powder mixture (Reproduced with permission from Ref. [39]. Copyright (2022), Elsevier); (e) backscattered SEM for the surface deformation during a compressive test at a strain of $15 \%$ using an elemental powder mixture (Reproduced with permission from Ref. [37]. Copyright (2019), Elsevier), and (f) nanoindentation load-displacement curves for inhomogeneous regions (Reproduced with permission from Ref. [38]. Copyright (2021), Elsevier). the FCG rate [260]. As reflected in Fig. 28(a), the machined L-PBFfabricated Ti-6Al-4V alloys present similar FCG rates to the as-built counterparts. However, as seen in Fig. 28(b), after the heat treatment, the crack growth resistance of L-PBF-fabricated Ti-6Al-4V alloys has been effectively improved due to the decreased level of RS. In addition, Cain et al. [261] also have indicated that FCG rate can be effectively decreased in the L-PBF-fabricated Ti-6Al-4V alloy after low-temperature stress relief and annealing heat treatments. Similarly, Leuders et al. [259] have reported that the crack propagation threshold of Ti-6Al-4V alloys manufactured by L-PBF can be effectively improved following stress relief through heat treatment. This improvement can be ascribed to the reduction of RS in the heat-treated part, leading to decreased stress accumulation and improved crack growth resistance. Additionally, because of the special producing process, Ti alloys fabricated by the EB-PBF method obtain relatively smaller RS. Therefore, EB-PBF-fabricated Ti alloys exhibit relatively higher fatigue toughness than the Ti alloys manufactured by L-PBF and DED methods. This can be confirmed in Fig. 28(b), where compared with the L-PBF-fabricated Ti-6Al-4V alloys, the EB-PBF-fabricated Ti-6Al-4V alloys present a lower FCG rate. \subsection*{5.5.2. Mechanical properties related to micro defects} The influence of chemical composition changes on the mechanical performances of produced parts indicates the importance of ad- dressing this aspect. As reflected in Fig. 23(b) previously, VelascoCastro et al. [243] have reported an improvement in hardness with increasing 0 content in the L-PBF-built Ti-6Al-4V alloy. Moreover, as shown in Fig. 29(a), Na et al. [244] also found that the hardness can be enhanced with the increasing concentration of $\mathrm{O}$ and $\mathrm{N}$ in the CP-Ti alloy fabricated by L-PBF. Apart from hardness, changes in chemical composition also have an influence on properties related to overall plastic deformation in AM-built Ti alloys. In order to exclude the influence of RS and porosity, Dietrich et al. [242] have investigated the mechanical properties of L-PBF-built Ti-6Al-4V alloys after heat treatments (stress relief) and HIP. Fig. 29(b) shows that an increase in $\mathrm{O}$ concentration resulted in the improvement of both YS and UTS of Ti-6Al-4V samples, while the elongation decreased. Furthermore, Carroll et al. [142] have found that DED-built Ti-6Al-4V samples with larger $O$ content exhibited both higher YS and UTS, but their ductility was $1.6 \%$ lower than the samples compared to samples with less $\mathrm{O}$ content. These results agree well with a study by Wang et al. [39], which revealed that the Ti-35Nb alloy fabricated via L-PBF with an O content of $0.30 \mathrm{wt} . \%$ exhibited a higher tensile YS of $636 \pm 80 \mathrm{MPa}$. However, it also exhibited a significantly lower ductility compared to the counterpart with an O of 0.12 wt.\%, which had a YS of $485 \pm 28 \mathrm{MPa}$. The strengthening mechanism of interstitial elements in Ti alloys is similar to the solid solution strengthening. According to Kornilov et al. [246], $\mathrm{O}$ and $\mathrm{N}$ elements can easily occupy voids\\ within the Ti crystal structure and form solutes in the alloys. Fig. 30(a) shows the interstitial 0 element occupying octahedral and tetrahedral holes in Ti-35Ta alloy with a bcc lattice structure, where the increasing $\mathrm{O}$ content results in an increased lattice constant in the Ti-35Ta (Fig. 30(b)). The presence of interstitial solute atoms leads to lattice distortion, which impedes dislocations slip and low temperature twinning $[263,264]$. This phenomenon has been observed in research conducted by Ref. [265], where integrated differential phase contrast (iDPC)-scanning transmission electron microscopy (STEM) and high-angle annular darkfield (HAADF)-STEM technologies were used. Fig. 30(c) and (d) show that the dislocation movement was impeded in the area containing a large amount of the 0 interstitial array. Apart from the dislocation movement, according to Zaefferer [266], deformation twinning also can be totally inhibited in CP-Ti alloy with 2000 ppm of $O$ content. Therefore, both the hardness and strength can be improved by the increased interstitial elements in Ti alloys. However, this effect can come at the expense of reduced ductility. During the deformation, Ti alloys undergo a combination of twinning and slip mechanisms, which significantly influence the ductility of the material [267-271]. Therefore, with the increasing 0 content, AM-fabricated Ti alloys tend to present the relatively lower ductility. Fig. 31 provides valuable insights into the inhomogeneity issues affecting the mechanical properties of Ti-35Nb alloy manufactured by L-PBF, thus allowing for an investigation into the impact of inhomogeneity on the mechanical properties of Ti alloys. Specifically, Fig. 31(a) shows the tensile fracture surfaces of L-PBFfabricated Ti-35Nb alloy with a homogeneous chemical distribution. The presence of elongated necking (Fig. 31(a)) and homogeneous ductile fine dimples (Fig. 31(b)) confirms higher tensile ductility $(23.5 \% \pm 2.2 \%)$. As described earlier, the inhomogeneity resulting from undissolved $\mathrm{Nb}$ particles in Ti-Nb alloys has a substantial impact on the mechanical properties. Further examination, as depicted in Fig. 31(c) and (d), reveals the presence of microcracks and unmelted powders. The fracture surface is dominated by transgranular fracture and smooth cleavage facets, leading to a low ductility of $2.2 \% \pm 1.4 \%$. In Fig. 31(e), the surface compressive deformation provides additional evidence of an inhomogeneous $\beta$ phase microstructure. The weaker interface between the $\beta$ matrix and undissolved $\mathrm{Nb}$ particles hinders shear band propagation near the interfaces. The nanoindentation load-displacement curves presented in Fig. 31(f) exhibit distinct response behaviors within the inhomogeneous microstructure of Ti-Nb $\beta$ phase, undissolved $\mathrm{Nb}$ particle, and their interface. The Ti-Nb $\beta$ phase displays the smallest penetration depth (displacement), indicating higher hardness and strength, whereas $\mathrm{Nb}$ particle exhibits the lowest hardness. This further confirms the weak bonding of the interfaces, which contributes to crack initiation and premature tensile and compressive failure. Improving homogeneity has the potential to significantly enhance ductility. These findings in Fig. 31 highlight the detrimental effects of inhomogeneity and emphasize the importance of addressing this issue in order to optimize the mechanical performance of Ti alloys. \section*{6. Conclusions} Additive manufacturing (AM) provides a complementary, rather than a replacement, approach to traditional manufacturing processes due to its enhanced flexibility in fabricating shape-complex parts and solving machining challenges, resulting in reduced lead times for custom designs. This article has systematically examined phase transformation, grain size and morphology, as well as defects, and discusses their impacts on the mechanical properties of Ti alloys manufactured by three commonly used powder-type AM techniques: laser powder bed fusion (L-PBF), electron beam powder bed fusion (EB-PBF) and directed energy deposition (DED). The differences in production processes lead to distinct microstructures and defects in Ti alloys fabricated by these AM techniques. The formation of non-equilibrium phases $\left(\alpha^{\prime}, \alpha^{\prime \prime}, \omega\right)$ usually can impede the dislocation movement, which can delay the deformation of Ti alloys. Therefore, the L-PBF-fabricated Ti alloys generally exhibit relatively higher strength and hardness compared with counterparts fabricated by DED and EB-PBF methods. However, nonequilibrium phases are unfavorable for the elastic modulus and ductility of Ti alloys. Grain size and morphology also impact the mechanical properties of Ti alloys. Smaller grain size (more grain boundaries) can prevent further propagation of dislocations and decrease stress concentration, resulting in higher strength and hardness in Ti alloys fabricated by L-PBF and DED methods compared to EB-PBFfabricated counterparts. Additionally, the columnar to equiaxed transition (CET) can be potentially achieved in AM-built Ti alloys by optimizing and adjusting printing parameters to modify the ratio of $G / R$ (thermal gradient over solidification rate) or by promoting heterogeneous nucleation via the addition of refining elements. Compared to the columnar grains, the AM-built Ti alloys consisting of equiaxed grain can present both higher strength and ductility because of decreased stress concentration and grain size. Despite the numerous benefits of AM, it remains challenging to produce Ti alloys without defects at the current stage. AMfabricated $\mathrm{Ti}$ alloys typically contain macro defects (e.g., porosity, surface roughness, residual stress) and micro defects (e.g., chemical composition changes and inhomogeneity issues). These macro defects, such as porosity and surface roughness, can obviously impact the fatigue property and overall performance of AM-fabricated Ti alloys by inducing fatigue crack initiation, while residual stress (RS) promotes crack growth and significantly influences fatigue toughness. In addition to defects, the introduction of oxygen $(\mathrm{O})$ and nitrogen $(\mathrm{N})$ atoms during the AM production process can alter the chemical composition of Ti alloys and result in lattice distortion, which impedes the dislocations slip and twinning, ultimately leading to improved strength but decreased ductility. Moreover, chemical inhomogeneity in Ti alloys can have a negative impact on their mechanical properties, as undissolved elements can induce microstructural inhomogeneity, alter the phase composition, and weaken bonding between the matrix and undissolved particles, ultimately leading to reduced ductility and strength. This article provides a comprehensive understanding of the mechanical properties of $\mathrm{Ti}$ alloys manufactured using various $\mathrm{AM}$ techniques, covering both macro and micro perspectives. Although the mechanical properties of AM-fabricated Ti alloys are influenced by several factors, understanding the relationship between these properties and the AM techniques used remains a significant challenge. As such, this review highlights the need for further research to optimize processing parameters and microstructure design to achieve a more desirable balance of mechanical properties. In addition, efforts to develop novel materials and alloys tailored for AM processes will unlock new possibilities for creating advanced components with superior properties. To fully unlock the potential of emerging multi-laser AM in improving production efficiency and utilize the capabilities of multi-material AM to create high-quality multifunctional Ti alloys for diverse industrial applications, it is important to prioritize and conduct more research in those areas. Researchers also need to improve their understanding of the fundamental mechanisms governing the relationships between powder characteristics, processing conditions, underlying principles of AM techniques, microstructure evolution, and resultant mechanical properties. \section*{Declaration of Competing Interest} The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. \section*{CRediT authorship contribution statement} H.Y. Ma: Writing - original draft, Writing - review \& editing. J.C. Wang: Writing - original draft, Writing - review \& editing. P. Qin: Writing - review \& editing. Y.J. Liu: Writing - review \& editing. L.Y. Chen: Writing - review \& editing. L.Q. Wang: Writing review \& editing. L.C. Zhang: Conceptualization, Writing - original draft, Writing - review \& editing, Supervision. \section*{Acknowledgements} The authors would like to acknowledge the financial support provided by the industrial grant (No. G1006320). J.C Wang is grateful for the support of the Forrest Research Foundation PhD scholarship. The authors would like to thank the Australian Government Research Training Program Scholarship. The authors also acknowledge the facilities, and the scientific and technical assistance of the Australian Microscopy \& Microanalysis Research Facility at the centre for Microscopy, characterisation \& Analysis, The University of Western Australia, a facility funded by the University, State and Commonwealth Governments. \section*{References} [1] A. Sargeant, T. Goswami, Mater. Des. 27 (2006) 287-307. [2] P. Majumdar, S.B. Singh, M. Chakraborty, J. Mech. Behav. Biomed. Mater. 4 (2011) 1132-1144 [3] Y.J. Liu, H.L. Wang, S.J. Li, S.G. Wang, W.J. Wang, W.T. Hou, Y.L. Hao, R. 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A 802 (2021) 140677. \begin{itemize} \item \end{itemize} \end{document} \documentclass[10pt]{article} \usepackage[utf8]{inputenc} \usepackage[T1]{fontenc} \usepackage{amsmath} \usepackage{amsfonts} \usepackage{amssymb} \usepackage[version=4]{mhchem} \usepackage{stmaryrd} \usepackage{graphicx} \usepackage[export]{adjustbox} \graphicspath{ {./images/} } \usepackage{hyperref} \hypersetup{colorlinks=true, linkcolor=blue, filecolor=magenta, urlcolor=cyan,} \urlstyle{same} \usepackage{multirow} \title{A comprehensive review on laser powder bed fusion of steels: Processing, microstructure, defects and control methods, mechanical properties, current challenges and future trends } \author{} \date{} %New command to display footnote whose markers will always be hidden \let\svthefootnote\thefootnote \newcommand\blfootnotetext[1]{% \let\thefootnote\relax\footnote{#1}% \addtocounter{footnote}{-1}% \let\thefootnote\svthefootnote% } %Overriding the \footnotetext command to hide the marker if its value is `0` \let\svfootnotetext\footnotetext \renewcommand\footnotetext[2][?]{% \if\relax#1\relax% \ifnum\value{footnote}=0\blfootnotetext{#2}\else\svfootnotetext{#2}\fi% \else% \if?#1\ifnum\value{footnote}=0\blfootnotetext{#2}\else\svfootnotetext{#2}\fi% \else\svfootnotetext[#1]{#2}\fi% \fi } \begin{document} \maketitle Review \textbackslash author\{\\ Shubhavardhan Ramadurga Narasimharaju ${ }^{a}$, Wenhan Zeng ${ }^{a}$, Tian Long See ${ }^{b}$, Zicheng Zhu ${ }^{c}$, \\ Paul Scott ${ }^{\mathrm{a}}$, Xiangqian Jiang $\left(\right.$ Jane $^{\mathrm{a}}$, Shan Lou ${ }^{\mathrm{a}, "}$ \\ ${ }^{\text {a }}$ EPSRC Future Metrology Hub, University of Huddersfield, Huddersfield HD1 3DH, UK \\ b The Manufacturing Technology Centre (MTC), Ansty Park, Coventry CV7 9JU, UK \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-01} \end{center} \section*{A R T I C L E I N F O} \section*{Keywords:} Laser powder bed fusion process Steels Thermo-physical phenomena Microstructure Mechanical properties Post-process treatments Current challenges Future trends \begin{abstract} A B S T R A C T Laser Powder Bed Fusion process is regarded as the most versatile metal additive manufacturing process, which has been proven to manufacture near net shape up to $99.9 \%$ relative density, with geometrically complex and high-performance metallic parts at reduced time. Steels and iron-based alloys are the most predominant engineering materials used for structural and sub-structural applications. Availability of steels in more than 3500 grades with their wide range of properties including high strength, corrosion resistance, good ductility, low cost, recyclability etc., have put them in forefront of other metallic materials. However, LPBF process of steels and iron-based alloys have not been completely established in industrial applications due to: (i) limited insight available in regards to the processing conditions, (ii) lack of specific materials standards, and (iii) inadequate knowledge to correlate the process parameters and other technical obstacles such as dimensional accuracy from a design model to actual component, part variability, limited feedstock materials, manual post-processing and etc. Continued efforts have been made to address these issues. This review aims to provide an overview of steels and iron-based alloys used in LPBF process by summarizing their key process parameters, describing thermophysical phenomena that is strongly linked to the phase transformation and microstructure evolution during solidification, highlighting metallurgical defects and their potential control methods, along with the impact of various post-process treatments; all of this have a direct impact on the mechanical performance. Finally, a summary of LPBF processed steels and iron-based alloys with functional properties and their application perspectives are presented. This review can provide a foundation of knowledge on LPBF process of steels by identifying missing information from the existing literature. \end{abstract} \section*{1. Introduction} Since their inception, steels and iron-based alloys have been the leading engineering materials for structural and sub-structural applications [1]. The steels have become part of our day-to-day life, and their importance to our society is extensively revealed by their plenitude of applications. These applications include aerospace, automotive, medical, machinery, nuclear reactors, marine/oil and gas, shipbuilding, food and transportation, electronics and consumer applications [2,3]. According to the World Steel Association, there are over 3500 different grades of steel produced based on their applications, encompassing unique physical, chemical, and environmental properties [4]. Availability of steels in numerous grades has increased their array of properties including higher strength, higher corrosion resistance, good ductility and toughness, low cost and nearly $100 \%$ recyclability etc. [5]. Among the steels family, low carbon alloy stainless steels (SS), particularly 316L SS have been one of the most widely used type due to low cost, ease of processing, good corrosion resistance and excellent toughness even in severe working conditions. The outstanding combination of good corrosion resistance, higher strength and higher \footnotetext{\begin{itemize} \item Corresponding author. \end{itemize} E-mail addresses: \href{mailto:shubhavardhan.narasimharaju@hud.ac.uk}{shubhavardhan.narasimharaju@hud.ac.uk} (S.R. Narasimharaju), \href{mailto:Z.Wenhan@hud.ac.uk}{Z.Wenhan@hud.ac.uk} (W. Zeng), \href{mailto:TianLong.See@the-mtc.org}{TianLong.See@the-mtc.org} (T.L. See), \href{mailto:zicheng.zhu@strath.ac.uk}{zicheng.zhu@strath.ac.uk} (Z. Zhu), \href{mailto:p.j.scott@hud.ac.uk}{p.j.scott@hud.ac.uk} (P. Scott), \href{mailto:x.jiang@hud.ac.uk}{x.jiang@hud.ac.uk} (X. Jiang), \href{mailto:S.Lou@hud.ac.uk}{S.Lou@hud.ac.uk} (S. Lou). } mechanical properties are the important features of martensitic type steels. Martensitic type steels such as precipitation-hardened $(\mathrm{PH})$ steels (17-4PH \& 15-5PH) are basically used in aerospace, chemical, petrochemical, food processing, general metal working, oil \& gas, powerplant and injection molding industries [6]. The combination of good corrosion resistance with higher hardness, yield strength and ductility, good weldability and abrasion resistance are necessary for tools and die making industry, tool steels fulfill this criterion. Most commonly used tool steels in metal AM process are the carbon-free maraging steels (18Ni-300) [7]. In addition to splendid high temperature tensile properties, creep resistance and favorable irridation resistance makes oxide dispersion strengthened (ODS) steels perfect candidates for high temperature turbine blades and heat exchanger tube applications [8]. A taxonomic classification of steels along with their applications is shown in Fig. 1 [9]. In addition to major class of steels (tabulated in Table 1), some of the less studied steel types used in LPBF process are martensitic steels, TRIP/TWIP steels, silicon based (Fe-Si), nickel based (Fe-Ni), and cobalt based (Fe-Co) alloy steels, China low activation martensitic (CLAM) steel and etc. \subsection*{1.1. The scope of the review} This article is focused to fill the de facto gap by reviewing steels and iron-based alloys used in LPBF process. Firstly, the basics of thermophysical phenomena operative during LPBF process, solidification by phase transformation, and formation of metallurgical defects and their potential control methods are discussed. Secondly, microstructure, wear and surface texture characteristics, mechanical properties are reviewed.\\ Furthermore, the significance of post-process treatments on LPBF processed steel components are enumerated. In particular, we concentrate to critically review on how the typical LPBF process parameters have an absolute impact on the formation of; (i) different type (size, morphology) of microstructures, and (ii) process related metallurgical defects. Consequently, how these two combinations have the direct influence on wear and surface texture characteristics and finally on mechanical properties such as hardness, tensile and fatigue properties of asbuilt and post processed LPBF fabricated steels and iron-based alloys. This article also describes the current state of the art, technological challenges, and future trends, with special emphasis on AM, forecast of AM technology, and its applications in various industrial sectors. We intentionally do not discuss the details of all types of AM process, instead we restrict our review just to LPBF process. However, other AM processes such as Electron Beam Powder-Bed Fusion (EPBF), Directed Energy Deposition (DED) processes are equally capable of fabricating plethora of steels. Similarly, this review is largely limited to commonly used steels and iron-based alloys; the overwhelming majority ( $>90 \%$ ) of referred articles to review this article are concerned with LPBF process of steels. Except in a very few countable occasions, conventional process or other AM process of other metallic alloys have been cited where relevant. Additionally, this review does not extensively cover AM of other similar/dissimilar metal alloys or metal matrix composites (MMC). Discussing all of these would further lengthen this review excessively. \subsection*{1.2. The organization of the article} The goal of this paper is to provide a critical overview for readers to \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-02} \end{center} Fig. 1. Taxonomy scheme for various steels. [Based on data provided in Tables 11.1(b), 11.2(b), 11.3, and 11.4, [9]]. Table 1 Chemical composition, mechanical properties of major class of steels fabricated in LPBF process. \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|c|c|} \hline Type & Common name/grade & $\mathrm{C}$ & $\mathrm{Cr}$ & Mn & Mo & $\mathrm{Ni}$ & $\mathrm{Si}$ & $\mathrm{Ti}$ & Others & $\mathrm{TS}(\mathrm{MPa})$ & Elongation (\%) \\ \hline Austenitic stainless steel & AISI 316L & $<0.08$ & $\sim 17$ & $<2$ & $\sim 2.5$ & $\sim 13$ & $\sim 0.75$ & - & - & $\sim 310$ & $\sim 30$ \\ \hline Duplex stainless steel & SAF2705 & $<0.03$ & $\sim 25$ & $<1.2$ & $\sim 4$ & $\sim 7$ & $<0.8$ & - & - & $\sim 900$ & $\sim 25$ \\ \hline PH stainless steel & $17-4 \mathrm{PH}$ & $<0.07$ & $\sim 17$ & 1 & - & $\sim 5$ & 1 & - & - & $\sim 1400$ & $\sim 16$ \\ \hline Maraging steel & 18-Ni300 & $<0.03$ & $<0.5$ & $<0.1$ & $\sim 5$ & $\sim 18$ & $<0.1$ & $\sim 0.7$ & $\sim 9$ Co & $\sim 2050$ & $\sim 8$ \\ \hline Carbon bearing steel & AISI H13 & $\sim 0.4$ & $\sim 5$ & $\sim 0.5$ & 1.15 & & $\sim 1$ & - & - & $\sim 1600$ & $\sim 9$ \\ \hline ODS steel & PM200 & $\sim 0.07$ & 19 & 0.07 & 0.13 & 0.03 & - & 0.5 & $0.5 \mathrm{Y}_{2} \mathrm{O}_{3}$ & $\sim 875$ & $\sim 15$ \\ \hline \end{tabular} \end{center} gain profound knowledge about the LPBF process of various steels. The review first enlightens introduction to Steels, AM, LPBF and their respective applications. Section 2 is organized to describe the important process parameters, and complex thermophysical phenomena that influences the phase transformation, and evolution of microstructure in LPBF process. A thorough discussion on the defects formation, potential control methods, and common issues that arise during LPBF processing of various steels are addressed in Section 3. Section 4 seeks to critically examine the microstructure, wear and surface texture characteristics, mechanical behavior, i.e., hardness, tensile, and fatigue properties of LPBF of steels on various combined process parameters. Effect of postprocess treatments on LPBF processed steels are investigated in Section 5. Finally, Section 6 highlights the summary and future scope. It is therefore hoped that this review will help in understanding the current state of the LPBF technology, the scientific knowledge gaps and the research mostly required for the advancement and extension of LPBF process of steels. \subsection*{1.3. Steels in additive manufacturing and their applications} Currently, steels that are used in structural and automobile applications are mostly manufactured by conventional methods like casting, extrusion, and powder metallurgy [10,11] The products produced by these traditional processes have been widely used but many problems still persist. The reason pertained to slow cooling rates of casting process induce coarser microstructure, and the defects related to inherent characteristics (porosity, part shrinkage) can subsist simultaneously, which collectively annihilate the mechanical properties [12-14]. Besides, fabrication of steels in the standard process which is time consuming due to a series of independent processes (materials preparation, production and assembly) making it less flexible. With extensive development in manufacturing, special attention has to be paid towards structure-performance requirements of steel components. For example, cellular or lattice type steel structures that are primarily used in working at elevated temperatures under extreme environments (missiles, aircrafts applications) to thwart from oxidation, corrosion while retaining their mechanical integrity [15,16]. Constituent fabrication of complex, functionally graded materials (FGM) for structural components in AM offers greater advantage of saving time, costs and the flexibility (see Fig. 2). More importantly, AM process reduces the weight and stress concentration factors associated with other conventional welding and joining techniques $[17,18]$. Despite the fact, some of the traditional manufacturing issues still exist in AM process, but the comparative analysis reveals that AM process or LPBF process, have been successful in fabricating defect free (minimum number of process related metallurgical defects) good quality parts exhibiting excellent mechanical properties as compared to conventional processes like casting, extrusion processes [19]. The higher strength is attributed to the combined effect of (AM process induced) refined microstructure (dendritic, cellular type of grains), and potential high dislocation density caused during rapid solidification [19-31]. As the technology continues to advance exponentially, the manufacturing process is no longer about just producing physical products. A fundamental shift is imperative to meet the change in consumer demands, nature of products, and the economics of production and supply chain. Data-driven models using advanced machine \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-03} \end{center} Fig. 2. Correlation between Additive manufacturing (AM) key features and its advantages. learning algorithms, added sensors and connectivity are capable of revolutionizing conventional manufacturing into smarter manufacturing. Fabrication of smart steels products by utilizing the smart and robust AM technology that possess designer surface topography and mechanical performance, highly dense and dimensionally accurate, near net shape parts with reduced requirement of postprocessing is going to be a major research and development objective in the future. AM process is classified into a family of technologies where the material is added, rather than removed to produce an end product. Unlike traditional manufacturing process which involves materials being shaped or carved into required final components by parts of it being subtracted in a variety of ways. AM herein is perceived pole opposite; three dimensional (3D) components are built directly from 3D CAD file by means of an additive strategy-based depositing or melting successive layers of the feedstock materials in an enclosed chamber of the additive manufacturing system. AM is considered as the direct manufacturing technology that gives freedom to fabricate parts from the materials composed of metals, polymers, ceramics, and composites with complex features through external and internal layout, in addition to reduced material consumption [32,33]. The materials used in AM process can be in the form of powder, wire, sheet, etc. [34,35]. AM process is often described by other terms such as additive fabrication, additive technique, additive layer manufacturing, layer manufacturing, solid freeform fabrication and freeform fabrication [36]. Out of many AM processes, LPBF process is currently the most favoured powder bed fusion method which is used to fabricate metallic materials [37]. According to SmarTech Publishing's latest metal AM report "Additive Manufacturing with Metal Powders 2018", LPBF technology is one of the most used and studied AM method [38]. Forecast of AM technologies have been constantly driving the industry revenues resulting from hardware, materials, and software. This revenue growth is predicted (by the Wohler's report 2020) to be worth of \$US 16 billion in 2020, growing to \$US 40.8 billion in 2024 (see Fig. 3) [39]. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-04} \end{center} Fig. 3. Forecast of AM Industry growth (Wohler's report 2020) [39]. Metal AM technology has attracted many researchers and industries because of its distinctive applications. In recent years metal AM is used to fabricate end-use products of the medical devices (dental restorations, medical implants), aerospace and military applications, automobile industrial and consumer applications [40] (see Fig. 4a \& b). AM is also expanding its territory into aircraft maintenance and transportation sector by production of spare parts and refurbishing the damaged components [41-44]. \subsection*{1.4. Laser powder bed fusion process of steels} Laser powder bed fusion process is also known as Selective Laser Melting that uses a high-power laser beam to selectively melt the predefined contours in subsequent layers of powder. The molten metal pool rapidly solidifies by cooling [46]. Selected regions in each layer are melted by a laser beam, to form a 3D cross-section of the final part. Consequently, the underlying build platform is lowered down, followed by deposition of another layer of powder with the powder coater/wiper mechanism. This cycle is successively repeated until the threedimensional solid object is built. The unfused powder is removed and recycled, this entire process is carried out inside a chamber filled with atmospheric gas (Argon, nitrogen), to avoid oxidation (see Fig. 5). Some applications of LPBF process are shown in Fig. 6. LPBF fabricated products possess higher density with refined microstructure, which contributes to the excellent mechanical properties, superior surface quality and dimensionally accurate final parts. Such a layer-wise production approach offers LPBF process an edge over conventional process in enabling consolidated parts with elaborated internal features \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-04(1)} \end{center} Fig. 5. Schematic illustration of the LPBF process.\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-04(2)} Fig. 4. (a) Categories and (b) Industrial sectors of AM applications based on Wohler's report 2019 [45].\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-05} Fig. 6. Various LPBF produced metallic part applications; (a) orthopaedic implant, (b) car steering knuckle, (c) engine mount cooling channel, (d) aircraft engine blades (e) formula student racing engine [43,44]. for complex assembly, higher production rate, reduced design iterations, and quicker introduction of new products/protypes to the market which were previously considered unfeasible to manufacture fuctional end-use products promptly [47-52]. The transition from rapid prototyping to fabricating final products also display numerous technological barriers such as part variability, incomplete knowledge related to structureprocess-property correlation. Meanwhile, LPBF process undergoes complicated thermodynamic and heat transfer mechanisms. The surface finish of the scan track is uncontrollable and unpredictable during the printing process which eventually affects the final quality of LPBF products [53]. Oxidation of feedstock materials, process induced inevitable thermal residual stresses generated during complex thermophysical phenomena [54], are the most commonly occurring problems. Dimensional accuracy from a design model to the actual part is another issue faced by LPBF technology. Although as-built LPBF components can be directly used as functional parts, aforementioned inherent problems need to be addressed prior to the fabrication of standalone parts, which should overcome to render a reliable, scalable, and high throughput widely adopted LPBF technique as a viable fabrication process. The laser interaction with the metallic powder generally leads to the formation of a smaller size molten pool approximately $0.9-1.4 \mathrm{~mm}$ in length, $0.16-0.63 \mathrm{~mm}$ in depth, and $0.12-0.38 \mathrm{~mm}$ in width respectively depending upon various LPBF process parameters $[55,56]$. The cooling rates can reach up to $10^{3}-10^{8} \mathrm{~K} / \mathrm{s}$ due to very fast movement of the laser beam $[57,58]$, again relying on the LPBF processing parameters, type of the material used, and its various physical and chemical properties [59]. Such a high cooling rate can sometimes impede grain growth and segregation of alloying elements. Along with mixing and stirring action of Marangoni convection, and particle accumulated structure formation mechanism, a thin, continuous and unique meta-stable cellular microstructure or in some cases even amorphous microstructure is formed in the molten metal pool [60]. The thin continuous refined microstructure formed is responsible for the significant improvement of the mechanical performance of the LPBF processed steel components. It is important to have both small and large powder particles: finer particles are easily molten and favour a relatively good part density, design quality surface finish; whereas the larger particles benefit ductility, mechanical strength, hardness and toughness $[61,62]$. Moreover, LPBF produced components typically display anisotropic microstructure at different length scales. The anisotropic microstructure is generally formed by the rapid solidification process through conduction, convection and radiation, in the direction of heat dissipation [63]. Anisotropy largely depends on the type of scanning strategy employed, base plate temperature, and the build direction [64]. The quality of LPBF processed components depends on the selection of the right parameters combination. A broad spectrum of LPBF process parameters accountable for the complex physical phenomena that is ultimately responsible for final quality of LPBF parts is illustrated in Fig. 7. The summary of these parameters grouped as input parameters, process physics, and outputs. There are more than 150 parameters that need to be considered during LPBF process which are not discussed in detail here. However, some of the most important process parameters (laser power (LP), scan speed (SS), hatch spacing (HS), and layer thickness (LT), atmospheric chamber gas and pressure), and their impact on various physical and mechanical behaviours of LPBF steels are discussed. In the design parameter chart, the output represents the final quality of the LPBF processed product, listed as part geometry, microstructure, mechanical properties, defects, surface roughness, etc. It is widely known fact that LPBF processed parts primarily exhibit significant anisotrophy in microstructure as aforementioned. For example, components built in different directions i.e., parallel (e.g. horizontal) or perpendicular (e.g. vertical) to the substrate undergo a different thermal history which leads to anisotropic mechanical properties, and different surface texture (finish) [52,65]. The ratio between hatch spacing and spot size plays a major role on the process stability that affects the quality of the LPBF products [66]. As a result of opting smaller hatch spacing, a continuous and thin layer is formed due to heat accumulation and slow cooling process in a molten melt pool [67]. In contrast, fully dense, good quality LPBF products were produced even with the selection of large hatch spacing combined with unusually high\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-06} Fig. 7. A detailed process design parameter(s) of the LPBF process. energy density and increased processing scan speeds [68]. To attain the process stability and good quality LPBF products, it is recommended to choose the average hatch spacing to spot size ratio between 0.6 and 1.5 [69]. Lower energy input or larger layer thickness causes insufficient energy input penetration (to achieve effective overlap) between the melt track layers that lead to the formation of lack of fusion (LOF) or incomplete fusion hole defects $[70,71]$. Similarly, at a relatively of lower scan speed and at a fixed or higher laser power, the energy input is high resulting in higher thermal stresses and keyhole porosity defects $[71,72]$. The higher energy input induces greater temperature gradient, combined with the larger thermal residual stresses frequently causes thermal cracks [73,74]. Conversely, at a relatively lower laser power and at a higher scan speed, the supplied low energy input is not sufficient to completely melt the surrounding powder particles (by wetting) leading to the formation of balling defect [75]. It is also evident that higher energy density reduces product dimensional accuracy, making it difficult for process optimization, which may lead to compromise between the specimen dimensionality and defects [76]. From the already published work, adopting higher layer thickness resulted in an decrease in relative density. Consequently, a combination of LPBF layer thickness along with the scan speed influences the microhardness [77-79]. Selection of layer thickness more than $0.1 \mathrm{~mm}$ will lead to staircase defects on curved and inclined LPBF built surfaces [80]. In addition to conducting physical experiments, computational modelling of the LPBF process is extremely important to optimize the process parameters. These models are also helpful to predict the complex temperature field of molten melt pool, development of microstructure, residual stresses, distortion, warping and etc. Some of the researchers attempted to correlate experimental and modelling results of LPBF fabricated steels [81-83]. Childs et al. investigated the link between range of laser powers and scan speeds with respect to the formed melt tracks through experiments and modelling of LPBF of M2 tool steel, H13 tool steel, and 314S-HC stainless steels [81,82]. From their research it was clearly shown that, a perfect combination of higher laser power with the lower scanning speed is necessary to achieve stable the melt tracks, as these stable melt tracks are beneficial to manufacture fully dense LPBF parts. Badrossamay et al. studied LPBF process of M2 tool steel and 316L stainless steel through experiments and simulation. Their results revealed that the thermal history of the LPBF process was responsible in ascertaining the amount of melt under the laser fluence. Also, simulation result suggested that the laser absorbtivity may increase with increase in the scan speed. The maximum power and scanning speeds used were $200 \mathrm{~W}$ and $0.5 \mathrm{~mm} / \mathrm{s}$ respectively [83]. Li et al. developed a practical multiscale modelling for instant prediction of LPBF steel part distortion [84]. Equivalent heat source was developed by micro-scale laser scan model, local residual stress field was predicted in meso-scale layer hatch model, and finally residual stress model was utilized to predict the part distortion and residual stress in macro-scale part model [84]. Contuzzi et al. evaluated the influence of LPBF process parameters on temperature distribution in a three dimensional model. The simulated results showed good agreement with the real dimension of the melted zone. It was concluded that their simulated model could be used to optimize LPBF of steel process parameters; to predict the bonding between the melt tracks, and to characterize the best building strategy [85]. Peng et al. developed the energy demand model to manufacture the LPBF steel parts (free from porosities) using critical parameters (laser power, scan speed, layer thickness and hatch spacing). The authors reported that the higher power with higher scan speeds results in a relatively thicker layer with stable molten melt pools, thereby producing high densely parts. Hatch spacing could be selected based on the actual molten pool. They suggested that this combination effectively reduced energy density, and the corresponding energy demand [86]. Further details about various LPBF modelling methods are presented in relevant subsection 2.4. LPBF research on different types of steels and iron-based alloys have been carried out mainly to examine the appropriate processing parameters that are suitable to achieve fully dense high quality components and their resultant microstructure. However, the major concern is process insight and manipulation of exact role of (each parameter or combination of) process parameters on physical and mechanical behaviours, and thus compliance with the industrial standards of engineering parts fabricated through LPBF process is not well established. Ascertaining the mechanical properties and surface roughness which are influenced by the process design parameters is also very important that can be helpful\\ to predict the quality and service of the LPBF components [87]. The use of non-optimized LPBF process parameters contribute to poor mechanical properties due to the formation of various metallurgical defects. Mechanically sound products with relatively high density, refined microstructure, and good surface quality can be produced by choosing a suitable combination of optimal process parameters [69]. Further critical review on linking crucial LPBF parameters with the resultant microstructure, metallurgical defects and mechanical performance have been discussed in the following Sections 3 and 4. In addition to the most critical LPBF process parameters, metal powder features (particles size and grain distribution, packing density) plays a significant role in determining the final part quality. A decent amount of research has been carried out in this domain [61,88-91]. Spierings et al. investigated the impact of three different particle size distributions (PSD), and different layer thickness on the surface quality and mechanical properties of LPBF stainless steels. It was found that the smaller steel powders with PSD $\mathrm{D}_{50}$ of $15.2 \mu \mathrm{m}$ and $28.26 \mu \mathrm{m}$ needed a lesser heat input to achieve 99\%-part density than that of the powders with PSD $\mathrm{D}_{50}$ of $37.70 \mu \mathrm{m}$ [61]. It was attributed to the smaller particles that are easily melted, however, the bigger particles are helpful in undergoing higher elongation before failure. Authors concluded that the PSD not only affects the part density but also affect the surface quality and mechanical properties [61]. Similar kind of results were reported by Liu et al. [89]. They confirmed that the smaller powder particles displayed better flowability resulting in higher density, good surface quality and better strength and hardness [89]. Azizi et al. examined the powder recycling implications on powder characteristics by the virtue of size, distribution, flowability and density measurements [90]. The authors reported only difference in the flowability between the virgin and recycled powders, whereas rest of the characteristics like PSD, phase homogeneity and chemical composition remained unchanged [90]. Coe et al. most recently inspected the single mode and bimodal PSD of LPBF of stainless steels with wide range of energy densities. It is worth noting that bimodal powder with PSD $\mathrm{D}_{50 \mathrm{~L}}$ of $36.31 \mu \mathrm{m}$ provided slightly higher (2\%) tap density than the single mode spherical powders. In addition, bimodal powders utilized higher laser power ( $>203 \mathrm{~W}$ ) to reach $99 \%$ relative density. Also, as-built bi-modal powders parts showed marginally higher hardness. However, bimodal powders displayed poor flowability [91]. As in the case of LPBF processed steels; various steel and iron-based powders are used as precursor materials, however, there are still lot of uncertain bases which need to be addressed. For example, what is the ideal or universal powder grain size distribution that is best suited for different LPBF processing windows with respect to different types of steel powders. The correlation between the universality of various steels powder characteristics and the processing envelope to achieve highly dense parts, possessing outstanding mechanical properties and excellent surface quality is one of intriguing area that definitely need to be explored. \section*{2. Thermo physical phenomena of laser powder bed fusion process} LPBF process generally undergoes a highly complex phenomena, governed by the kinetics and thermodynamic mechanisms that occur at different spatial and temporal time scales. The important thermophysical phenomena that take place within the powder-bed, inside the molten melt pool and in the solidified phase of a typical LPBF process is explained in the following three subsections (Sections 2.1-2.3). \subsection*{2.1. Thermo physical phenomena during the laser-powder bed interaction} The focused laser beam irradiates the surface of the powder bed, leading to the formation of more complex heterogeneous heat transfer phenomena such as powder-bed radiation (between laser beam and the power particles), convection (between the powder bed and the environment), and finally heat conduction (between the powder bed and building substrate, and/or inside the powder bed) (see Fig. 8a) [92], depending upon various physical, and optical properties of the materials. The focused laser beam absorption is governed by multiple reflections off the oblique surface particles through pores, then it is penetrated and further scattered into a greater depth which can sometimes reach the range of the powder bed layer thickness as shown in the Fig. 8b [93-95]. The photon energy is converted into thermal energy which is dissipated across the powder bed. The spatial power density distribution of incident laser beam on the powder bed is generally assumed to follow Gaussian distribution, with the associated $2 \sigma$ (standard deviation) value usually being taken as the laser beam spot size. Typical laser spot diameters vary from 25 to 100 $\mu \mathrm{m}$ with the layer thickness lies between 25 and $50 \mu \mathrm{m}$ depending on the powder morphology and the build material [48]. The choice of lasers depends on the absorptivity of the powder materials [96]. For example; polymers, ceramics, and metal oxides are usually inclined towards the use of continuous $\mathrm{CO}_{2}$ lasers with a wavelength $10.6 \mu \mathrm{m}$, whereas other continuous fibre-lasers (Nd:YAG) with a wavelength:1.1 $\mu \mathrm{m}$ is normally used for processing the metals. In general, the nominal laser power and the laser scan velocities are in the range of $\mathrm{P} \approx 50-1000 \mathrm{~W}$, and $v \approx$ $0.1-3 \mathrm{~m} / \mathrm{s}$ [97]. The number of factors that influences the overall absorption and local energy distribution includes laser power, wavelength, polarization, angle of incidence, powder temperature, surface roughness, surface oxidation and inclusions/impurities $[88,98]$. \subsection*{2.2. Thermo physical phenomena within the molten melt pool} As soon as the focused laser beam strikes the local positions on the powder surface, the melting temperature is reached and the laser beam instantly melts the powder causing phase transition from solid to liquid droplets leading to the formation of a molten melt pool, (ideally) with a continuous melt track. The formed molten melt pool undergoes very complex physical phenomena driven by buoyancy, gravity, surface tension and capillary forces, due to high thermal gradients induced by the high velocity laser beam onto the metal powders [99]. The transfer of heat within the molten melt pool is dominated by thermo-capillaryconvection or Marangoni convection which drives the molten liquid metal from the hotter laser spot to the cold rear, (see Fig. 9) influenced by the temperature dependent surface tension [99,100], and particle accumulated structure (PAS) formation mechanism [60]. Surface tension, capillary forces, wetting behavior, as well as inertia effects are considered as the primary driving forces [101-103]. Viscousity and gravity forces are considered as secondary effects that influence the melt pool kinetics, thermodynamics, geometry as well as the surrounding powder morphology by attracting or rejecting individual powder grains [102]. The interaction of both primary and secondary forces would decide the stability and the final geometry of melt track. The shape of the molten melt pools are generally controlled by the surface tension and capillary flow, and thus it can be controlled by adjusting the laser processing parameters [104]. Formation of the molten metal pool is considered as the first point of solidification microstructure. \subsection*{2.3. Thermo physical phenomena within the solidified phase} Metallurgical microstructure is instantly established when the solidification of molten melt pool begins. The solidified microstructure determines the macroscopic properties of the final LPBF built product. The phase transformation of the solidified microstructure is distinguished by the grain morphology and grain texture which are influenced by the prevalent spatial temperature gradients, cooling rates, as well as the velocity of the solidification front [106]. The solidification process in LPBF process are classified into two regions; the first region consists of the temperature field which is in direct contact with the laser beam (fusion zone), and the heat affected zone (HAZ) [107]. The first region undergoes highly complex kinetic and thermodynamic mechanisms\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-08} (b) Fig. 8. Schematic illustration of thermophysical phenomena in LPBF process (a) various heat transfer phenomena, (b) interaction between laser beam and powder bed. within the molten melt pool with all the individual physical phenomena as explained in the previous sections. These distinct non-equilibrium condition during LPBF process results in formation of fine grain metastable microstructures [60], and compositions of the resulting phases, typically give rise to superior mechanical properties [108-110]. In the second region (HAZ), the thermal evolution is predominant in already deposited layers, located below the current layer and further away from the laser heat source which are exposed to prolonged repeated heating and cooling cycles resulting in solid phase transformations and grain coarsening [109-112]. \subsection*{2.4. Summary of relevant studies on thermophysical phenomena of LPBF process of steels} Based on the available literature, there are three kinds of computational models namely analytical models, empirical models and numerical models. Analytical models focus on the physics side of the process, and they tend to be beneficial to optimize process parameters. These models require high computation time to capture the complex thermophysical phenomena of the molten melt pool. Analytical methods are accountable to model the part of process physics with different prediction accuracy, they do not incorporate the multi-physics, and thus are less effective since more complex physics is involved during LPBF process. Fathi et al. studied a mathematical model accustoming a parabolic equation to build the molten melt pool's top surface during laser powder deposition [113]. The temperature distribution inside the clad and substrate was acquired by solving the heat conduction equation based on an infinitely fast-moving laser heat source. This model enabled to the predict molten melt pool depth, temperature field and the dilution as a function of clad height and width [113]. Mirkoohi et al. investigated a three-dimensional (3D) semi-elliptical model with moving heat source approach to predict the in-process temperature profile inside LPBF processed part [114]. The authors further studied the effect of time spacing (laser pulse), the impact of number of scans and hatch spacing on the thermal properties and the molten melt pool geometry. From this analytical model, few details were considered to predict the geometry of the molten melt pool more precisely and realistically [114]. Lee et al. developed a novel hybrid heat source model to predict and analyse melt pool characteristics including molten melt pool dimensions and melting modes of LPBF processed steels [115]. This formulated hybrid model \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-09} \end{center} Fig. 9. Schematic illustration of Marangoni convection (flow) during LPBF process [105]. considered different absorption mechanisms for the porous and densestate materials, and an effective absorptivity was employed to analyse melting mode transitions [115]. Empirical models are more case sensitive, but are time-efficient compared to numerical models. Numerical modelling method in LPBF process draws great research interest in both academia and industry fraternity. Numerical methods are used to simulate the laser interaction with powder particles, molten melt pool formation and the thermal stress field distributions in LPBF process [116-121]. Numerical methods in LPBF process are categorised as microscopic, mesoscopic, and macroscopic scale models. As the name indicates microscopic models usually deals with thermo-mechanical phenomena that takes place in microscopic level such as energy absorption, heat conduction, convection, radiation, thermo capillary effects, Marangoni effect and recoil pressure [116]. Additionally, microscopic models also comprised of stable/unstable phase transformations, microstructure evolution involving size, shape and the orientation of resultant grains in the molten melt pool. Ninpetch et al. developed a power scale computational fluid dynamics (CFD) model to study the thermal behavior, molten metal pool flow characteristics of LPBF processed steel, and also to analyse the influence of laser power, scan speed on the melt track formation [116]. It was revealed from the results that CFD model could be used to describe the complex thermophysical phenomena like heat transfer and molten melt pool characteristics, along with the laser scan track width and depth entity [116]. Lindroos et al. formulated a novel thermomechanical crystal plasticity model which demonstrated the microscale level structure evolution, residual stresses, and the strains in a single track LPBF processed H13 steels [117]. Their model effectively addressed the microscale residual stress anomalies that depend on molten melt pool thermal, and microstructural evolution, phase transformations and the interplay with the surrounding matrix of H13 steel. It was concluded this microscale model was exceptionally robust in predicting microstructural residual stresses and the deformation [117]. Mesoscale models are generally utilized to address the solitary grain and the complex thermo-hydrodynamic phenomena of molten melt pool during LPBF process. These types of models are helpful to study the bonding properties between the successive melt track layers, which determines the formation of process related metallurgical defects, responsible for the surface quality of the final part. L. Cao simulated mesoscale multi-layer multi path forming process to predict the molten melt pool behavior dynamics of LPBF built steel. Impact of three scan strategies on grain orientation, porosities and the surface morphology were investigated and compared with the experimental results [118]. It was found that the grain orientation of current formed layer of the first scan strategy and the third scan strategy was almost same as that of already formed layer, however, the reported grain orientation of current formed layer under the second scan strategy was significantly different from that of already formed layer. Additionally, they reported porosities and the surface morphology in scan strategy three was lower than the other two scan strategies [118]. Lie et al. developed a new ray tracing heat source 3D mesoscale simulation for LPBF processing of steels [119]. The simulated model analysed the laser interactions with the powder bed, considering the multiple laser reflections from the surface of the steel powder. It was found that the proposed ray tracing heat source model was able to simulate the laser heating process of LPBF process better than the conventional one. In addition, this model was successful in identifying the local defects such as balling [119]. Macroscopic models mainly focus on simulating the whole fabrication of LPBF parts. Macroscopic models are accountable for predicting the spatial temperature distributions, residual stresses, distortion, warping of LPBFed parts. Li et al. developed a geometry scalable predictive model across the microscale laser scan, mesoscale layer hatch and the macroscale part build-up to quickly predict the residual stresses and distortion with respect to different scanning strategies [122]. The model predictions were validated by experimental data, it was found that the geometry scalability law in context of layer thickness is achievable for the complex part geometries to predict the residual stresses and distortion without compromising the accuracy. The authors concluded that adapting orthogonal scanning pattern between the two adjacent layers was beneficial to reduce the residual stresses and distortion [122]. Shiomi et al. studied the distribution of residual stress model, and proposed base plate pre-heating, stress reliving heat treatment and laser re-scanning methods to reduce the residual stress formed during LPBF processing of steel [123]. Li et al. evaluated finite element analysis (FEA) model to predict the transient thermal stress field, and optimize LPBF process parameters to analyse these impact on residual stresses and deformation [124]. This FEA predictive numerical model could be used as an effective tool for the parametric study of LPBF process parameters, residual stresses and deformation [124]. Altogether, multi-scale computational models are therefore considered as the basic reliable tools to understand the complex thermophysical phenomena that occur in LPBF process. Concurrently, these basic reliable tools could serve as predecessor to design physical experiments. For further information related to computational modelling methods in context of LPBF process of steels; interested readers are requested to refer these articles [125-129]. Overall, publications on modelling studies of LPBF process are abundantly available, however, the research specific to the simulation of LPBF process of steels and iron-based alloys\\ are limited. \subsection*{2.5. Solidification theory of laser powder bed fusion process} To understand the formation of microstructure and property evolution of LPBF processed parts, it is crucial to ascertain solidification theory and the associated thermal behavior. During conventional welding or similar processes, nucleation begins at existing base-metal grains in the fusion line which act as a substrate, and these grains grow towards the centre of the weld by epitaxial growth (see Fig. 10). Since the molten melt pool is in intimate contact with the base-metal grains, it (molten melt pool) completely wets these basemetal grains [130]. Homogeneous nucleation typically requires larger time scales which is unachievable in LPBF process. Nucleation is commonly initiated at the solid-liquid interface between the base metal grains surface and liquid metal pool [131,132]. The solidification commences at the molten melt pool boundary and directed towards the centre of the melt pool itself [133]. LPBF process induces heterogeneous nucleation at the molten pool boundary, and epitaxy grains growth with columnar solidification front. These grains grow randomly in the direction perpendicular to the molten melt pool boundary, and along the maximum temperature gradient, that facilitates the maximum heat extraction and the highest degree of undercooling [134-136]. As a result of epitaxial nucleation, columnar dendrites or cells within each grain tends to grow in preferred crystallographic direction $\langle 100\rangle$ [137]. This is the conducive crystal growth direction or commonly observed solidification texture for cubic crystals including face-centred cubic (FCC), and body-centred cubic (BCC) metals [130,138]. The crystallographic orientation mainly depends on the scan strategies [131]. The growth of columnar grains in LPBF process of steels with a strong texture of $\langle 100\rangle$ preferentially aligned to the build direction (Z-axis) [138,140]. It is well established that directional solidification microstructure of metal alloys is determined by the effect of two apparent parameters: temperature gradient at the solid-liquid interface " $G$ ", and growth rate of the solidifying front (or solidification rate) "R" (see Fig. 11). $G$ and $R$ dominate the solidification microstructure together [137]. The ratio between temperature gradient and growth rate $(G / R)$ decides the morphology of the solidified grains, while the product of these two quantities $\left(G^{*} R\right)$ determines the cooling rate of the material within the solidification interval and therefore controls the size of the resulting microstructure [130,137]. The fast cooling and rapid solidification of the molten melt pool in LPBF process primarily depend on the energy density, and the scan velocity [141-143]. As shown in Fig. 11, the microstructure evolution by solidification undergoes a morphological transformation from the planar front to the equiaxed dendrites as the degree of constitutional supercooling increases [144,145]. The constitutional supercooling region does not exist in planar microstructure, as the temperature gradient $G_{a}$ at the front of solidliquid (S/L) interface is relatively high, but the actual temperature in liquid phase is higher than the liquidus temperature $\left(T_{L}\right)$ (see Fig. 12a). The embryos advancing are re-melted by the molten melt pool causing smooth interface without any solute segregation in grains. Due to the low temperature gradient $G_{b}$ of the liquidus phase, constitutional supercooling region is apparent, resulting in cellular microstructure. Numerous small, equidistant, and prismatic grains with hexagonal crosssection are formed at S/L interface which extends into the supercooled liquid due to unstable state of planar crystal interface (Fig. 12b) $[145,146]$. The constituent (solute) is rejected towards the lateral subgrain boundaries, as the corresponding $\mathrm{T}_{\mathrm{L}}$ of sub-grain boundary decreases. When temperature gradient $G_{c}$ is further decreased, already formed cellular crystalline microstructure penetrates deep inside the liquid for a prolonged depth, also results in constitutional supercooling in transverse direction (Fig. 12c). The coexistence of columnar or equiaxed dendrites, along with the liquid phase in a sensitive region is called as mushy zone [146]. It is quite appealing to notice that a very high degree of constitutional supercooling in this mushy zone (Fig. 12d). This phenomenon is attributed to temperature gradient, crystallization rate, and the Gaussian distribution of the laser energy, as well as distribution of the supercooling of the molten melt pool in different zones [147]. Hence, the formation of different types of grains is expected in the solidified microstructure [148,149]. Further information about the solidification theory of LPBF process can be referred to [150,151]. It is evident that laser power, scanning velocity, and different building directions affects the grain features of LPBF built parts. Elongated grains (Fig. 13a) are prevalent in the building direction, while the equiaxed grains are apparent in the transverse direction (see Fig. 13b) [152]. The faster cooling rate sometimes affect the sub-structure grain boundary formation, resulting higher hardness and wear resistance due to evenly distributed fine dendrites on a surface $[153,154]$. \section*{3. Formation of metallurgical defects and their potential control methods} Formation of metallurgical defects such as; balling, porosities, keyholes, cracks, metal inclusions, residual stresses, warping, delamination, oxidation, loss of alloying elements, denudation etc., and surface asperities namely; staircase effect, partially-melted/un-melted particles, spatters, re-entrant features [155] etc., are commonly observed during metal LPBF process (see Fig. 14). Incorrect selection of process parameters would likely introduce inevitable metallurgical \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-10} \end{center} Fusion line Fig. 10. Schematic illustration of epitaxial growth in LPBF process, similar to conventional welding [130]. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-11} \end{center} Growth rate, $\mathrm{R}$ Fig. 11. Effect of temperature gradient $G$ and growth rate $R$ on the morphology and size of solidification microstructure [130]. defects and surface asperities into LPBF fabricated part, which causes adverse effect on the resultant microstructure, surface texture, physical and mechanical properties [156]. \subsection*{3.1. Balling} Balling phenomenon is described by the Plateau-Rayleigh capillary instability, which occurs when the deposited melt track sometimes tends to break up into half-cylindrical shape or into spherical balls [157]. This phenomenon depends on the process variables such as scanning speed, surface tension, viscosity and density of the materials deposited [158]. Balling phenomena is one of the critical surface defects which is considered as the severe processing defect in LPBF process [159]. Surface tension and wettability have a greater impact on the formation of molten melt pool. The combination of surface tension and capillary forces drive the molten pool to shrink into its lower surface energy state (a sphere), when coalescence of individual melt tracks is in poor contact with underlying substrate results in the formation of balling defect [157]. In other words, balling defect can also occur when the liquid phase present along the surface and grain boundaries of molten melt pool fails to completely wet the remaining solid particles and the underlying substrate due to the presence of surface impurity [103]. The balling defect leads to pores, higher surface roughness, reduced density, causes lack of fusion between the powder particles/layers, imparts irregular melt tracks, and in some extreme conditions causes obstruction to the deposition process [159]. Severe balling phenomenon on a certain melt track layer inevitably leads to the formation of humping or ripple effect [160]. These ripples can carry forward onto the next layer resulting in lack of coalescence between the layers causing poor metallurgical bonding, and induce low part density. Ripple defect contribute to stacking of materials that can have serious impact on the surface quality of the scan track resulting in poor surface roughness of LPBF built 316L stainless steels [160]. When the laser beam incident on the powder bed, melting starts instantly at the local positions of powder particles surface. The phase transition from solid to the liquid molten 'cluster' is formed between the surrounding powder particles, causing reduction in surface area that gives rise to agglomeration. The selected laser spot size is usually bigger than the (starting) particles size. As the powder particles are melted together, smaller agglomerates gradually grow and are bound to form significantly bigger agglomerates (coarsening). As this process continues, a further reduction in surface tension of the molten melt pool tends to form a ball-shaped structures (balling). The dimensions of these formed balling structures are several times bigger than the original particle size (see Fig. 15) [159]. High surface tension and viscosity are the two important hydrodynamic forces that enhance balling initiation. Higher laser energy density induces more heat to form a bigger geometry molten melt pool and a wider region of contact with the substrate. The bigger and wider molten melt pool decreases the viscosity and increases the liquid metal flowability (wettability) thereby, limiting the tendency of balling (see Fig. 16) [161]. However, employing extreme laser power and scan velocity give rise to various detrimental effects. Excess heat input causes vaporization by over-heating the molten melt pool. As shown in Fig. 17a $\& b$, intense vaporization is generally observed at the top surface of the molten melt pool due to Gaussian beam heating and the highest recoil pressure right underneath the laser beam. The combination of excessive heating and the higher recoil pressure lead to the ejection of metal vapour jet plume in the form of hot spatters, un-melted powder particles that converted into powder splashes [158]. Laser re-melting can be employed on each of the fully molten metal layer to enhance the microstructure, thereby overcoming balling phenomena. Laser remelting is also helpful to minimize the spatters by rewetting the substrate at the expense of longer production times [162,163]. Similarly, preheating the base plate can improve the flowability between liquid metal and the substrate that results in the formation of a better metallurgical bond, and subsequently reduces the (balling) contraction effect arising from surface tension [164]. Nevertheless, excessive preheating chamber temperature causes droplet spatters which again lead to the formation cluster of partially melted powders obstructing the molten melt pool wettability. The preheating temperature during LPBF process of steels ranges from 80 to $900{ }^{\circ} \mathrm{C}$ [165]. \subsection*{3.2. Porosity} The degree of metal powders compactness is generally low. In \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-12(1)} \end{center} Fig. 12. Effect of constitutional supercooling on solidification mode: (a) planar; (b) cellular; (c) columnar dendritic; (d) equiaxed dendritic (S, $L$, and $M$ denote solid, liquid, and mushy zone, respectively) [130].\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-12} Fig. 13. Electron backscattered diffraction (EBSD) images of (a) elongated grains in the build direction and (b) equiaxed grains in the transverse direction [24]. addition, existing gas in the powder particles can easily diffuse into molten melt pool which cannot escape out of the molten melt pool surface due to rapid cooling and solidification. Thus, porosity is formed in LPBF fabricated steel parts [167]. Conversely, the gas solubility in liquid metal is commonly high at elevated temperatures which also contribute to the formation of pores. The porosity defects in LPBF process could be classified into incomplete fusion holes, keyhole/or depression defect and voids [167-170]. Incomplete fusion holes are related to insufficient energy input that fails to completely melt the metal powders and inadequate penetration of liquid metal into previously solidified layer causing poor metallurgical bonding [171]. The lack of fusion defects can range up to a few hundreds of microns which basically are irregular in shape, and are commonly formed at the melt track layers interface. If the supplied heat \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-13(1)} \end{center} Fig. 14. List of various metallurgical defects and surface asperities emerge during LPBF process. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-13} \end{center} Fig. 15. Schematic illustration of balling phenomena [159]. input is low, then the formed width of the molten pool becomes too small. The less wider molten melt pool lead to an insufficient overlap between the melt tracks. This insufficient overlap gives rise to the formation of un-melted powders at melt tracks interface (Fig. 18a \& b). As a result of incomplete fusion holes, the surface of this location becomes rough which directly obstructs the flow of molten pool causing interlayer defects. These interlayer defects gradually propagate as the process continues, finally to form a multi-layer defect $[167,168]$. The keyhole pores are usually spherical (Fig. $18 \mathrm{c} \& \mathrm{f}$ ) in shape caused by the gas bubbles trapped inside the powder particles in the powder mass. The keyhole pores are also attributed to very high laser energy density [31], which leads to the vaporization of low melting point elements within the alloy in the form of gas bubbles. The vapour bubbles can be trapped by fast moving laser beam and sometimes be easily dragged to the bottom of the molten melt pool by convective currents. The fast solidification rate does not allow these gas bubbles to arise and escape from the molten melt pool $[172,173]$. The spherical pores are formed due to trapped gases inside the powders during the powder atomization process or inside the molten pool during LPBF processes. In some cases, keyhole pore is also referred as depression defect that can exist at the end of a melt track with a width almost equal to the laser spot size (Fig. 18e) [169-171]. End-hole is ascribed to very high scan velocities, where laser irradiation time is not sufficient for a deep keyhole formation, instead, an open pore is created at the end of the melt track surface. End hole pore is usually induced by the dominant downward recoil pressure that is exponentially dependent on the temperature of the molten melt pool region which is directly under the laser beam $[169,174]$. Void formation is not entirely limited to low laser energy input. Perhaps, it also depends on the stability of the melt track. Voids could be either trapped gas pores, lack-of fusion holes or keyhole pore induced porosities [158]. Voids are characterized by inside layered morphology associated to molten melt pool boundaries (Fig. 18d). Void defects normally originate from the higher residual stresses generated by the rapid cooling of the molten melt pool, also sometimes could nurture the formation of cracks along the melt pool boundaries, leading to final segregation and void formation [167]. There is a strong possibility of voids or open porosity to occur at a higher scanning speeds due to inability of the liquid metal flow to completely fill the surrounding area, where the shielding gas is originally present (see Fig. 19). Insufficient filling of the neighbouring gaseous region and rapid cooling rates leads to the generation of voids or open porosity of several hundreds of microns at the surface and distributed along the overlapping gaps [173]. Porosities can lead to serious metallurgical defects, yield lower part density, and adversely affect the surface texture and mechanical performance of LPBF fabricated steels. The strategies used to suppress \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-14} \end{center} Fig. 16. Single track process map for the first layer of stainless-steel grade 316L [161]. balling are equally applicable to limit the porosities. For example, substrate preheating, [176] and employing laser re-melting reduces porosity. Selection of adequate process parameters that produce sufficient liquid metal and the larger molten metal pool lifetime is considered to be beneficial to eliminate the surrounding gas pores (region) during LPBF process. \subsection*{3.3. Residual stress and cracking} Residual stress is a characteristic of the thermal manufacturing processes, and parts produced by LPBF process are especially vulnerable to residual stresses. Residual stresses can also lead to the formation of various building defects associated with LPBF parts failure [177]. Higher temperature gradients and densification ratio which are attributed to the LPBF process, tend to create higher residual stresses. High thermal stresses cause surface defects and porosity that normally occur around the melt pool. In extreme cases, higher residual stresses results in part distortion, shrinkage, cracking, warping and delamination of LPBF produced part from its support structures. As a result of this, resultant mechanical properties, part density, dimensional accuracy of LPBF parts tends to be substantially compromised [84,153,177]. Thermal stresses generally occur from the temperature gradient or the solidification-induced shrinkage of adjacent laser melted zones in solidified material, thereby a decrease in thermal stresses would also result in a decreased residual stresses. Thermal stresses are mainly responsible for cracking. Based on the expansion behavior of the material heating or cooling, thermal stresses formed during LPBF process are classified into (i) temperature gradient mechanism (TGM) in the solid substrate (ii) cool-down phase of the melted top layers [178]. In the first case, the top layers of the solid substrate expand thermally when it experiences the high thermal energy gradients induced by the laser beam. The thermal expansion is restricted by the colder underlying solidified layers. This induces elastic compressive stresses in the top layers of the substrate. The thermal expansion may exceed the yield stress of the material and upend the plastic deformation of top layers in the direction of the laser energy source (Fig. 20a). However, when it reaches the yield stress point, the compressive stresses in the material causes plastic deformation of the top layers. As the plastically deformed layers cool down, printed material layer contracts and bend in the opposite direction (see Fig. 20b). As a result, the compressive stresses are converted into residual tensile stresses that induce cracking in the LPBF processed parts [178]. In the second case, already melted upper layer temperature is higher at the beginning as compared to the underlying layer. When the molten melt pool is cooled and solidified, upper layer tends to shrink to a greater extent due to thermal contraction. Although, this deformation is again inhibited by the underlying colder layers. Thus, tensile stresses are introduced in the upper layer and the compressive stresses in the bottom layers [178-180]. Due to the complexity of LPBF process and the difficulty in experimental measurement, finite element simulation methods are commonly used to predict the distribution and evolution of residual stresses [84]. Cracking in LPBF process can be divided into solidification cracking and liquation cracking (Fig. 21a \& b). Solidification cracking occurs in the terminal stages of the solidification when dendrites have almost fully grown into equiaxed grains, which are separated by a small residual liquid strip in the form of grain-boundary films in mushy zone. At this point, molten melt pool can be rather weak and thus susceptible to the cracking under tensile stresses. In simple terms, solidification cracking occurs inside molten melt pool or in fusion zone [181], when the liquid flowability is limited by the increased viscosity at a lower temperature, and the inter-dendritic liquid flow is obstructed by the solidified dendrite arms. Solidification cracking occurs when the localized tensile stresses developed across the adjoining grains overpower the ultimate tensile strength (UTS) of the completely solidified material at a certain point and temperature [182]. It is found that the effect of solidification cracking on the final clad properties is unaccountable as it commonly occurs at the top deposit surface. Solidification cracking can be eliminated by adopting laser re-melting or by machining. Liquation cracking needs to be carefully monitored as it remains in heat affected zone once it is formed [183]. Liquation cracking initiates from the weaker region, i.e. partially melted zone or at the heat affected zone (HAZ) in pre-layers, propagating through the intergranular region with the further deposition proceeding layer by layer [184]. Liquation cracking is also called as hot \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-15(1)} \end{center} (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-15} \end{center} (b) Fig. 17. Schematic illustration of spatters formation (a) hot/droplet spatter and (b) powder spatter [166]. cracking, mostly occur in alloys with high contents of alloying elements. These alloys precipitate several low-melting eutectics in HAZ and intersection regions between the layers that are re-melted above the eutectic temperature (solidus temperature) [130]. Once a liquation crack is formed, it becomes an initiation site for the crack propagation and crack gradually expands as the deposition progresses. Liquation cracking tendency depends greatly on the grain boundary misorientation, that is influenced by the stability of liquation films and local stress concentration [184]. Ductility-dip-cracking (DDC) is a key mechanism of crack initiation/ formation in the presence of high angle grain boundaries in LPBF process. It occurs at a modest temperature where ductility and tensile properties are relatively low [185]. Due to lack of diffusion in a nonequilibrium rapid solidification process, the solidus and liquidus temperatures decrease, and the temperature range of solidification becomes wider initiating DDC in LPBF process [186]. The faster melting and solidifying rates in LPBF process results in tensile residual stresses because the localized high laser energy input, coupled with the lower thermal conductivity of the powder particles [136]. A higher temperature gradient is developed right next to the laser spot. Comparing austenitic SS and low-carbon steels, the former is more susceptible to solidification cracking than the later one because of their lower thermal conduction and higher thermal expansion coefficients. Furthermore, some of the alloying impurities, like Sulphur (S), phosphorous $(\mathrm{P})$, and silicon $(\mathrm{Si})$, have a serious impact on cracking in $\mathrm{SS}$ materials. The cracking sensitivity can be reduced by decreasing $\mathrm{S}+\mathrm{P}+$ $\mathrm{Si}$ content. It is also interesting to note that a considerable amount of nitrogen value is detrimental to the solidification cracking of stainless steels [187]. The high carbon steels usually composed of a continuous martensite phase, whereas in SS, a continuous phase is often in the form\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-16} Fig. 18. SEM images of porosity defects observed in 316L LPBF samples: (a) low and (b) high magnification of insufficient fusion defect; (c) gas pore; (d) void/cavity defect; [167], (e) end-of track hole [175]. (f) An array keyhole pores at the bottom of melt tracks [169]. of retained austenite, that helps in preventing from cracking. An extreme cracking in M2 (medium-alloyed tungsten molybdenum steel) and H13 tool steels can be prevented by preheating partially or fully. The preheating becomes beneficial to suppress martensite formation. Rapid solidification generally results in the formation of finer microstructure; however, it is not sufficient to curb the segregation. But a low melting phase is sufficiently non-uniform to avoid segregation and cracking. This type of cracking is observed generally in high copper alloy 17-4 PH SS [188]. Similar cracking has also been observed in high silicon steel which was influenced by higher laser energy input [136]. In order to control the thermal stresses and cracking, the following necessary steps can be adapted. Higher heat input results in the formation of higher thermal residual stresses that causes cracking. Hence, the formation of cracks also depends on the selection of optimum range process parameters [182]. The presence of low-melting alloy elements along the grain boundaries can induce severe grain-boundary liquation cracks. Introducing some alloying elements which tend to limit the solidification temperature range can be beneficial to alter the chemical composition of the molten pool and thereby, preventing from the cracking [189]. Base plate preheating is the new enhancing tool added to LPBF process of steels, that aims to lower the thermal gradients, minimize residual stresses, which in turn results in the fabrication of higher density parts with superior mechanical and physical properties $[176,180]$. Higher cooling rates are generally avoided as they tend to induce thermal strains and reduce the time available for the liquid metal to fill the cracks [185]. \subsection*{3.4. Oxidation} The environment of the LPBF processing chamber is very important to fabricate oxides-free parts. Despite using protective inert environments and a shielding inert gas flow to limit the oxygen content in the \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-17(1)} \end{center} Fig. 19. Scanning electron microscope (SEM) image showing open porosity or voids [173]. working chamber, there is always a chance of small percentage of unwanted oxygen content $(\sim 0.1-0.2 \%)$ present during LPBF process [190]. This is due to the invisible air filling between the powder particles. Similar to the conventional metal casting process, sources of oxides formation in LPBF process arise from oxygen existing in the surrounding atmosphere entrapped inside the porosity of the powder particles. The entrapped oxygen being unable to completely vaporized from the surface caused by the extreme intermix irregular strong laser fluence flow convections of the molten metal pool [190]. Passive oxide films on the powder's surface before the melting can also be formed due to a large area being affiliated with a powder mass [191]. Oxygen content present in the powder could directly be translated into LPBF fabricated specimen\\ [192]. $\mathrm{Ti}_{3} \mathrm{O}_{5}, \mathrm{Al}_{2} \mathrm{O}_{3}, \mathrm{Cr}_{2} \mathrm{O}_{3}, \mathrm{Fe}_{2} \mathrm{O}_{3}, \mathrm{NiCr}_{2} \mathrm{O}_{4}, \mathrm{NiFe}_{2} \mathrm{O}_{4}$ are the common oxide phases formed during LPBF of maraging steels, stainless steels, and Inconel 718 metal powders respectively [191,193,194]. Generally, alloying elements in steels such as $\mathrm{Mn}, \mathrm{Si}, \mathrm{Ti}$ and $\mathrm{Al}$ display higher affinity to oxygen. These elements can be selectively oxidized on the surface of LPBF built part [190]. Maraging steel 18Ni(300), Ti and Al have the highest affinity to oxygen. Oxide phase is generally more stable than the nitride in the steels molten melt pool. A portion of Ti from the maraging steel reacts with the nitrogen to form small cubic TiN particles. TiN is most likely to be formed by higher $\mathrm{N}_{2}$ supply from the atmosphere (see Fig. 22). In addition, a combined oxide phase containing mainly $\mathrm{Ti}_{3} \mathrm{O}_{5}$ and $\mathrm{Al}_{2} \mathrm{O}_{3}$ can also be formed. The mechanism of formation of oxides, nitrides, and carbides is similar to other types of steels. The formation of nanometer range oxide films can be easily evaporated during intense stirring action of the molten melt pool by a laser beam, causing negligible damage to LPBF processed parts. On the other hand, micrometer range oxides films $(10-100 \mu \mathrm{m})$ of irregular geometry formed cannot be completely vaporized by stirring action of the laser beam and Marangoni flow. The oxide layer can grow thicker with increasing oxygen content in the atmosphere, at the same time as the layer re-melting. When re-melting of a new layer begins, the oxide film formed previously breaks down, and part of this oxide hovers on top of the newly formed layer, with the rest trapped inside the LPBF fabricated component. The trapped oxide leads to the formation of oxide inclusion [194]. The oxide inclusions become a site for some partially melted/unmelted powders entrapment. The oxide residues can have a substantial negative impact on heating, melting, and fusion of powder particles, thereby affecting the stability of the molten melt pool [190]. Thick oxide inclusions increases the surface tension effects, limit the absorption of the laser energy and wetting of substrate, obstructs molten pool flowability. These oxide inclusions also result in the formation of \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-17} \end{center} Fig. 20. Schematic of thermal gradient mechanism of residual stress in LPBF: process (a) heating; (b) cooling [84]. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-17(2)} \end{center} Fig. 21. Morphology of cracking (a) Liquation and (b) solidification cracking [183].\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-18} Fig. 22. Light optical micrographs at different magnifications of the LPBFed parts built with laser re-melting under technical pure N2 atmosphere. Top (left) and side views (right) are showing the melt pool shapes and the dark grey oxides containing white parent powder particles (indicated by the white arrows) and yellow TiN inclusions (indicated by the black arrows) [194]. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.) metallurgical defects such as balling, insufficient melting between powder particles, reduces the part densification, induces cracking and consequently lowering the mechanical properties [191-193]. It is worth noting that any pickup of moisture from the environment by the feedstock powders paves way to the introduction of oxygen content into the LPBF system [166]. It is well documented in the existing literature that metal powder characteristics like flowability, tap density, compressibility, grain shape and the size distributions have significant effect on the final quality of LPBF built parts [195,196]. Hoeges et al. investigated the impact of different powder atomization methods on the quality of LPBF processed maraging steel parts [197]. High-pressure water atomization was used to produce maraging steel with niobium instead of titanium. Niobium has a lower affinity to oxygen, which is beneficial in preventing the formation of stable solid oxides/inclusions, and also optimizes the flowability of maraging steels [197]. Formation of oxide inclusions due to the oxide contamination of the powder has been discussed in LPBF process of 17-4 PH steel [198]. The major issues associated with water atomized powder characteristics include irregular particle shape, lower tap densities, and oxidized surfaces. Most of the researchers used gas atomized powders in the literature. To minimize oxidation, clean and dry powders must be used, despite maintaining sufficiently low oxygen partial pressure. However, the surface oxidation can sometimes become advantageous. An appreciable increase in absorption of $\mathrm{CO}_{2}$ laser radiation on a surface of oxidized metal powders, compared with normal powders (without oxide growth) which strongly reflected the $10.6 \mu \mathrm{m}$ radiation [191]. Formation of nanometer scale, continuous and thermodynamically stable oxides films on the surface of 316L, H13, P20 and 18Ni300 steel powders resulted in improved laser absorptivity [199]. Similarly, formation of secondary phase nano oxide particles (oxide dispersion strengthened) during LPBF processing of steels resulted in higher part density, better mechanical and physical properties $[200,201]$. \subsection*{3.5. Loss of alloying elements} In LPBF process and other laser processing technologies, vaporization is basically intense in a small region right underneath the laser beam where the temperature is high. At a very high laser fluence, the temperature at the surface of the molten melt pool is higher than the boiling point of steels, that contributes to vaporization. Vaporization leads to loss of alloying elements, resulted from the concentration and pressure gradients. The concentration of vaporized alloying elements on molten melt pool surface is higher than that inside the shielding gas [202]. The intensity of vapour pressure at molten melt pool surface is higher than the surrounding environment pressure, thus the surplus pressure drives vapours containing alloying elements to eject away from the surface [202,203]. The vaporization and segregation of alloying elements change the chemical composition of LPBF processed steels. For example, nickel, manganese concentrations were significantly reduced, while the increase in silicon and iron alloying elements concentrations were recorded during LPBF process of Invar 36 steel [158]. Similarly, nickel, manganese and chromium alloy concentrations were decreased with an increase in silicon, molybdenum alloying elements in LPBF process of 316L stainless steels. This phenomena was attributed to the concentrations of alloying elements with lower boiling temperatures decreased, while the concentrations of the other alloying elements with higher boiling temperatures increased, except the base alloying element iron [158]. Loss of alloying elements reduces part density, causes microstructural defects such as keyhole, pores, spatters, voids, cracks, un-melted tracks and exhibit a lower mechanical performance of LPBF fabricated parts $[20,158]$. Hence, minimizing the loss of alloying elements is considered as an important criterion during the laser parameters\\ optimization process. Although the laser energy density is a key process parameter, scan speed also plays an equally important role in vaporization. Vaporization of alloying the elements can be minimized by careful selection of laser power and scanning speed. \subsection*{3.6. Denudation} Powder denudation is the apparent depletion of powder particles around the solidified melt track (see Fig. 23). In a typical LPBF processing environment, the denudation is caused by the intense evaporation of the metal vapour plumes from the molten melt pool. The intense vaporization causes the pressure to drop inside the vapour plume and produces an ingoing flow of ambient gas towards the centre of the melt track known as Bernoulli's effect. This inward ambient gas flow is enough to sweep in the powder particles along its flow, which can be included in the molten melt pool or ejected with the vapour plume [204]. Adopting increased laser power, higher scan speed and atmospheric gas pressure (argon gas) results in higher powder particles depletion. On the other hand, denudation also occurs if the laser fluence is not sufficient to completely melt the powder particles, and surface tension tends to pull the partly melted powder particles into the molten pool [100]. Powder denudation leads to porosity and accumulation of un-melted/partially melted particles between the melt track layers causing rough surface [94]. Careful selection of hatch spacing is also important to refrain from linear void structures associated with powder denudation effects [205]. Denudation is critical to process optimization; hence, it is always recommended to identify the suitable process parameters that result in reduced denudation [204,206]. \subsection*{3.7. Environmental effects} Argon (Ar), Nitrogen $\left(\mathrm{N}_{2}\right.$ ), Helium (He) are the three most commonly used protective shield inert gases during LPBF process. In some cases, hydrogen $\left(\mathrm{H}_{2}\right)$ is also used as a deoxidizer to provide required protective environments. LPBF processed steel components produced under Ar and $\mathrm{N}_{2}$ environments exhibited near full density values, while the parts produced under He environment exhibited density around 90\% using the same processing parameters [207]. The reported lower dense part produced under He environment can be attributed to the shielding gas effect. The higher plasma plumes were generated in He environment above the molten melt pool restricting the laser interaction. $\mathrm{He}$ and $\mathrm{H}_{2}$ environments could block the laser irradiation, resulting in less dense parts due to the transport of low laser fluence. The formation of higher plasma plumes is due to low specific gravity. These plumes can sometimes completely obstruct the laser irradiation causing porosity defects [207]. The lower plasma plumes generated under Ar and $N_{2}$ environments maintain good contact between the laser beam and the metal powders, which resulted in near full density parts with the values over $99 \%$ [207]. However, it is commonly believed that use of $\mathrm{N}_{2}$ can react with the alloying elements present in the metal powders, forming unwanted nitrides in the solidified microstructure that tend to display detrimental impact on mechanical properties of LPBF fabricated parts. The difference in the final product densities can be related to the plasma plumes generated, that obstruct the laser beam. This can be combated by using $\mathrm{Ar}$ as the shielding gas and supplying the sufficient and continuous energy input, which can overcome the energy losses of metal vaporization and ionization processes [208]. Similarly, employing low atmospheric pressure during LPBF process of steels offers less resistance to metal vapours which causes a large number of free powder spatters (Fig. 24a). Therefore, strong environment pressure is recommended. Metal vapours that exist from the surface of the molten pool have to fight against the strong protective environment, which results in less powder spatters (Fig. 24b) [203]. \subsection*{3.8. Common issues associated with LPBF process of steels} In addition to the already discussed different process induced metallurgical defects, there are other most common issues that arise during LPBF fabricating of steels components are as follows: \begin{enumerate} \item LPBF process of steels commonly result in the formation of anisotropic microstructure along the build direction, especially orientation of defects at the interface of build layers, which affect the elongation and deter the mechanical properties [209]. \end{enumerate} \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-19} \end{center} Fig. 23. Confocal height microscope image of denuded zones around melt tracks for different laser power and $2 \mathrm{~m} / \mathrm{s}$ scan speed [204].\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-20} Fig. 24. X-ray images showing spatter counts for the same powder bed thickness with different environment pressure (a) weak environment pressure (b) strong environment pressure [203]. \begin{enumerate} \setcounter{enumi}{1} \item The chemical composition of the parent steels is going to influence the crack-susceptibility. Low melting alloy elements such as sulphur and phosphorous cause solidification cracking, while manganese can lead to localized depletion, because of its high vapour pressure. Other alloying elements such as silicon, titanium can cause irregular porosities. \item Hard and brittle high-carbon martensite is expected to form during LPBF process of low carbon steels, due to subsequent rapid cooling, which significantly degrades the mechanical properties. \item LPBF processing of high carbon steels is more difficult than lower carbon steels, due to the fact that higher residual stresses are induced during rapid cooling by solidification shrinkage and thermal contraction. In addition to LPBF process induced residual stresses, there is higher possibility of forming undesirable (martensite) microstructure. Combination of martensite and hydrogen promotes hydrogen cracking. \item The higher carbon content promotes higher hardness levels and lower toughness and hence a greater susceptibility to hydrogen cracking in LPBF process of heat treatable low alloy steels. \item The possibility of formation of oxide films and passive carbides on the powder's surface before melting is greater due to the affiliation with powder mass in a large area which affects the surface quality of LPBF fabricated steel parts. \item Formation of large molten melt pool attracts more powder particles that reduces wettability due to presence of oxides and carbides; which favours the formation of defects. In addition, larger melt pool induces higher grains boundaries which becomes a site for liquation cracking. \item Poor flowability of steels powders can block the spreading of powder particles, that affects the continuity of layers thickness and induces surface roughness in LPBF produced part. Exothermic oxidation of steel powders increases the volume of the molten pool leading to a high degree of melt track instability and balling defect. \item In a broader view, it is difficult to produce large components for aerospace, marine, and other industrial applications as the existing LPBF systems are limited to manufacture small and medium size parts due to building chamber size constraint $(300 \mathrm{~mm} \times 300 \mathrm{~mm} \times$ $350 \mathrm{~mm}$ ). \end{enumerate} Based on the existing literature on LPBF of different steels, three LPBF processing windows have been proposed such as the lower processing window, higher processing window and finally the optimum processing window (see Fig. 25). Additionally, impact of the respective processing windows on the final part quality is outlined. Laser power;100-200 W, scan speed; 500-1500 mm/s, layer thickness; 40-60 $\mu \mathrm{m}$ and hatch spacing; $75-100 \mu \mathrm{m}$, chamber gas pressure at $0.1 \mathrm{~atm}$, oxygen content less than $0.1 \mathrm{vol} \%$ are deemed to be optimum processing window parameters, as mostly tend to satisfy all the required constraints (refined microstructure $<1 \mu \mathrm{m}$, fewer defects, part density $>98 \%$, surface roughness $<25 \mu \mathrm{m}$, hardness $>600 \mathrm{HV}$, yield strength $>750 \mathrm{MPa}$, and tensile strength $>550 \mathrm{MPa}$ etc.). Selection of these parameters does not necessarily yield the same results as quoted, concurrently, choosing parameters outside this range could yield excellent results as universality of LBPF machines are not defined. It is believed that this proposed LPBF processing windows gives an overall basic idea (of the role) of most important parameters on the final part quality, and would also act as reference while selecting the appropriate or right combination or the optimized set of process parameters to achieve the superior final part quality. \begin{enumerate} \setcounter{enumi}{3} \item Microstructure, wear and surface texture characteristics,\\ mechanical properties of LPBF processed steels \end{enumerate} \subsection*{4.1. Microstructure characteristics} Microstructure evolution during LPBF is not trivial. It is impossible to attribute the microstructure characteristics of a specific type of steel to all other types of steels. However, it is necessary to understand the general aspects of microstructure evolution in LPBF process of steels for further research. Tan et al. studied the microstructure evolution of LPBF process of maraging steels in both horizontal and vertical planes $[141,210]$. The authors noticed a massive submicron sized hexagonal cellular grains uniformly distributed at the centre, and a needle-shaped elongated grains prevalent at the boundaries of the melting tracks (perpendicular to the scanning direction) [141]. These microstructure characteristics would form in response to the instant melting and rapid solidification at higher cooling rates during LPBF processing of maraging steels (see Fig. 26a). In a horizontal plane, heat input decreases exponentially when the solidification rate $(R)$ is increased. This is due to the temperature dependent thermal flux generated by laser fluence would be much higher at the centre of melt track, as compared to the thermal flux at the boundaries as a result of heat dissipation [210]. Owing to the simultaneous action of higher heat dissipation and faster cooling rate, the temperature of liquid metal $\left(\mathrm{T}_{\mathrm{L}}\right)$ at this point reaches well below the melting point $\left(\mathrm{T}_{\mathrm{M}}\right)$ at the centre, and the degree of undercooling ( $\triangle T=T_{M}-T_{L}$ ) is sufficiently high enough for the new\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-21} Fig. 25. Label of LPBF processing windows and their effects on final part quality. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-21(1)} \end{center} Fig. 26. Microstructural evolutions of LPBF fabricated specimens: (a) the characteristic morphologies of the horizontal and vertical cross-sections; (b) the schematics and formation mechanism of the cellular crystals and elongated acicular crystals; (c) schematics and formation mechanism of the microstructures in the molten pool and overlapped area [210]. grains to nucleate in random orientations [210]. Furthermore, the growth rate of the crystal nucleus is consistent in all directions resulting in easy formation of equiaxial crystal grains. The equiaxial crystals exhibit hexagonal cellular structures as seen in Fig. 26b. The formation mechanism of various crystals morphologies in a vertical plane is shown in Fig. 26c. Planar solidification structure could be observed at the bottom of the molten melt pool ( $\mathrm{G}$ is maximum $\& \mathrm{R} \sim 0$ ). As $\mathrm{G} / \mathrm{R}$ ratio decreases with the gradual increase in $R$, ascending from the bottom of\\ melt pool in layer stacking/or building direction cellular dendritic structure is visible. A further decrease in $\mathrm{G} / \mathrm{R}$ value to reach the middle of the molten melt pool, the cellular structure is prevalent followed by a finer/coarser equiaxed crystal that is predominantly evident at the boundaries of the molten metal pool [141,149,210]. Boes et al. reported the heterogeneous dendritic microstructure consisting of low thermal gradients induced fine equiaxed grains, and the elongated dendrites influenced by higher thermal gradients at lower solidification rate [211]. Microstructures of LPBF built 316L steel parts characterized by the columnar grains of austenite with intercellular segregation of $\mathrm{Mo}, \mathrm{Cr}$ and $\mathrm{Si}$ alloying elements, resulted in the formation of non-equilibrium ferrite [140,212]. The occurrence of sub-grain cellular structure (less than $1 \mu \mathrm{m}$ ) can be mainly related to the microsegregation of primary elements such as Mo, V and C, due to the Marangoni convection and the difference in temperature between the inside and outside of the molten metal pool [192]. Columnar grains with ferrite content $\sim 68.8 \%$, and grain orientation is predominantly in $\langle 001\rangle$ direction with an average grain length to width ratio of 11.5:1 has been reported during LPBF process of duplex stainless steels [213]. The microstructure of LPBF process of duplex steels was largely composed of ferritic with small traces of austenite and nitrides (presumably $\mathrm{Cr} 2 \mathrm{~N}$ ) nucleating at grain boundaries. It is worth noting that the growth morphology of austenite is along the grain boundaries or of Widmanstatten type [213]. A needle shaped nano precipitate martensites with width $\sim 200 \mathrm{~nm}$ and $15-50 \mathrm{~nm}$ in length were observed at $450-510{ }^{\circ} \mathrm{C}$ in LPBF processing of maraging steels $[153,210]$. A very fine microstructure $(<2 \mu \mathrm{m}$ or less) mainly consisting of $\alpha-\mathrm{Fe}(\mathrm{M})$ phase (M, $\mathrm{Cr}, \mathrm{Ni}, \mathrm{Mo})$ formed during LPBF process of nickel-molybdenum alloy steels [214]. The microstructure of the LPBF built hot work steel characterized by $\alpha$-Fe dendritic cells decorated at the grain boundaries by the carbon rich $\gamma$-Fe regions [215]. LPBF fabricated 316L stainless steels displayed a finer and equiaxed grain, which resulted in superior mechanical properties without compromising ductility [212]. (A schematic illustrating typical microstructure at various length scales formed during LPBF process of 316L SS is shown in Fig. 27a). Wang et al. accredited this combined property to superior nature of the microstructure composed of solidification cells, low \& high angle grain boundaries, dislocations, and oxide inclusions [212] (Fig. 27b-h). LPBF process of austenitic SS are almost extensively restricted to 316L SS and 304L SS. 316L and 304L SS are in a composition range where solidification front is dominated either with a primary ( $\delta$ ) ferritic\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-22} Fig. 27. Schematic illustration of typical microstructure of LPBF 316L SS. (a) Label of discovered microstructure at various length scales, (b) an electron backscatter diffraction (EBSD) inverse pole figure (IPF) revealing grain orientations, (c) SEM image showing fusion boundaries, high-angle grain boundaries (HAGBs), and solidification cellular structures, (d) transmission electron microscopy (TEM) image of solidification cells, (e) high-angle annular dark-field (HAADF) scanning TEM image (STEM) of the solidification cells shown in d, (f) EBSD acquired with a 1- $\mu$ m size ( $\mathrm{g}$ ) EBSD image of superimposed HAGBs and low-angle grain boundaries (LAGBs). Legend representation, HAGBs ( $>10^{\circ}$ ) coloured in blue and LAGBs ( $2-10^{\circ}$ ) coloured in red. Fraction of HAGBs and LAGBs are $\sim 59 \%$ and $\sim 41 \%$, (h) Kernel average misorientation (KAM) map to demonstrate local misorientation across individual grain, (i) HAADF STEM image showing segregation of Mo and Cr alloying elements in to the solidified cellular structure and low angle grain boundaries, while EDS confirms the corresponding Fe, Mo, and Cr this segregation. EDS map also confirms that these particles are predominantly rich with Si, O, and Mn [212]. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)\\ phase or with a primary austenitic $(\gamma)$ phase. LPBF processing of stainless steels exhibits fully refined austenitic microstructure, with the columnar solidification grains $\sim 1 \mu \mathrm{m}$ diameters or less [5,162,216-218]. The number of solidified columnar grains may vary from tens and/or hundreds that are very similar to the crystal orientation, collectively form a single austenite grain i.e. a material volume responsible for high angle grain boundaries (Fig. 27b \& c) The grains formed in as-built LPBF SS samples are finer than those of conventional processes [180,184,216-218]. LPBF process of this type of steels is fully austenitic and there is no conclusive evidence of any solid-state phase transformations [180,216,219-221]. The intercellular regions shows an enrichment with $\mathrm{Cr}$ and Mo, which are, however, not sufficient to stabilize the ferrite [218,212] (Fig. 27i). A strong fibre texture with $\langle 001\rangle$ crystallographic direction aligned along the build direction (i.e. against fast heat dissipation direction) was revealed during LPBF process of steels $[180,221]$. The strong texture is caused by the $\langle 001\rangle$ crystallographic direction, as it is the fastest growing direction in the solidification of cubic metals, and hence dendrites or cells grow aligned with the temperature gradient [222]. Z. Sun et al. employed a modified laser scan strategy by adopting relatively high laser power with smaller hatch spacing to improve the mechanical properties of LPBF processed 316L SS [223]. This modified approach lead to the formation of new $\langle 110\rangle$ crystallographic texture along the build direction instead of a regular $\langle 001\rangle$ texture. The modified $\langle 110\rangle$ crystallographic grain orientation favours twinning effect under deformation, as a result of this the material experiences higher strain hardening rates which profits in achieving superior mechanical properties (ductility and UTS) [223]. H. Sun et al. moved a step forward to show that it is possible to regulate crystallographic texture by carefully controlling the process parameters during LPBF processing of 316L steel. They reported crystallographic lamellar microstructure $\langle 100\rangle$ and $\langle 110\rangle$ oriented grains along the build direction [224]. As already mentioned texture control could be a reliable tool to control anisotropic microstructure in yield and tensile strength [225]. However, the strain hardening behavior is predominantly dependent on grain morphology, resulting in anisotropy in ductility despite the reduced crystallographic texture [225]. LPBF process of maraging steels displayed different solidification microstructure with cellular/dendritic sizes $\sim 0.3-2 \mu \mathrm{m}$ as compared to conventional built maraging steels [226,227]. The cellular structure in LPBF processing of maraging steels is a result of microsegregation during solidification which enriches some of the alloying elements in the inter-dendritic regions. The microstructure of LPBF produced $\mathrm{H} 13$ tool steel consists of solidification cells/dendrites with retained austenite located in the inter-dendritic regions. The observed size of the cells/dendrites was in the range of $0.5 \mu \mathrm{m}-2 \mu \mathrm{m}$. It is worth noting that there is only limited information available in the literature regarding the crystallographic texture of $\mathrm{H} 13$ tool steels and maraging steels. This could be probably related to the very weak crystallographic texture [228]. As-built LPBF processed 17-4 PH steel displayed a high fraction of austenite phase or even fully austenitic microstructure. Facchini et al. reported LPBF of 17-4 PH stainless steel contained $72 \%$ austenite and $28 \%$ martensite [229]. In addition, the presence of little traces of Nbrich carbides was unsure [188]. TEM investigations confirmed the presence of retained austenite between martensite discs. LPBF process of 17-4 PH steels (including austenitic, martensitic and ferritic steels) usually display strong crystallographic grain orientation in $\langle 001\rangle$ direction aligned along the building direction (z-axis) [230,231]. LPBF process of (TWIP/TRIP steels) high-manganese steel was investigated by [232], the microstructure consisted of mainly austenite, together with $\alpha$ - and $\varepsilon$-martensite, along the small quantity of Mn segregation was observed as compared to cast $\left(\mathrm{X}_{30} \mathrm{Mn}_{22}\right)$ steels [232]. LPBF synthesized mechanically-alloyed (ODS steels) PM2000 steels revealed that a homogeneous distribution of retained oxides. However, the average sizes of these retained oxides were in the range of $48 \mathrm{~nm}-61$ $\mathrm{nm}$, which are significantly coarser than the conventionally-produced\\ PM2000 steels (30 nm) [233]. Similar microstructure wholly ferritic in nature was reported during LPBF built mechanically alloyed PM2000 or MA956 steels [234-236]. The fully ferritic microstructure exhibited strong fibre texture with the $\langle 001\rangle$ direction parallel to the build direction. The ferritic microstructure revealed a homogeneous distribution of both finer and coarser oxides [234]. The presence of both finer and coarser oxides could be attributed to agglomeration of nanometersized oxides [236]. The strong crystallographic texture results in anisotropic mechanical performance of LPBF processed ODS steels [234]. The crystal structure of the oxides is sometimes represented by $\mathrm{Y}_{2} \mathrm{Ti}_{2} \mathrm{O}_{7}$ or $\mathrm{Y}_{4} \mathrm{Al}_{2} \mathrm{O}_{9}$ [219]. LPBF fabrication of $\mathrm{Fe}-14 \mathrm{Cr}-1 \mathrm{~W}$ powder mechanically alloyed with $\mathrm{Y}_{2} \mathrm{O}_{3}$, and $\mathrm{TiH}_{2}$ reported a similar microstructure to the one described above for LPBF process of PM2000 [201]. \subsection*{4.2. Wear and surface texture characteristics} As a result of complex thermophysical mechanism LPBF process undergoes, rougher surface finish is induced. Defects and surface asperities like thermal cracks, spatters, un-melted/partially-melted, ripple effect, staircase effect, surface and sub-surface porosities, re-entrant features emerge on LPBF parts surface which are responsible for causing the unfavorable surface finish or surface texture [155]. A thorough investigation of currently available literature on wear and the surface texture characteristics of LPBF process of steels reveal that the research is still in its early stages. Presently it is hard to relate the wear and surface texture characteristics of LPBF processed parts to the real applications. However, to expand LPBF applications into frictional pairs, it is paramount to study the wear performance of LPBF process of steels under various contact conditions [154]. Wear is defined as the loss or displacement of material from a contacting surface. The wear rate of LPBF processed steels linearly depends on the volume percentage of the porosity. Reported wear rate was 6-17\% higher than bulk steels for less dense LPBF process of steels with the presence of porosities (see Fig. 28). It is indeed possible to achieve equivalent or superior wear resistance than conventional steels if LPBF built steel components are fully dense with minimum numbers of surface defects [237]. Similarly, the higher hardness, perfectly dense, plus good wear resistance could be accomplished when LPBF processed parts exhibit least surface asperities $[238,239]$. The principal wear phenomena act as a site for crack initiation and crack growth, originating from the pre-existed surface defects that subsequently leads to the premature failure of the component at lower applied loads [239]. Surface texture at this stage is generally used to study the basic capabilities of LPBF processes, an application to specific requirements is not completely introduced. Nevertheless, surface roughness plays a key role in determining the mechanical, tribological and functional properties of LPBF processed steel components. Surface texture is defined as the geometrical irregularities exist on the surface, excluding the geometrical imperfections that contribute to the form or shape of the surface [240]. LPBF built surfaces containing surface asperities and other particles features are often characterized by using 3D optical profilometers and X-ray computer tomography (XCT), allowing the captured data to be used for 3D surface texture characterization. Narasimharaju et al. more recently investigated the impact of various build surface inclinations with respect to 3D surface texture parameters [155]. It was found that varying surface inclination combined with staircase effect and un-melted/partially melted particles exhibit a strong correlationship. Staircase effect was evident between 3 and $45^{\circ}$; above $45^{\circ}$ the staircase surface was supplanted by un-meted/partially melted particles at $90^{\circ}$ [155]. Similar kind of research was carried out by Gogolewski et al., apart from investigating multiscale analysis of surface texture quality of models for LPBF built steel [241]. Horizontal built LPBF components are governed by balling, ripple effect, spatters, while staircase (stair-steps) effect, un-melted/partially melted particles are linked with the curved or inclined surface of LPBF processed parts (see Fig. 29 a \& b). Staircase effect could be minimized by adaptively \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-24(1)} \end{center} Fig. 28. Comparison of wear rate in dry wear test condition at $120 \mathrm{rpm}, 10 \mathrm{~N}$ (a) bulk $316 \mathrm{~L}$ and (b) LPBF $316 \mathrm{~L}$ ( $175 \mathrm{~mm} / \mathrm{s}, 150 \mathrm{~W}$ ) samples. "O" marks in (a) indicate tribo-oxide film [238]. reducing the layer thickness between the melt track layers [80,242]. Lou et al. successfully examined the novel material ratio (Mr) curve as an effective analysis tool to differentiate two AM (LPBF process, high speed sintering) surface topographies, and allowing surface texture to be linked with process control and functional performance [243]. The recesses of 3D Mr curves are caused by re-entrant features (surface pores). Authors identified Vvv (valley void volume) parameter determined by the Mr2 ratio to characterize (height position) the open surface pore [243]. Lower energy density gives rise to shattered, rough and scattered porous worn surface with cracks due to insufficient fusion of powders. Similarly, excessive energy density leads to ejection of hot spatters and redepositing on the LPBF processed part surface resulting in higher surface roughness. Wear resistance and surface finish can be improved by selecting the optimized LPBF process parameters including smaller layer thickness, in addition to adopting laser re-melting, and suitable post-processing methods [244]. Partially melted powder particles on the interior surfaces can be eliminated while the surface finish and texture could be substantially improved (at least $45 \%$ Ra value) by employing chemical-abrasive flow polishing techniques [245]. Additionally, reinforcement of tungsten carbides during LPBF process of maraging steel resulted in the formation of a thin carbide layer that significantly reduced the wear rate by $>1500$ times [246]. \subsection*{4.3. Mechanical properties of LPBF fabricated steels} \subsection*{4.3.1. Hardness and tensile properties} The present studies on mechanical properties of LPBF process of steels are mostly concentrated on evaluating hardness, tensile performance and fatigue properties. Tensile and hardness properties are summarized in Table 2. Schematic overview of basic mechanical properties of most common steels used in LPBF processes and conventional processes is shown in Fig. 30. This figure intends to provide a broad overview of the results reported in the literature but does not holds good for all classes/cases of steels, and sometimes considerable dependence of the material properties on LPBF processing conditions. From the existing literature, average Vickers hardness values for LPBF processed steels range from 408 to $900 \mathrm{HV}$, which is certainly higher than wrought materials [26,213]. Residual stresses are sometimes tend to benefit the LPBF fabrication of steels. They can also improve the hardness values of a LPBF component if at a reasonable level [247]. The increase in hardness values also improves the wear resistance of LPBF built parts [248]. The refined microstructure of LPBF processed tool steel samples consisted of low martensite phase, and high content of fine carbides and the alloying elements (V, Mo, C), that are much more homogeneously dispersed in the material as compared to the as-cast state which resulted in higher hardness values [26]. The microstructure of LPBF built samples determines the mechanical properties and the difference in tensile properties along various directions is mainly due to the easy introduction of metallurgical defects into the bonding area between two adjacent melt track layers. The tensile properties of the LPBF fabricated samples along the vertical direction are inferior, as compared to those samples built in the horizontal direction [165]. In order to obtain the higher tensile properties, besides the position of the sample in the horizontal direction, the laser fluence also plays an equally important role. For a low laser fluence ( $104.17 \mathrm{~J} / \mathrm{mm}^{3}$ ), unsurprisingly, resulted in higher porosities (lack of fusion holes or crater like voids). The porosities\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-24} Fig. 29. (a) Alicona G4 image showing staircase effect [155], (b) SEM image of inclined surface illustrating un-melted/partially melted powders stuck at the step edges [242]. Table 2 Hardness and tensile properties of LPBF of steels from different literatures. \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|} \hline Materials & Condition & Hardness (HV) & Yield strength (YS) & Ultimate tensile strength (UTS) & Elongation (\%) & Reference \\ \hline \multirow[t]{4}{*}{Maraging steel (MS)} & LBPF & - & $\sim 915$ & $\sim 1165$ & $\sim 12.44$ & $[210]$ \\ \hline & LPBF aged & - & $\sim 1967$ & $\sim 2014$ & $\sim 3.28$ & $[210]$ \\ \hline & LPBF solution & - & $\sim 962$ & $\sim 1025$ & $\sim 14.40$ & [210] \\ \hline & LPBF solution aged & - & $\sim 1882$ & $\sim 1943$ & $\sim 5.60$ & $[210]$ \\ \hline 316L stainless steel (SS) & LPBF & - & 455 & 579 & $\sim 50$ & $[25]$ \\ \hline FeCrMoVC tool steel & LPBF & $\sim 900$ & - & - & - & $[26]$ \\ \hline \multirow[t]{4}{*}{316L SS} & $\operatorname{LPBF} \Psi=0$ & - & $\sim 494$ & $\sim 640$ & $\sim 56.7$ & $[27]$ \\ \hline & $\mathrm{LPBF} \Psi=45$ & - & $\sim 498$ & $\sim 606$ & $\sim 59.9$ & $[27]$ \\ \hline & $\operatorname{LPBF} \Psi=60$ & - & $\sim 536$ & $\sim 601$ & $\sim 62.7$ & $[27]$ \\ \hline & $\operatorname{LPBF} \Psi=90$ & - & $\sim 489$ & $\sim 548$ & $\sim 43.7$ & $[27]$ \\ \hline 316L SS & LPBF & $\sim 281$ & - & $\sim 590$ & $\sim 21.1$ & [170] \\ \hline \multirow[t]{2}{*}{Austenitic SS} & LPBF single & - & $\sim 346$ & $\sim 921$ & $\sim 69.9$ & $[22]$ \\ \hline & LPBF CLM & - & $\sim 387$ & $\sim 924$ & $\sim 67.5$ & $[22]$ \\ \hline 316L SS & LPBF low power & $\sim 241$ & $\sim 500$ & $\sim 625$ & $\sim 47$ & [268] \\ \hline ASTM A131 & $\mathrm{LPBF} 250 \mathrm{~mm} / \mathrm{s}$ & - & $\sim 938$ & $\sim 1037$ & $\sim 4.5$ & [171] \\ \hline steel & LPBF 300 mm/s & $\sim 241$ & $\sim 850$ & $\sim 1050$ & $\sim 4.75$ & $[171]$ \\ \hline 316L SS & LPBF & - & $\sim 1100$ & $\sim 1200$ & $\sim 20$ & $[21]$ \\ \hline 316L SS & LPBF & - & $\sim 517$ & $\sim 633$ & $\sim 74$ & $[23]$ \\ \hline 304L SS & LPBF & - & $\sim 485$ & $\sim 712$ & $\sim 61$ & $[24]$ \\ \hline Invar 36 & LPBF & - & $\sim 350$ & $\sim 400$ & $\sim 64$ & $[20]$ \\ \hline 316L SS & LPBF heat treated (HT) & - & $\sim 550$ & $\sim 620$ & $\sim 90$ & $[20]$ \\ \hline 17-4 PH SS & LPBF opt parameters & $\sim 355$ & $\sim 650$ & $\sim 940$ & $\sim 4$ & [269] \\ \hline 17-4 PH SS & LPBF & $\sim 395$ & $\sim 750$ & $\sim 950$ & $\sim 3.6$ & [269] \\ \hline 17-4 PH SS & LPBF & $\sim 475$ & $\sim 940$ & $\sim 1150$ & $\sim 2.8$ & [269] \\ \hline Maraging steel & LPBF & - & $\sim 750$ & $\sim 1200$ & $\sim 17$ & $[270]$ \\ \hline MS-10\%WC & LPBF & - & $\sim 650$ & $\sim 1000$ & $\sim 7.5$ & [270] \\ \hline 17-4 PH GA & LPBF $60 \mu$ s time & - & $\sim 1116$ & $\sim 1358$ & $\sim 5.1$ & $[271]$ \\ \hline 17-4 PH WA & LPBF $80 \mu$ s time & - & $\sim 500$ & $\sim 990$ & $\sim 3.3$ & [271] \\ \hline 316L SS & LPBF & $\sim 202$ & - & $\sim 750$ & - & $[272]$ \\ \hline \multirow[t]{3}{*}{316L SS} & LPBF HT @ 650C 2 h & $\sim 210$ & - & $\sim 700$ & - & [272] \\ \hline & LPBF & $\sim 209$ & - & - & - & [273] \\ \hline & LPBF HT & $\sim 215$ & - & - & - & [273] \\ \hline \multirow[t]{3}{*}{CLAM steel} & LPBF HT HIP & - & - & $\sim 966$ & $\sim 5$ & $[274]$ \\ \hline & LPBF & - & - & $\sim 757$ & $\sim 9$ & [274] \\ \hline & LPBF $573 \mathrm{~K}$ & - & - & $\sim 694$ & $\sim 18$ & [274] \\ \hline \multirow[t]{4}{*}{316L SS} & LPBF $873 \mathrm{~K}$ & - & $\sim 550$ & $\sim 1016$ & - & $[275]$ \\ \hline & LPBF $1273 \mathrm{~K}$ & - & $\sim 459$ & $\sim 969$ & - & $[275]$ \\ \hline & LPBF $1673 \mathrm{~K}$ & - & $\sim 440$ & $\sim 941$ & - & $[275]$ \\ \hline & LPBF/plain carbon steel substrate & - & $\sim 347$ & $\sim 836$ & - & $[275]$ \\ \hline \multirow[t]{4}{*}{Maraging steel} & LPBF/MS substrate & - & $\sim 174$ & $\sim 712$ & - & $[275]$ \\ \hline & LPBF/H13 substrate & $\sim 450$ & - & $\sim 2100$ & $\sim 15$ & $[90]$ \\ \hline & & $\sim 286$ & - & $\sim 1200$ & $\sim 13$ & [90] \\ \hline & & $\sim 608$ & - & $\sim 1180$ & $\sim 11$ & [90] \\ \hline \end{tabular} \end{center} act as the main sites for crack initiation triggering brittle fracture with limited plastic deformation, causing cracks propagation under tensile loading conditions (see Fig. 31a). It was reported that using optimized energy density $\left(125 \mathrm{~J} / \mathrm{mm}^{3}, 156.25 \mathrm{~J} / \mathrm{mm}^{3}\right.$ ) the part density reached to its maximum, and the obtained microstructure displayed decent refined dimples with numerous grain boundaries that would block dislocations movements causing the material to resist deformation resulting in higher yield strength and tensile strength (see Fig. 31b \& c) [170]. It is worth noting that LPBF fabricated steels are strengthened without losing their ductility, unlike work-hardening that improves the tensile strength by sacrificing ductility. Adapting excess energy density $\left(178.57 \mathrm{~J} / \mathrm{mm}^{3}\right)$ resulted in decreased toughness due to high degree of overheating of the molten melt pool, causing larger and shallow dimples with lower resistance to dislocations (see Fig. 31d) [170,249]. As a result of finer microstructural texture, the mechanical properties of the LPBF manufactured steels have been improved. In addition, refined microstructure provides higher resistance to the dislocation motions and other mechanisms of plastic deformation, such as sliding [167]. Owing to the high density of low-angle grain boundaries, and the fine cellular microstructures associated with LPBF processing, the yield strength (YS) of 316L stainless steel is greatly improved [19]. The unique development of crystallographic lamellar microstructure (CLM) via strengthening of LPBF built 316L steel resulted in higher YS, ultimate tensile strength (UTS), and significantly higher ductility [22]. The grain refinement of the nano-cellular structures, presence of nano-size carbides along with the negative residual stress resulted in superior YS, UTS and higher ductility [20] in LPBF fabricated steels in comparison with other conventional manufacturing like standard casting, extrusion (wrought) and laser engineered net shape processes [23,24]. Similar higher yield strength and better ultimate tensile strength and compressive strength (CS) $[\mathrm{YS}=455-640 \mathrm{MPa}, \mathrm{UTS}=579-2100 \mathrm{MPa}, \mathrm{CS}=$ $3796 \mathrm{MPa}$ ] have been achieved in LPBF fabricated steels [26,27,249]. In some cases, higher elongation [24,25], and higher toughness are reported, the reason for this is attributed to the stress induced austenite-tomartensite transformation [28,58]. It is well acknowledged that LPBF processed stainless steels often display superior YS and UTS [180,212,217,218,225,250-252] as compared to conventionally manufactured steels. The reported YS values and UTS values from the literature show high variability and are in the range between 350 and $600 \mathrm{MPa}$, and 480-800 MPa respectively. Typical YS and UTS values for conventionally processed stainless steels lie in the range of 230-290 MPa and 580-590 MPa [253]. Additionally, LPBF processed steels display a higher yield to tensile strength ratio $[180,217,250,252]$. However, some of the researchers reported lower fracture toughness and elongation to fracture; as low as $12 \%$ for LPBF built stainless steels as compared to the wrought material typically; $40-50 \%[217,254]$. Majority of the studies reported a higher elongation to failure and fracture toughness up to $67 \%$ $[180,212,218,225,252,255]$. The higher yield strength and tensile strength have been related to the finer microstructure and dislocation \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-26} \end{center} 150300450600750900105012001350150016501800195021002250 \section*{Ultimate Tensile Strength (MPa)} Fig. 30. Schematic illustration of basic mechanical properties of commonly processed steels in LPBF process and conventional process. Steels type is indicated by the field colour, whereas the field border represent the process type. (TWIP/TRIP stands for twinning/transformation-induced plasticity, PH-precipitation-hardening and ODS-oxide dispersion-strengthened).\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-26(1)} Fig. 31. Typical SEM images taken from the tensile fracture surfaces of LPBF-processed 316L specimens at different laser energy densities of (a) $104.17 \mathrm{~J} / \mathrm{mm}^{3}$; (b) $125.00 \mathrm{~J} / \mathrm{mm}^{3}$, (c) $156.25 \mathrm{~J} / \mathrm{mm}^{3}$ and (d) $178.57 \mathrm{~J} / \mathrm{mm}^{3}$ [170].\\ substructure as per Hall-Petch relation [212,218,225,251]. Also, in LPBF built steel components, the defects in the microstructure (brittle phases/inclusions, porosity) have a strong negative impact on the elongation to fracture. LPBF fabricated duplex SS 2507 resulted in much higher YS and UTS, along with a moderate elongation at fracture $(\sim 8 \%)$, and a ductile fracture mode compared to conventionally produced ones. The higher mechanical strength and average ductility of LPBF built sample was the result of exclusive ferritic microstructure [256]. TRIP/TWIP steels (unlike austenitic stainless steels) are considered fully austenitic metastable steels which display transformation-inducedplasticity (TRIP) or twinning-induced-plastic deformation (TWIP). Haase et al. studied LPBF process of high-manganese steel (X30Mn22) [232]. From their study it was revealed that TRIP/TWIP effect was certainly functional when subjected to tensile deformation of the material, and also described its anisotropy arising from the strong (LPBFtypical) fibre texture. The tested YS and UTS were found to be higher in all the directions as compared to cast and rolled standard steels (302-416 MPa vs. $275 \mathrm{MPa}$ and 906-1065 MPa vs. $894 \mathrm{MPa}$, respectively), but the elongation at fracture was lower ( $24-31 \%$ vs. $52 \%$ ). These type of steels offer high work hardening ability which makes them attractive for applications where high energy absorption, high strain hardening rates, and high ductility are required [232]. Owing to the presence of martensite and austenite in different proportions in 17-4 PH steel microstructure, hardness and mechanical properties of LPBF processed 17-4 PH steel dispersed over a wide range of values. It is worth noting that lower mechanical properties were not only due to the presence of softer austenite phase but also due to the precipitation reaction that takes place in martensite during ageing $[32,257,258]$. Overall LPBF process of 17-PH steels are typically softer and less strong than wrought and age hardened materials [259]. In general, the mechanical properties of LPBF maraging steel are comparable to conventionally produced steel materials, but not entirely identical. LPBF produced maraging steels displays equal or slightly better YS and UTS as compared to conventional ones despite the finer microstructure resulting from LPBF process [77,210,260,261]. Hardness values recorded for LPBF produced $\mathrm{H}-13$ tool steels range from $570 \mathrm{HV}-680 \mathrm{HV}$, (and marginally higher $745 \mathrm{HV}$ when measured in the skin area [262]) [192,263,264]. These values are close to or even superior than as-quenched wrought $\mathrm{H} 13$ steels, contemplating the fullymartensitic state. Many of the recorded YS and UTS values for LPBF processed maraging steel samples are significantly lower because of the extreme brittleness of this state which leads to the premature failure of tensile test specimens [68,262,264-266]. The strong crystallographic texture $\langle 001\rangle$ of LPBF processed ODS steels lead to an anisotropic mechanical behavior, i.e., ductile fracture when strained in the build direction, but brittle trans-granular fracture when strained perpendicular to the build direction [234]. Employing additional post process heat treatments, the tensile strength numbers reached somewhat closer/equal to the conventionally produced ODS steels [234]. The difference in mechanical properties of LPBF manufactured steels available in numerous grades depending upon on specific applications is attributed to its wide range of technological parameters, which lead to the formation of anisotropy of cellular dendrite microstructure and some deviation in part densities of steel samples [227,267]. \subsection*{4.3.2. Fatigue properties} LPBF processed steels are exposed to a dynamic loading condition in many functional industrial applications; Hence, a thorough understanding of fatigue behavior and characteristics is requisite to evaluate their fatigue life. However, there is only a limited number of studies available in the literature that are dealing with the fatigue properties of LPBF built steels. The most important parameters that affect the fatigue property of the LPBF processing of steels are surface finish and the building direction $[167,249,276]$. Furthermore, process parameters along with fatigue testing conditions also influence the fatigue life of LPBF built steels $[277,278]$. The fatigue limit of LPBF fabricated part mainly depends on its surface finish. It is commonly believed that fatigue crack initiation starts at the surface of metallic materials. Similar to conventionally manufactured steels, LPBF manufactured steels are greatly affected by the rough surface finish, as well as other surface defects caused by micropores, surface cracks and un-melted and partially melted powder particles that are stuck on the surface. Additionally, the unstable molten melt pool aggravates the surface roughness $[279,280]$. The higher surface roughness $\left(R_{a}\right)$ paves the way for the higher local stresses under the dynamic loading conditions, which result in lower fatigue limits and consequently reduces the fatigue life of LPBF processed steels [281]. High cycle fatigue (HCF) limit is strongly dependent on the surface roughness related defects compared to low cycle fatigue (LCF). Hence, the HCF performance of LPBF steels can be improved by decreasing the surface roughness and the defects that occur on the part surface [281-283]. Another important parameter that has a considerable impact on fatigue properties is the build direction. The direction in which the load is applied to the built layers during LPBF process defines the fatigue strength $[249,284]$. The build direction governs the size, shape and the distribution of the LPBF processed defects, such as insufficient fusion holes and porosities that are elongated perpendicular to building direction [285,286] (see Fig. 32a-d). The horizontal built components (build direction normal to the loading axis) are exposed to longer intertime intervals which experiences higher cooling rates and faster solidification (see Fig. 32d). Thus, a formation of finer microstructure and higher distribution of smaller scale porosities which causes less stress flow and concentrations around the defect that results in better fatigue limits as compared to vertical built component (build direction parallel to loading axis). The stress concentrations are maximum in vertical components due to comparatively weak interfacial bonding between successive layers and the axis of linear and planar bigger size defects (Fig. 32c). The arrangement of these defects normal to loading direction provides easy access paths for voids to grow bigger and coalescence, causing failure at lower fatigue limits. Irrespective of the building direction, LPBF built parts are generally more susceptible to the fatigue characteristics [167,277,287,288]. LPBF process induces distinct (elongated grain structure along the build direction) microstructure as compared to conventional process of steels. Also, LPBF process stimulates higher surface roughness, which is particularly detrimental to their fatigue performance [289,290]. Nehzadfar et al. studied fatigue behavior of LPBF process of steels under the influence of process induced defects such as voids, oxides, un-melted powder particles, and other surface defects causing higher roughness. Other researchers [249,291] reported that in addition to crystallographic texture and the anisotropic distribution of process defects, LPBF processed 17-4 $\mathrm{PH}$ is highly reliant on the build orientation under both low-cycle and high-cycle fatigue properties. Fatigue limit of LPBF built components depend on directions of load applied. For parallel loading to the direction of grain growth, crack path was highly tortuous, resulting in slow crack propagation and yielding high fatigue limit [290]. On the other case loading perpendicular to the grain long axis, the crack propagation along grain boundaries was straight and easy without any resistance to the applied load [180,217]. Croccolo et al. found that the fatigue properties of LPBF processed maraging steels were isotropic, and with a fatigue limit of $600 \mathrm{MPa}$ [292], which is approximately equal to $1 / 3$ rd of the static yield strength; which is in line with fatigue limit obtained for conventionally-produced maraging steels [293,294]. Isotropy is caused because of the weak texture in this type of steels as a consequence of the martensitic phase transformation. The achieved fatigue life of LPBF built H13 tool steel was significantly below that of conventionally produced steel [265]. This can be related to LPBF process induced residual stress and the higher sample surface roughness [264]. The higher surface roughness of LPBF processed steels is particularly detrimental to the fatigue performance. It is found that by employing\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-28} (d) Horizontally built Fig. 32. Schematic illustration of LPBF build directions and stress concentrations associated with it, (a) vertically built, (b) horizontal built LPBF specimens, (c) higher stress concentrations around the defect in the vertical sample and (d) fewer stress concentrations in the horizontal sample [249]. post-process surface finishing treatment, fatigue limit can be doubled [180,282]. Stress relief heat treatments and hot-isostatic-pressing (HIP) both have only a marginal effect on fatigue performance, but primarily, after surface finishing post-process treatment, the fatigue limit lies in the range of conventionally processed steels [180,217,282]. The biggest challenge of LPBF process pose is the selection of optimum set of processing parameters from a wide array of parameters. Each parameter presents its own impact on the final properties of the manufactured part. Thus, controlling and estimating the characteristics of the final product is a very daunting task. Inappropriate energy densities result in the formation of unfavorable defects, which impart local stress concentrations during cyclic loading and leads to premature fatigue failure. Even by selecting optimum laser fluence, few small entrapped spherical gas pores inevitably occur in the LPBF processed parts. However, the effect of these pores on fatigue life of LPBF processed 316L steels is unaccountable, as they are less sensitive to notch due to its higher ductility and more resistant towards defects and residual stresses [295]. LPBF process induced defects formed due to lower or extremely higher energy densities are more detrimental to HCF because of their higher level of stress concentrations [296]. \section*{5. Effect of post process treatments on LPBF process of steels} The impact of post process treatments on microstructure, surface texture and mechanical properties of LPBF built steels have been studied by many researchers $[69,210,247,273-275,287-303]$. The most commonly used post-process heat treatment processes are annealing (with vacuum, argon etc.), solution heat treatment (in air cooling, water quenching etc.), and ageing. In addition, a typical industrial densification post processing method: hot-isostatic pressing (HIP) is accustomed to drastically reduces the micro-defects, and effectively improves the uniformity of microstructure. HIP refines the microstructure by the dislocation migration, and recrystallization of grains, thereby enhancing mechanical properties of LPBF processed steels [274]. Unlike hot working processes, cold working post processing method such as shot peening can be employed. Shot peening induces compressive residual stresses, grain refinement and macro strain, thereby improving the surface roughness, microhardness, compressive yield strength and wear resistance [247]. Other post-processing methods that are mainly focused to enhance the surface integrity characteristics of LPBF built steels are finish machining (FM), vibratory surface finishing (VSF), drag finishing (DF), laser polishing, magnetic field-assisted finishing, grinding, sandblasting and electro-polishing [300]. The surface roughness of LPBF built steels can be reduced (from $8.2 \mu \mathrm{m}$ to $0.05 \mu \mathrm{m}$ ) by grinding [240]. Lower surface roughness (by 48.72\%) was reported after sandblasting in two phases [304]. Gas atomized maraging steel powders generally contains only $(\alpha)$ martensite phase. LPBF built specimens contain a large number of nano precipitates embedded in the matrix of columnar martensites along with the traces of austenite $(\gamma)$ phase [297-299]. The inherent heating and rapid cooling during LPBF process cause the phase transformation from martensite to austenite. Despite the revision of martensite to a more stable austenite phase transformation, the size and the number of austenite increases during ageing process [297]. The ageing postprocess treatment provides ample time for the initiation of intermetallic compounds that are heterogeneously precipitated in to dislocations resulting in diffusion and grain growth. Solution heat treatment effectively dissolves alloying elements into a supersaturated austenite solution. Solution heat treatment (at $820-850{ }^{\circ} \mathrm{C}$ for $1-2 \mathrm{~h}$ ) above the austenite finish $\left(750{ }^{\circ} \mathrm{C}\right)$ temperature with the age hardening (at $460-520{ }^{\circ} \mathrm{C}$ for $5-24 \mathrm{~h}$ ) leads to the formation of intermetallic precipitates [298]. These formed intermetallic precipitates induces uniform dissolution of alloying elements into austenite solid solution. Consequently, cooling the austenite results in the formation of complete\\ martensite by eliminating austenite [298,299]. The average grain size remains the same for both as-built LPBF and aged steel specimens, while the martensitic matrix grain growth and the grain orientation substantially changes in case of solution treated-aged parts. Nanoprecipitates consists of spherical nanoparticles with an amorphous outer shell and crystalline core structure (see Fig. 33a), a line scan showing the atomic composition of each element in the precipitate is shown in the Fig. 33b. Maraging steels are generally strengthened by $\mathrm{Ni}_{3} \mathrm{Ti}$ precipitate phase. In some cases, $\mathrm{Ti}$ is replaced partially by other elements such as Mo, Co or Al, depending on the composition of the alloy [297]. Salman et al. revealed a single-phase austenite in as-built LPBF followed by post-process heat treated (annealed) sample at various temperatures for 316L stainless steels (see Fig. 34a-f) [275]. Finer equiaxed sub-grained (nano-precipitate with amorphous structure rich in Mn and $\mathrm{Si})$ characteristic cellular microstructure embedded into 316L matrix resulted in both as-built LPBF and annealed SS samples. They did not report any changes in random crystallographic orientation in microstructure except the lone difference being the average cell size. Size of the cells was gradually increased with increasing annealing temperature. Higher annealing temperature caused grains and cells to grow bigger until the cellular microstructure was no longer be observed at higher temperatures ( $T \geq 1273 \mathrm{~K}$ ) [275]. Sun et al. compared the microstructure of wrought and LPBF produced 17-4 PH steels in as-built, solution heat treated $\left(1038^{\circ} \mathrm{C}\right.$ for $\left.4 \mathrm{~h}\right)$ and aged ( $482^{\circ} \mathrm{C}$ for an hour) samples. From their research it was found that, both solution and ageing heat treatments have no significant impact on the initial microstructure of wrought as well as-built LPBF 174 PH steel samples [231]. Few other researchers have studied the solution heat treated and aged microstructures of LPBF processed 17-4 PH steel and their investigation yielded similar results as one to that of Sun et al. [249,305]. Heat treatment of LPBF built duplex stainless steels was examined by [306]. Their observations reveal that recrystallization occurs in the temperatures ranging between 900 and $1200{ }^{\circ} \mathrm{C}$, whist, the maximum austenite fraction was achieved at the intermediate temperature $1000^{\circ} \mathrm{C}$ [306]. They also reported that missing nitrogen during the process but without quantifying it. Fig. 35a-d gives a clear picture of the comparison between as-built LPBF and heat treated (recrystallized micro- and nanostructure) of duplex stainless steels. LPBF processed duplex stainless steel displayed almost fully ferritic in as-built condition. On the contrary, conventionally processed steel displayed small austenite grains and chromium nitride precipitates along the grain boundaries (Fig. 35b \& c). Applying suitable post process treatment partially recovers the desired duplex austenite/martensite microstructure (see Fig. 35e \& f) [306]. Post process heat treatment at $500{ }^{\circ} \mathrm{C}$ dissolves the retained austenite, at the same time the dominant cellular microstructure disperses at the tempering temperatures above $600-700{ }^{\circ} \mathrm{C}$ in LPBF processed H-13 steel [192,307]. Also, Martensitic microstructure, similar to that of conventionally processed tool steels (without any trace of retained austenite) could be achieved by complete austenitization followed by quenching [192,308,309]. As-built LPBF samples that are tempered directly at low tempering temperatures did not exhibit any drop-in hardness value due to softening of the martensite, coincidentally which is also the characteristic of conventionally manufactured $\mathrm{H}-13$ tool steels [192,307]. Furthermore, secondary hardness peak transited to higher temperatures can be attributed to the large formation of carbon and carbide embedded in the retained austenite that is stable up to relatively high temperatures [192,307]. Secondary carbides could be formed only after this austenite is completely decomposed. As-built LPBF samples displayed a relatively lower tensile yield strength and UTS as compared to heat treated and conventionally manufactured steels. This could be attributed to the extreme brittleness that lead to premature failure without any deformation [68,262,264-266]. Again, tempering effect did not display any changes in as-built LPBF steels. Instead, the tensile value remained same as wrought and heat-treated tool steels [265]. However, the ductility obtained was much lower in as-built as well as aged samples, presumably due to the surviving process-induced defects [68,262,264-266]. Salman et al. investigated the changes induced by the various post process heat treatments on the mechanical properties of LPBF built 316L stainless steels under tensile loading conditions [275]. They revealed that the tensile strength of the 316L samples decreased with increasing annealing temperature as a result of the microstructural coarsening (see Fig. 36a). As-built LPBF samples exhibited an excellent combination of strength and ductility, along with the plastic deformation exceeding $50 \%$. This can be attributed to the complex cellular microstructure and subgrains along with the misorientation between the grains, cells, cell walls and subgrains; which would prevent the formation of higher local stresses, repel dislocation movements and defers premature fracture until ultimate tensile stress is reached [275]. Degradation of yield strength and ultimate tensile strength with increasing annealing\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-29} Fig. 33. Atom probe tomography (APT) images of LPBF maraging steels for the aged sample at $510{ }^{\circ} \mathrm{C}$ for $2 \mathrm{~h}$ (a) spherical precipitate enriched in Ti and $\mathrm{Ni}$, and (b) line scan showing the atomic concentration of each element [297].\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-30} Fig. 34. EBSD Grains maps of 316L stainless steel for: (a) as-built LPBF samples and specimens annealed at (b) $573 \mathrm{~K}$, (c) $873 \mathrm{~K}$, (d) $1273 \mathrm{~K}$, (e) $1373 \mathrm{~K}$ and (f) 1673 K [275].\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-30(1)} Fig. 35. TEM images of LPBF produced duplex stainless steel in the (a) as-built (d) heat treated condition, (b and c) electron backscatter diffraction (EBSD) inverse pole figure, and phase mappings for as-built condition, while (e and f) for heat treated condition [306].\\ \includegraphics[max width=\textwidth, center]{2024_04_13_3097074fadaf43f17ccag-31(1)} Fig. 36. (a) Tensile stress-strain curves of LPBF processed 316L stainless steel at different heat treatment conditions, and (b) effect of heat treatment on the yield and ultimate tensile strength of the different samples [275]. temperature is shown in Fig. 36b. This is due to microstructural variations that occur under various post process heat treatments. No preferred orientation of the grains was observed as annealing caused the growth of grains and cells while decreasing the dislocation network [275]. Conde et al. reported enormous improvement in hardness values $(\sim 60 \%)$, and bending strength ( $\sim 73 \%)$ similar to that of UTS while the substantial decrease in ductility can be noticed in age hardened martensite steel sample as compared to the as-built LPBF sample [298]. This remarkable hardness and tensile strength enhancement can be ascribed to the precipitation hardening and strengthening by the formation of fine precipitates of intermetallic compounds such as $\mathrm{Ni}_{3} \mathrm{Ti}$ in the martensite matrix. To overcome the loss of ductility, solution treatment combined with ageing or hot isostatic pressing can be employed which would result in better overall mechanical strength [298], along with considerable improvement in elongation (ductility) that lies within the standard ranges [69,210,274,297]. A stable microstructure was reported after conducting the stress relieving through post process heat treatments with temperature up to $650{ }^{\circ} \mathrm{C}[180]$. As a result of unchanged microstructure, there was no significant effect on crack propagation or the fatigue life of austenitic steels [180]. However, at higher annealing temperatures (above $800{ }^{\circ} \mathrm{C}$ in a furnace or HIP), resulted in partial recrystallization of austenitic steels. HIP induces partial recrystallization which is predominantly dual-mode isotropic. Consequently, this isotropic microstructure eventually lead to isotropic crack propagation under fatigue loading. Similarly, Saeidi et al. found that there were no changes in microstructure at annealing heat treatment at $800{ }^{\circ} \mathrm{C}$ [310]. In addition, they noticed unstable phase transformation in the form of sharp edged $\delta$-ferrite at annealing temperatures over $1150{ }^{\circ} \mathrm{C}$. This needle shaped $\delta$-ferrite was stable during slow cooling to room temperature as indicated in the equilibrium phase diagram. Combination of recrystallization and coarsening phenomena renounce tensile strength of the steels [310]. Applying appropriate post process solution treatment followed by subsequent ageing treatment helps in transforming the retained austenite to martensite. Higher YS can be obtained as compared to conventional processed 17-4 PH steels but with the compromised ductility [305,311,312]. Unfortunately, the fatigue life of LPBF processed 17-4 steel has minimal advantage of undergoing any post process heat treatments under all conditions. At high strain amplitudes, when fatigue life is short, the impact of defects is weaker than at longer fatigue life and low strain amplitudes. Since internal defects act as crack initiation site as well as dominating the mechanical behavior of the strong, heat-treated 17-4 steel sample displays a higher fatigue strength in the low-cycle regime as compared to the as-produced sample due to higher internal defects (Fig. 37) [249]. The fatigue life of as-produced LPBF H13 steels was significantly lower than that of conventionally produced \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-31} \end{center} Fig. 37. Experimental date and curves of Strain-life fatigue for LPBF 17-4 PH SS in different building conditions [249]. steels [265]. Becker et al. investigated the fatigue crack growth rates of peak aged material LPBF processed maraging steels, the authors found that fatigue crack growth rates were isotropic and equal to conventionally-processed maraging steels [261]. Isotropy is due to the weak texture orientation in these type of steels which is a consequence of the martensitic phase transformation [261]. This could be correlated to the LPBF process induced residual stress and higher surface roughness [264]. Surface machining had limited impact on improving the fatigue strength as it remained significantly lower than that of reference material (50\% failure probability at 107 cycles strength of $283 \mathrm{MPa}$ compared to $600 \mathrm{MPa}$ ) [264]. The improved fatigue properties were attained when the stress relief and austenite decomposition heat treatment at $600{ }^{\circ} \mathrm{C}$ were employed [265]. Sagar et al. studied the impact of different build directions and post process heat treatments (ageing and over ageing) on HCF and LCF of LPBF built 15-5 PH steels [313]. Based on their findings, aged specimens subjected to LCF performed better than as-built specimens in both vertical and horizontal build directions, whereas, the performance of the same samples subjected to HCF was poor compared to as-built LPBF parts. This was accredited to the ageing treatment; which resulted in precipitation strengthening of the matrix through copper-rich precipitation, however, this precipitation strengthening could lead the specimen more defect sensitive in HCF regime [313]. Also, overaged samples tend to be less sensitive to the defects than aged specimens in HCF\\ regime. This could be attributed to the over ageing heat treatment, which resulted in a more ductile LPBF built sample through microstructure grain coarsening of copper-rich precipitates in addition to the increased amount of retained austenite [313]. In general, the higher surface roughness caused by irregular shaped defects e.g. voids, partially-melted/un-melted etc. in LPBF processed samples are more sensitive to the fatigue life than its wrought counterparts. However, relatively lesser number of defects were reported for the vertically built LPBF specimens [313]. The removal of residual stresses via stress relief (SR) heat treatment $\left(5 \mathrm{~h}\right.$ at $470{ }^{\circ} \mathrm{C}$ ) does not necessarily improve the fatigue behavior. Instead, removing critical crack initiators by machining the surface of LPBF processed 316L steel significantly improved the fatigue performance. The superior fatigue performance was achieved through machined samples with and without SR heat treatment when compared to conventionally fabricated 316L steels. SR treatment coupled with the machining is recommended to obtain desired fatigue performance when cyclically loaded at high stresses. Overall, the post process treatments has very minimum effect on the fatigue performance of LPBF processed steels $[287,302]$. \section*{6. Summary and future trends} \subsection*{6.1. Steels in LPBF process} An outline of already published and ongoing research reveals that LPBF process influences the time related spatial variations of molten melt pool, thermal gradients induced vortex flow, solid-liquid interface velocity which leads to the formation of spatial microstructure and diverse mechanical properties within a specific geometry and the processing conditions for a wide variety of steels. The limiting factors for research and development of LPBF process of steels are mainly related to: Firstly, LPBF technique is not fully standardized and reliable for the industries to completely adapt as it is a maturing technology. However, some of the steels structural components have been approved to be used in the service. Secondly, it is difficult to regulate the metallurgical defects, microstructure evolution, and efficacy of mechanical performance driven studies have not been consistent. LPBF processing of steels research is largely concentrated on the limited type of steels. Current research trend is mostly focused on LPBF process optimization, studying the impact of post-process treatments on the microstructural changes and mechanical behavior (under static loading) of various alloying compositions, and commercially available stainless steels [314-317], tool steels [26,30,215], and maraging steels [297-299]. LPBF processing of alloy composition steels intensifies the solid-solution limit of the alloying elements in the molten melt pool and leads to the formation of unfavorable microstructure. Similarly, carbon bearing tool steels experiences relatively uncontrolled in-situ-tempering during LPBF processing. It is extremely important to accurately control the interplay between LPBF process parameters and the different type of steels, and possibly redesign/modify the new type of steels whose composition is best suited for the LPBF processing features, concurrently fulfill the desired functional properties. Some alloy combinations of stainless steels (i.e. 304 SS, 410 SS, 420 SS, 430 SS, Inox904L SS), maraging steels and other low alloy steels are not completely explored in LPBF process. Meanwhile, there is a continuous scope for expanding LPBF process of steels. A further investigation of other types of steels such as iron-based and nickel-based superalloys, single crystal alloys, and cobalt chromium alloy steels is much-needed. These materials are of high interest in wide range of aerospace and biomedical applications. As a general rule, steels and iron-based alloys are intended for structural applications. However, there is an increase in demand utilizing LPBF process for fabricating steels parts with functional properties. Fe-Al (low-density low-carbon ferritic) steels display outstanding strength to weight ratio. Owing to excellent resistance to corrosion, wear and oxidation LPBF built Fe-Al steels are considered for high temperature system's functional applications [164]. This include jet engines, turbine blades, heat exchangers piping, holding fixtures of a heat treatment furnaces and etc. The next category of functional LPBF processed steels are invar steels. The unique features on invar steels (FeNi or FeCoNi steels) posses over a wide temperature range and a negligible coefficient of thermal expansion, makes them perfect candidates in high temperature applications such as household appliances, electronic devices and aircraft controllers. Electrical steels are the other important category of LPBF functional components that are already in service. These electric steels are soft-magnetic materials that demand a specific crystallographic texture (i.e. Goss texture, $\{011\}\langle 100\rangle$ ) to achieve a low hysteresis losses [318], and have high permeability which means that the electrical current needed to produce magnetization should be as minimal as possible. This exclusive characteristics makes electric steels best suited for generators, alternators, amplifiers, transformers and iron electric motor applications. Furthermore, electrical steels are subjected to the series of rolling and annealing treatments to achieve Goss orientation; induces better power and higher permeability properties in the rolling direction. However, this (rolling and annealing treatments) requirement can be easily substituted by LPBF processing. On the other hand, by increasing $\mathrm{Si}$ content in $\mathrm{Fe}-\mathrm{Si}$-alloy which is again a soft magnetic polycrystalline metal alloy. LPBF processing of modified Fe-Si alloy steel exhibit a higher electrical resistivity and a lower hysteresis loss compared to the conventional electrical steels [319]. Modifying bulk Fe by adding Mn content enhances corrosion resistance along with displaying good cytocompatibility during LPBF process of Fe-Mn steels [320]. A basic research investigation is not sufficient, an advanced and thorough research is vital before using Fe-Mn scaffolds as functional components. Some of the other steels and iron-based functional components that are considered for use in laser based additive manufacturing includes amorphous, nanocrystalline, and magnetocaloric materials [321]. A number of LPBF processed steels and ironbased alloy components have already been in use for economic reasons, and further growth is expected in the near future. Development of LPBF process steel lattice structures with unique thermal, mechanical, electrical and acoustic features opened the door for cellular light weight structures applications in aerospace and biomedical industries. Cellular structures offers an exciting opportunity especially in design light weight applications due to their high strength to stiffness ratio provided by the porous structure. The light weight applications includes personal protective equipment (PPE), conformal cooling channels, thermal controllers, bone scaffolds, antimicrobial functionality possessing medical implants, sports equipment and etc. There have been good number of researchers attempted to study the microstructure relating to the mechanical properties of LPBF processed steel lattice structures. However, comprehensive analysis including FEM prediction of defects, structure and property relationship of LPBF steel lattice structures, and their overall performance capabilities is one area of the future scope to be explored in upcoming days. Functional graded steel lattice structures which display varying densities across the structure in contrast to regular lattice structures with uniform density throughout the structure is another exciting topic requires further research. Future developments of LPBF process of steels in the healthcare sector is focused on various features, such as biocompatibility, corrosion resistance, mechanical properties, printing properties, biomimetics, and degradation. The development of new type of steels as biomaterials with various compositions to achieve reprogrammed mechanical properties and functions would become a reliable method for building various organs and tissues with diverse mechanical requirements. For example, type 302 stainless steel was introduced solely for its application in orthopaedic surgery. Additionally, 316L type stainless steel is most commonly used in surgical procedures to replace biological tissue or to assist in stabilizing a biological component like bone tissue to aid the healing process. Stents, screws, plates, scaffolds are some of the commonly used functional components of LPBF produced 316L stainless\\ steels [322]. This type of stainless steel is the most corrosion resistant when it is in direct contact with biological fluid. A surgical implant must not be vulnerable to corrosion when positioned inside the human body to avoid any possible chance of infection. Hence, 316L stainless-steel implant is particularly effective when it is used in cold-worked condition, due to the non-existence of any inclusion in this material. Steel materials with inclusions also contain sulphur which is a key alloying element to encourage corrosion in steels. 316L stainless steel used in surgical implants contains approximately $\sim 17$ to $\sim 19 \%$ of chromium and $\sim 14 \%$ nickel. Corrosion resistance can also be achieved with the carbon, but only when the carbon is in a solid solution state. Chemical composition of stainless steel can be altered by adding chromium $(\sim 16 \%)$ to become corrosion resistant. Similarly, the addition of carbon and nickel $(\sim 7 \%)$ helps to stabilize the austenite in stainless steels. Adding Pd into TWIP steels can significantly improve corrosion resistance by forming decent intermetallic compounds. Similarly, adding silver into TWIP steels results in improved mechanical strength by establishing $\varepsilon$-martensite during deformation. Also, molybdenum is added to the steels implants that act as a protective layer sheltering the metal from exposure to an acidic environment. It is important to note that including ferrite element into stainless steel gives the metal a magnetic property, which is not ideal for surgical implants as it could obstruct the Magnetic Resonance Imaging (MRI) equipment. One of the most evident problems with using magnetic implants is their susceptibility to heating which could change the shape or structural position of this steel implants. The potential solution to resolve this magnetic property is by adding Mn in to steels, which enhances MRI compatibility by promoting austenitic growth. LPBF process of steels and iron-based alloys as biodegradable medical implants is the one of the exciting and novel research areas that can be further explored. Meanwhile, enhancing mechanical properties along with maintaining biocompatibility as well as good corrosion resistance of LPBF processed steels implants is another challenge which require detailed attention. Development and fabrication of complex biodegradable LPBF processed implants with porous architecture imparts poor surface quality. A suitable post-processing is necessary to overcome this issue. Addressing the link between post-processing and the poor surface finish of these implants is extremely paramount which is another important topic for the future examination. LPBF process of steels induces process related higher residual stresses, inevitable internal defects such as porosities, balling and thermal cracks that result in higher surface roughness. Metallurgical defects or any irregularities present at the surface pose an adverse effect on the final part geometry, and consequently, lead to poor surface quality and mechanical performance of the LPBF fabricated parts. Hence, unique and effective statistical approaches are needed that would consider of all these interdependencies of process related parameters while a minimum number of experiments are to be conducted. A considerable number of researchers have attempted to optimize the LPBF process by altering parameters and used many new approaches like employing multiple laser beams, laser de-focusing, laser re-melting, or adapting different scan strategies, substrate preheating and/or using a hybrid substrate to boost the efficiency of LPBF technology. However, there is still lack of detailed scientific understanding; how these inevitable metallurgical defects associated with LPBF process are going to behave when subjected to the dynamic loading conditions i.e. fatigue properties. A very limited research has been carried out focused on the influence of process parameters on the surface quality and fatigue properties of an LPBF processed product to optimize the given process in terms of time, cost, and properties of a product. Also, from the existing literature, most of the experiments that were carried out are based on recommended parameter settings provided by LPBF/AM machine vendors, which might induce uncertainty in the outcome of the process depending on the operator or expert knowledge of the vendors. Mathematical models based on FEM and regression analysis are developed to predict the process performance. However, these methods are usually not sufficient mainly because they lack the ability to extrapolate the given input data or information. Overall, the vast majority of available studies in the literature have investigated optimal process parameters for steels using simulation and/or experimental approach. The limitation of this approach is, if we change the experimental conditions (e.g., materials, process, system, or environment conditions), the resulting optimal process parameters settings may no longer be valid and applicable. New experiments are required to revalidate the samples. In addition, it is worth noting that very little attention has been paid towards the systematic study of surface texture characterization. The surface texture of LPBF processed steel parts impart anisotropic or sometimes isotropic. Surface finish is generally very sensitive to the mechanical performance especially under dynamic loading (fatigue testing) conditions. As in case of HCF limit, which is predominantly dependent on the surface texture of the LPBF fabricated steel products. \subsection*{6.2. LPBF technology} LPBF process of metals is gaining popularity while displaying significant growth along with taking greater steps in novel and advanced technologies to make it more competent, cost-effective. LPBF technology offers new opportunities in production paradigms, versatility of fabricating complex structures ranging from various scientific and technical innovation industrial applications to the retail and personal products. Due to higher freedom of design, LPBF process is very efficient in producing individual or customized products mainly in healthcare, aerospace, and automotive industrial applications. Medical implants, lattice type structures, high temperature resistant, and high strength combined with lower weight large components are among the examples. LPBF technology is also included in the group of sustainable and efficient production processes which helps in saving resources and protecting the environment. Sustainability studies carried out on LPBF process displays that significant reduction in material waste and fuel consumption are the two other principal benefits. Moreover, eco-design in LPBF process provides this opportunity where the environmental issues are considered in each design and fabrication phase, in accordance with, various eco-design tools, for example, life cycle analysis (LCA) method, could become handy to quantify the environmental impact associated with the LPBF built products. Despite several exciting opportunities and advantages offered by LPBF technique, there are still some obstacles that act against its rapid growth, such as size limitation, production time, limited number of materials, machine and production costs. LPBF technology is not completely matured yet and is still in its infancy stages. The technology needs to be further developed and standardized in the coming years, including the availability of the LPBF machines, and feedstock materials at a lower cost. Also, expanding the capabilities of the LPBF machines so that they are more autonomous, faster, more accurate and able to mass-produce design surface quality components in addition to the superior mechanical properties. In particular, innovative research and development steps are to be taken up by governments, academics, public and private sectors to design and improve the speed and accuracy of LPBF machines and raise the number of its applicable metal alloys and new methods while maintaining its economic viability. LPBF process needs to overcome some of these specific technical challenges before it becomes a reality for operational use in the industry. LPBF process lacks robustness, repeatability, reliability and process monitoring because two machines from the same manufacturer (or from different manufacturers) can yield different part quality. It is a requirement to understand, and efficiently be able to control the effects of machine-to-machine variability. It is also paramount to develop specifications and industrial standards for the processing of LPBF steel components. Development of integrated processes, through the advancement of technologies for monitoring and control of production processes needs to be prioritized. Find/propose alternatives using existing conventional qualification methods based on validated models,\\ probabilistic methods and standard parts. New standard and advanced non-destructive (NDT) capable of sensing LPBF processed critical defects with a high degree of accuracy shall also be developed. New design guidelines with innovative and customized structural features are needed to build optimized components in structural terms and weight, which is essential for the validation of virtual models based on physical models, to predict the characteristics of microstructure, surface texture, mechanical properties and corrosion properties. Existing design tools are not capable of capitalizing the full advantage from the process due to the compatibility of LPBF process specific features with present LPBF machines as the design rules are not fully checked and established. In addition, designers are forced to follow the design rules set for conventional manufacturing processes. Along with the product level testing, qualification process, methodologies for part verification and the product assurance of LPBF built parts need to be established. LPBF technology can be enhanced by introducing genetic algorithms, artificial intelligence, machine learning, and other similar computer automated systems, which are helpful to optimize process parameters. Also, these intelligent LPBF systems are beneficial to predict the geometry of the molten melt pool, microstructure, surface finish and mechanical properties by eradicating time-consuming, expensive trial and error methods to carry out the physical experiments. For example, introducing trained computer vision or unsupervised machine learning algorithm, operative on a small to medium size training data base of image patches will be beneficial to detect and classify the anomalies that occur during LPBF process. Similarly, developing new algorithms that are able to automatically alter the LPBF processed part geometry by compensating CAD model would be interesting. A significant update of the existing CAD software is required as it holds more important information than the normal STL files. Thereby, upgraded software could make reliable and meaningful alterations based on the example constraints already stored in the CAD file. This novel approach can be applied not only for the simple parts but also be useful for the complicated parts including internal surfaces, overhangs of curved and inclined LPBF processed components. Prediction, maintenance and regulation the dimensional tolerance variations of LPBF produced components are extremely important. Geometrical tolerances must kept to the lowest minimum tolerance range. The existing statistical analysis to develop linear models for the tolerance prediction is not sufficient. Adapting machine learning techniques for the prediction of dimensional features would be very valuable. LPBF product quality can be improved by minimizing or eliminating the defects. Metallurgical defects could be controlled by employing advanced numerical modelling and simulation methods, as well as realtime defects elimination by in-situ detection. In addition, coupling near infrared image (NIR) camera within LPBF systems would become advantageous to analyse each layer characteristics before and after it is selectively melted and deposited. The NIR data that provides information about the possible location of metallurgical defects and the surface dimensions can be utilized as an important tool to mitigate LPBF process related defects. Innovative hybrid LPBF process (combining LPBF technology with other AM process like cold spraying) would become more beneficial to fabricate high quality functional products with improved properties. It is essential for LPBF technology to advance in the direction of multiple materials fabrication of different products simultaneously. Developing intelligent materials should be concurrent with some advances in LPBF technique itself. With this advancement, smart materials can be produced into complex and multifunctional structures with higher precision and particular responsiveness. Altering the steel alloy by mixing with another metallic powder at various compositions on demand to create a flexible system is another interesting aspect in LPBF of steels. This flexible prototype system can work for in-situ alloying of various elements, in the meantime helpful to produce composite materials. There has been sincere effort to investigate the LPBF feasibility on fabricating multi-materials ( $\mathrm{Fe}$ and Al-12Si metallic powders) [323]. These LPBF composite materials applications are found in hybrid or translational junctional elements; nodes in heterologous frame space. However, a detailed work is needed in further to investigate the various processing strategies to enhance the part quality as well as component metallurgy [323]. \subsection*{6.3. Materials} The other key challenges faced by metal LPBF technology is the cost of materials and the array of materials. To tackle this, materials suppliers need to invest highly in the research and development of new LPBF materials to claim their stake in the global market. It is estimated that the cost of materials shall be driven down in the near future and a wider array of materials would be made widely available. In addition, the demand for industrial-grade materials such metals and its alloys also driving growth, particularly across the critical industrial sectors like aerospace, automotive, shipbuilding, transportation, and etc. Hence, there is a continuous rise in demand for materials like high-performance low weight functionally graded materials. New materials shall be developed to optimize the production process and the final quality of LPBF processed components. Physical (morphology, flowability, particle size distribution, humidity, moisture content), and chemical properties (impurities level) of LPBF powders (including steel powders) are not fully defined to achieve functional parts quality. Currently, there is no powder handling or recycling specifications to ensure traceability and avoid contamination of feedstock materials, which needs to established at the earliest convenience. \subsection*{6.4. Post processing} The post-processing in LPBF technology is cited as "hidden dirty secret" because it is a necessary part of LPBF process workflow. Post processing has naturally been manual and highly labor intensive. Simplifying the entire LPBF process starting from designing a 3D computer model that goes for a print, to the finished part, which is ready to use has therefore been key imperative for the industries. An imperative that is been driven by complete automation. The ability to automate the post-processing stage comes with several benefits, such as saving costs in both manual labor force and machine running cost as well as being able to reduce the overall production times. For example, automated postprocessing systems could be comprised of automated cleaning, polishing, injection molding to achieve smoother surface finish. The effect of existing post process heat treatments (stress relieving, solution/ageing, vacuum furnaces, hot-isostatic-pressing), surface treatments (mechanical and electrochemical polishing, abrasive flow polishing), and cleaning procedures (jet blasting, sand blasting) applied to LPBF processed components on the final properties is still unclear, which needs to be addressed. It is utmost important for the researchers and manufacturers to shift their focus on design and development of specific application orientated-optimized part geometries, compositions, and functionality of AM/LPBF steels or other metal components. Overall, more and more researchers and experts in AM fraternity (including LPBF) are turning their interest towards fabricating bulk structures using customized large-format additive manufacturing machines, hoping to create large structures all at once, to avoid assembly or post processing methods that are expensive and time-consuming. All in all, the versatility at which LPBF and/or AM offers as non-conventional approach providing new opportunities for mass customization of complex parts by saving time, costs, and establishing process efficiency; which is playing a major role in branding AM is leading the subsequent main industrial revolution 4.0. Industrial manufacturing companies tend to enter the highly decentralized industrial revolution. However, AM experts presume that neither LPBF technology nor any other technique in AM process is going to substitute traditional manufacturing processes completely. In the coming years, a greater correlation between the machines, materials and software is expected as manufacturers demand continuous \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_3097074fadaf43f17ccag-35} \end{center} Fig. 38. Schematic illustration to show the correlation between the machines, materials, and software work together to ensure an end-to-end seamless LPBF process workflow. workflows and systems that collaborates harmoniously (see Fig. 38). All of these elements would consequently need to come together, work hand-in-hand to ensure an end-to-end seamless LPBF process workflow. \section*{Declaration of competing interest} The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. \section*{Acknowledgements} Dr. Shan Lou would also like to thank the EPSRC (EP/S000453/1) for funding this work. 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Manuf Lett 2017;11:8-11. \begin{itemize} \item \end{itemize} \end{document} % This LaTeX document needs to be compiled with XeLaTeX. \documentclass[10pt]{article} \usepackage[utf8]{inputenc} \usepackage{ucharclasses} \usepackage{amsmath} \usepackage{amsfonts} \usepackage{amssymb} \usepackage[version=4]{mhchem} \usepackage{stmaryrd} \usepackage{hyperref} \hypersetup{colorlinks=true, linkcolor=blue, filecolor=magenta, urlcolor=cyan,} \urlstyle{same} \usepackage{graphicx} \usepackage[export]{adjustbox} \graphicspath{ {./images/} } \usepackage{polyglossia} \usepackage{fontspec} \setmainlanguage{english} \setotherlanguages{romanian, hindi} \newfontfamily\hindifont{Noto Serif Devanagari} \newfontfamily\lgcfont{CMU Serif} \setDefaultTransitions{\lgcfont}{} \setTransitionsFor{Hindi}{\hindifont}{\lgcfont} \title{Review } \author{Alberico Talignani ${ }^{\mathrm{a}, 1}$, Raiyan Seede ${ }^{\mathrm{b}, 1}$, Austin Whitt ${ }^{\mathrm{b}}$, Shiqi Zheng ${ }^{\mathrm{a}}$, Jianchao $\mathrm{Ye}^{\mathrm{c}}$,\\ Ibrahim Karaman $^{\text {b,*, }}$, Michael M. Kirka ${ }^{\text {d,* }}$, Yutai Katoh ${ }^{\text {e, }}$, Y. Morris Wang ${ }^{\text {a,* }}$\\ a Department of Materials Science and Engineering, University of California, Los Angeles, CA 90049, USA\\ ${ }^{\mathrm{b}}$ Department of Materials Science and Engineering, Texas A\&M University, College Station, TX 77843, USA\\ ' Materials Science Division, Lawrence Livermore National Laboratory, Livermore, CA 94005, USA\\ ${ }^{\mathrm{d}}$ Manufacturing Science Division, Oak Ridge National Laboratory, Oak Ridge, TN 37830, USA\\ ${ }^{\mathrm{e}}$ Materials Science \& Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37830, USA} \date{} %New command to display footnote whose markers will always be hidden \let\svthefootnote\thefootnote \newcommand\blfootnotetext[1]{% \let\thefootnote\relax\footnote{#1}% \addtocounter{footnote}{-1}% \let\thefootnote\svthefootnote% } %Overriding the \footnotetext command to hide the marker if its value is `0` \let\svfootnotetext\footnotetext \renewcommand\footnotetext[2][?]{% \if\relax#1\relax% \ifnum\value{footnote}=0\blfootnotetext{#2}\else\svfootnotetext{#2}\fi% \else% \if?#1\ifnum\value{footnote}=0\blfootnotetext{#2}\else\svfootnotetext{#2}\fi% \else\svfootnotetext[#1]{#2}\fi% \fi } \begin{document} \maketitle \section*{A review on additive manufacturing of refractory tungsten and tungsten alloys} \section*{A R T I C L E I N F O} \section*{Keywords:} Tungsten Tungsten alloys Additive manufacturing Laser powder-bed-fusion Laser directed-energy-deposition Electron beam powder-bed-fusion \begin{abstract} A B S T R A C T We review the progress of additive manufacturing effort on refractory metal tungsten and tungsten alloys. These materials are excellent candidates for a variety of high temperature applications but extremely challenging to fabricate via additive manufacturing due to a series of existing issues during the manufacturing. We outline these issues and discuss the current understanding and progress to tackle them. Laser powder-bed-fusion, laser directed-energy-deposition, and electron beam powder-bed-fusion are three common techniques that have been applied to additively manufacture pure tungsten. This overview discusses current observations and understanding on the issues associated with each of these techniques. We identify future research opportunities in additive manufacturing of refractory metals. \end{abstract} \section*{1. Introduction} Due to their high density, excellent thermal conductivity and high temperature capabilities, high strength and hardness, and minimal sputtering yield and hydrogen interactions, refractory metal tungsten and tungsten alloys have a broad range of potential applications, including plasma facing components for fusion reactors [1,2], fusion targets [3], armor penetrators [4], and nuclear space power and propulsion [5]. For many of these applications, additive manufacturing (AM) offers unique geometrical design freedom and rapid prototyping capability, which is unparalleled by conventional manufacturing techniques. Moreover, AM offers additional potentials to fabricate functionally graded transitions from tungsten to various dissimilar materials. Due to its extremely high melting temperature (for pure tungsten, the melting temperature $\mathrm{T}_{m}=3422^{\circ} \mathrm{C}$ ) and brittle nature, tungsten is notoriously difficult to fabricate via either laser- or electron-beam-based AM techniques. Nevertheless, encouraging progress has been made in the past few years in AM tungsten and its alloys. Given the rising importance of refractory metals in various high temperature applications for harsh environments, this article aims at providing a timely overview of recent development in AM of refractory metals, in particular, tungsten and tungsten alloys. To date, laser powder-bed-fusion (L-PBF) [sometimes also known as selective laser melting (SLM)], laser directed-energy-deposition (LDED), and electron beam powder-bed-fusion (EB-PBF) [or electron beam melting (EBM)] are the most common techniques to fabricate tungsten materials. The first two utilize laser energy to melt tungsten powder, and the latter electron energy. This review is arranged according to the materials made by the above three AM techniques, each of which has its own unique set of promises, challenges, and opportunities. As processing conditions determine the manufacturing defects, microstructure, and subsequent mechanical properties, Table 1 summarizes some key features and limitations of each AM technique, which help readers to better understand the microstructural origins of each type of materials and subsequent challenges involved in each approach. Notably, L-PBF offers substantially higher cooling rate and stronger temperature gradient compared to other two techniques, and thus may influence the residual stresses and cracking behavior. This review focuses on pure tungsten, as it is arguably one of the most challenging materials for AM. We contend, however, that many challenges and issues encountered in tungsten \footnotetext{\begin{itemize} \item Corresponding authors. \end{itemize} E-mail addresses: \href{mailto:ikaraman@tamu.edu}{ikaraman@tamu.edu} (I. Karaman), \href{mailto:kirkamm@ornl.gov}{kirkamm@ornl.gov} (M.M. Kirka), \href{mailto:ymwang@ucla.edu}{ymwang@ucla.edu} (Y.M. Wang). 1 These authors contributed equally to this work. } Table 1 A summary of some key processing features [7] of three commonly used AM techniques for tungsten and tungsten alloys; i.e., laser powder-bed-fusion (L-PBF), laser directed-energy-deposition (L-DED), and electron beam powder-bed-fusion (EB-PBF). \begin{center} \begin{tabular}{llll} \hline & L-PBF & L-DED & EB-PBF \\ \hline Source power $(\mathrm{W})$ & $10^{2}-10^{3}$ & $10^{2}-10^{4}$ & $10^{2}-10^{3}$ \\ Beam size $(\mu \mathrm{m})$ & $30-200$ & $10^{2}-10^{3}$ & $10^{2}-10^{3}$ \\ Scanning speed $(\mathrm{mm} / \mathrm{s})$ & $10^{1}-10^{3}$ & $10-10^{2}$ & $10^{1}-10^{3}$ \\ Cooling rate $(\mathrm{K} / \mathrm{s})$ & $10^{5}-10^{7}$ & $10^{2}-10^{5}$ & $10^{3}-10^{4}$ \\ Temperature gradient (K/ & $10^{6}-10^{7}$ & $10^{4}-10^{6}$ & - \\ $\quad$ m) & & & \\ Environment & Argon, & Argon & Vacuum, trace \\ & nitrogen & & helium \\ Material waste & High & Minimal & High \\ Pre-sintering & No & No & Yes \\ Spattering & Yes & No & No \\ \hline \end{tabular} \end{center} manufacturing are likely applicable to a general class of refractory metals including high entropy alloys [6], which are prone to cracking during manufacturing. The review ends with our recommendations on the future opportunities for AM refractory metals, in the hope to spur future research in these interesting materials. \section*{2. Laser powder-bed-fusion} \subsection*{2.1. Method} L-PBF is a well-known additive manufacturing technique for metals and alloys and sometimes ceramics. During this process, the powder is deposited layer-by-layer ( $\sim 20-150 \mu \mathrm{m}$ thick), and a laser beam (either continuous or pulsed wave) is applied to selectively melt the desired region. Some critical parameters that influence the build quality of materials include build layer thickness, laser power, scan speed, and hatch spacing. These parameters influence the volumetric energy density of processing conditions. Laser absorptivity is another important variable that determines the percentage of laser energy coupled into the powder layer. Notably, L-PBF is typically performed in an argon environment with oxygen levels ranging from tens to hundreds of ppm. Compared to L-DED and EB-PBF, the beam size of L-PBF is appreciably smaller, leading to a higher cooling rate and stronger temperature gradient, Table 1. Powder spattering and denudation phenomenon are common features of L-PBF processes, which cause processing defects/ pores that are difficult to eliminate $[8,9]$. \subsection*{2.2. Cracking} The ability of powerful lasers to melt essentially any types of metals makes L-PBF a natural choice to fabricate tungsten. However, cracking has been the biggest challenge in L-PBF W. None of the available literature has reported crack-free samples except when a femtosecond laser source was used [10]. As such, understanding the cracking mechanisms during L-PBF processes has been a central focus of recent studies. Generally speaking, two types of cracks have been observed in L-PBF W: longitudinal (with the crack direction parallel to the laser scanning direction) and branched or transverse cracks (with the crack inclined to the laser scan direction) [11-16]. The crack nucleation and propagation in L-PBF W are considered to be associated with the high ductile-to-brittle transition temperature (DBTT) $\left(\sim 200-400^{\circ} \mathrm{C}\right)$ of tungsten. The direct evidence supporting the above proposition is the appreciable time delay between the \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-02(4)} \end{center} (c) Frame $150, \mathrm{t}=3.00 \mathrm{~ms}$ \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-02(3)} \end{center} Frame $329, \mathrm{t}=6.58 \mathrm{~ms}$ \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-02(1)} \end{center} Frame $459, \mathrm{t}=9.18 \mathrm{~ms}$ \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-02(2)} \end{center} Frame $638, \mathrm{t}=12.76 \mathrm{~ms}$ \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-02} \end{center} Fig. 1. Scanning electron micrographs showing cracking behavior of tungsten bare plate during single-track experiments. (a) The laser power was set at $\mathrm{P}=300 \mathrm{~W}$ and speed $v=300 \mathrm{~mm} / \mathrm{s}$. Melt pool is marked with a dashed line. A branched crack is shown by the black arrow. (b) $\mathrm{P}=450 \mathrm{~W}$ and $\mathrm{v}=100 \mathrm{~mm} / \mathrm{s}$. The melt pool is well defined by the grains. Branched cracks are observed [12]. (c) Images taken from in-situ high speed camera experiments with $\mathrm{P}=300 \mathrm{~W}$ and $\mathrm{v}=100 \mathrm{~mm} / \mathrm{s}$. Cracks are highlighted inside red rectangles. The amount of time elapsed is shown in each image [15].\\ solidification and the appearance of cracks, as recorded by an in-situ high speed camera during single track experiments, Fig. 1 [12,15]. Another important observation is that cracks tend to propagate along high angle grain boundaries (HAGBs) [11-16], examples of which are shown in Fig. 2 [11]. This behavior can be attributed to the sensitivity of grain boundaries (GBs) in tungsten to impurities. Oxygen is a known impurity in tungsten powder and has been reported ranging from 30 to $370 \mathrm{ppm}[12,17]$. Several groups attributed the formation of cracks to aggregation of tungsten oxides during solidification [11,14,18,19]. However, cracks are not fully eliminated even when the oxygen level is very low [12], suggesting that impurities might not be the only factor influencing the cracking behavior. A systematic study of oxygen or other impurities such as hydrogen effects on the brittleness of L-PBF W remains missing. Tungsten powder size, shape, and distribution have also been reported to influence the cracking behavior of L-PBF materials [20]. However, a systematic study is needed in order to fully clarify this phenomenon. To further understand the role played by residual stresses in cracking of L-PBF W, electron backscatter diffraction (EBSD) studies have been performed on cross-sections of printed tungsten samples [13, 16, 17, 19, 21-25]. A correlation between the density of HAGBs and cracks was observed. Although HAGBs are more prone to cracking than low-angle GBs (LAGBs), the formed cracks help to relieve intergranular stresses [13,19,23,24]. As shown in Fig. 3 [19], most cracks are observed along HAGBs, and perhaps even more importantly, the regions right next to the cracks have lower Kernel Average Misorientation (KAM) values compared to regions in which cracks are absent - evidence that most of the plastic deformation experienced by the material is concentrated near HAGBs [19]. Similar observations were made by another group, Fig. 4 [24], where cracks tend to appear near HAGBs (instead of LAGBs). \subsection*{2.3. Strategies to mitigate cracks} \subsection*{2.3.1. Alloying} To suppress cracks in L-PBF W, incorporation of rare earth or other elements into pure tungsten has been explored. Researchers mixed pure tungsten powder with $1 \mathrm{wt} \%, 5 \mathrm{wt} \%$, and $10 \mathrm{wt} \%$ of Ta powder [23]. As shown in Fig. 5 [23], addition of $5 \mathrm{wt} \%$ Ta significantly decreased the grain size. However, no further grain refinement was observed with $10 \mathrm{wt} \% \mathrm{Ta}$. The refinement of grain size appears to reduce the cracks. The same approach was reported by another group [26,27], where alloying with Ta was found to reduce the average crack length per unit area by $30.7 \%$. One possible reason is that Ta has higher electron affinity to oxygen than $\mathrm{W}$ (the formation Gibbs free energy of $\mathrm{Ta}_{2} \mathrm{O}_{5}$ is $-1904 \mathrm{~kJ} / \mathrm{mol} \mathrm{vs}-761.5 \mathrm{~kJ} / \mathrm{mol}$ for $\mathrm{WO}_{3}$ ). Thus, Ta has the tendency of attracting oxygen and mitigating the impurity segregation along GBs. In a similar approach, $5 \mathrm{wt} \%$ of $\mathrm{Nb}$ was added to tungsten powder during L-PBF (the formation Gibbs free energies of $\mathrm{Nb}_{2} \mathrm{O}_{5}, \mathrm{NbO}_{2}$, and $\mathrm{NbO}$ are $-921 \mathrm{~kJ} / \mathrm{mol},-771 \mathrm{~kJ} / \mathrm{mol},-416 \mathrm{~kJ} / \mathrm{mol}$, respectively), and it was also found effective in suppressing cracks [24]. Although these alloying approaches achieved a certain degree of success, the underlying mechanisms have been poorly understood and require further investigations. In a different study, $0.5 \mathrm{wt} \%$ of $\mathrm{ZrC}$ was added to tungsten. Grain refinement was observed and a reduction of crack density as high as $88.7 \%$ was achieved [25]. The beneficial effect of yttrium oxides $\left(\mathrm{Y}_{2} \mathrm{O}_{3}\right)$ has also been studied. In this case, no significant change of average grain size was noticed, crack reduction was still observed and attributed to tungsten grain shape changes [19]. A comparison between micro- and nano-sized $\mathrm{Y}_{2} \mathrm{O}_{3}$ was also carried out. The addition of nano- $\mathrm{Y}_{2} \mathrm{O}_{3}$ was found to reduce cracks due to the formation of a large fraction of LAGBs, whereas a reduction in hardness was seen with the addition of micro-sized $\mathrm{Y}_{2} \mathrm{O}_{3}$ [19]. In contrast, a separate single track experiment [28] found the addition of $\mathrm{Al}_{2} \mathrm{O}_{3}, \mathrm{Y}_{2} \mathrm{O}_{3}$, and $\mathrm{ZrO}_{2}$ to have no influence on suppressing cracks. The above results suggest that there is no consensus in the scientific community in terms of the choice of alloying element and other additives. \subsection*{2.3.2. Remelting, scanning strategies, and substrate heating} Other reported strategies to suppress cracks in L-PBF W include\\ \includegraphics[max width=\textwidth, center]{2024_04_13_b812db9e8fc840c3b338g-03} Fig. 2. Electron backscatter diffraction (EBSD) images of two sides of a L-PBF W sample showing surface cracks. (a, c) As-printed W cross sections on different axes. Longitudinal, branched, and parallel build direction (BD) cracks are visible (red and blue arrows). (b, d) EBSD inverse pole figure (IPF) maps of (a, c). Scan tracks and scanning directions (SD) are visible, and 'ladder-shaped' grains and cracks along the grain boundaries are seen. The effect of rotation by 67 plus remelting between each layer is shown [11].\\ \includegraphics[max width=\textwidth, center]{2024_04_13_b812db9e8fc840c3b338g-04} Fig. 3. Scanning electron micrographs (SEM), electron backscatter diffraction (EBSD), and Kernel average misorientation (KAM) data from the surface of a L-PBF W sample. Cracks are observed along high angle grain boundaries in L-PBF W. (a) SEM image of the L-PBF W sample. (b) EBSD inverse pole figure (IPF) map. (c) Image quality map. (d) KAM map [19]. scanning strategy adjustment, remelting, and substrate heating. These processes aim to alleviate or minimize the residual stresses formed in tungsten during printing, which are due to the high temperature gradients near the melt pool. Remelting plus rotating strategy has been studied in a series of tungsten builds [11]. It was found that rotating 67 between each layer randomized the grain orientation and shape, thus reducing the so-called "ladder-shaped" structure formed by grains without rotation (see Fig. 2). This process hindered the formation of cracks since the ladder-shaped grains provide crack-formation sites. While only one example is given, almost all works in the literature adopted the scan vector rotation strategy in between build layers most used 67 so as to minimize scan alignment in the same orientation, while others opted for either $45^{\circ}$ or $90^{\circ}[23,29]$. Remelting refers to the process of scanning a track more than once before recoating the sample with fresh powder. In conjunction with rotation strategies, remelting eliminated the columnar grains and helped to suppress longitudinal cracks [11]. As a result, the remelt sample was found to have smaller grain sizes and the average surface roughness was reduced [30]. Nonetheless, the combination of scan rotation and remelting was not sufficient to fully suppress cracks [11,30,31]. In addition, it was suggested that remelting may impact the density of the sample compared to a non-remelted reference [30]. This phenomenon is not well understood. Substrate heating (up to $1000{ }^{\circ} \mathrm{C}$ ) is another strategy to suppress cracks in L-PBF W [29]. The purpose of substrate heating is two-fold: to reduce the temperature gradient (and thus residual stresses during L-PBF) and to bypass the DBTT of tungsten. Given that the DBTT is above room temperature $\left(200-400^{\circ} \mathrm{C}\right)$, embrittlement is likely to occur during solidification and cooling $[12,31,32]$. Theoretically, if the substrate is preheated above the DBTT, the screw dislocations in tungsten will have enough mobility and accommodate the plastic strain induced by the temperature gradient during melting and subsequent cooling [29]. Another important aspect is that maintaining an impurity-free environment during printing is crucial, as the DBTT of tungsten can be theoretically shifted by $200{ }^{\circ} \mathrm{C}$ between $10 \mathrm{ppm}$ and $50 \mathrm{ppm}$ of oxygen content [12]. Despite tremendous effort [29], substrate heating between 80 and $1000^{\circ} \mathrm{C}$ has not been enough to fully eliminate cracking. A detailed study of how to optimize the preheating conditions (e.g., temperature and cooling rate) is needed in the future in order to fully understand their influences on cracking behavior.\\ \includegraphics[max width=\textwidth, center]{2024_04_13_b812db9e8fc840c3b338g-05(1)} Fig. 4. Electron microscopy was used to characterize surface cracks in Nb-alloyed W. (a) Scanning electron microscope (SEM) image of W-5 wt\% Nb alloy. Cracks are present and pointed by arrows. (b) Electron backscatter diffraction (EBSD) inverse pole figure (IPF) map of image (a), cracks are again indicated with arrows. (c) Inverse pole figure (IPF) grain boundary (GB) distribution map. Red lines represent high angle GBs. Cracks appear only on red lines. (d) GB distribution plot. Those images are adopted from [24].\\ \includegraphics[max width=\textwidth, center]{2024_04_13_b812db9e8fc840c3b338g-05} Fig. 5. Electron backscatter diffraction (EBSD) images of three Ta-alloyed L-PBF W samples. From left to right, normal direction inverse pole figure (IPF) maps of cross-sections from pure W, W-5 wt\%Ta, and W-10 wt\% are shown. Grains become more refined between (a) and (b), while a smaller change is noticed in (c) [23]. \subsection*{2.4. Processing parameter windows} Aside from cracks, achieving a high relative density is crucial in LPBF W. As summarized in Table $2[7,9,10,12-19,22-26,27]$, the relative density of tungsten obtained in the literature ranges from $\sim 80 \%$ to $\sim 99 \%$ [14]. To obtain good processing parameters for high density samples, laser power versus scan speed graphs have been used $[21,32$, 33]. In such a graph, different processing zones are marked according to a combination of laser power and speed, such as irregular crack region, regular crack region, balling region, warped region, and dense region $[21,32,33]$. Similar graphs have also been adopted to study cracking behavior in single-track experiments [12,32,33]. However, it has been demonstrated in the literature that such graphs are only useful for a specific type of L-PBF machine and with a fixed build layer thickness. An energy input diagram or normalized enthalpy diagram would be more useful to identify processing windows for tungsten and offer better Table 2 A summary of laser processing parameters for L-PBF W and resultant mechanical properties in the literature. \begin{center} \begin{tabular}{|c|c|c|c|c|c|c|c|c|c|c|c|} \hline Ref. & Machine & \begin{tabular}{l} Laser \\ Power \\ (W) \\ \end{tabular} & \begin{tabular}{l} Energy \\ Density $^{1}$ \\ $\left(\mathrm{~J} / \mathbf{m m}^{3}\right)$ \\ \end{tabular} & \begin{tabular}{l} Beam/ \\ Hatch \\ $(\mu m)$ \\ \end{tabular} & \begin{tabular}{l} Layer \\ Thick. \\ $(\mu m)$ \\ \end{tabular} & \begin{tabular}{l} Max Density \\ $(\%)$ \\ \end{tabular} & \begin{tabular}{l} Hardness $^{4}$ \\ $(\mathrm{GPa})$ \\ \end{tabular} & \begin{tabular}{l} Preheating \\ $\left({ }^{\circ} \mathrm{C}\right)$ \\ \end{tabular} & \begin{tabular}{l} Powder \\ Size $(\mu m)$ \\ \end{tabular} & Alloying & Cracking \\ \hline [31] & \begin{tabular}{l} Renishaw \\ AM250 \\ \end{tabular} & \begin{tabular}{l} 200 \\ (pulsed) \\ \end{tabular} & - & $75 / 90$ & - & 82.90 & - & No & 19.4 & No & Yes \\ \hline [23] & \begin{tabular}{l} Custom \\ Built, DMP \\ 320 \\ \end{tabular} & 300 & $150-900$ & 90 & - & $\sim 81-98.7$ & - & $0-400$ & - & \begin{tabular}{l} No and \\ Yes (Ta) \\ \end{tabular} & \begin{tabular}{l} Yes. Cracks are less \\ evident with alloying. \\ Both transverse and \\ longitudinal cracks are \\ observed. \\ \end{tabular} \\ \hline $[21]$ & \begin{tabular}{l} Renishaw \\ AM400 \\ \end{tabular} & \begin{tabular}{l} $150-400$ \\ (pulsed) \\ \end{tabular} & $88-1185$ & \begin{tabular}{l} $75 /$ \\ $75-150$ \\ \end{tabular} & 30 & $\sim 80-96$ & $\sim 2-3.79$ & No & 28 & No & Yes \\ \hline $[11]$ & \begin{tabular}{l} Renishaw \\ AM400 \\ \end{tabular} & \begin{tabular}{l} 400 \\ (pulsed) \\ \end{tabular} & 474 & $75 / 100$ & 30 & $92.5-96.5$ & - & No & 28 & No & \begin{tabular}{l} Yes, cracks are longer \\ than $1 \mathrm{~mm}$ (along \\ HAGBs). \\ Transverse and \\ longitudinal cracks are \\ observed. \\ \end{tabular} \\ \hline [13] & \begin{tabular}{l} Renishaw \\ 125 \\ \end{tabular} & \begin{tabular}{l} 200 \\ (pulsed) \\ \end{tabular} & $641-930$ & \begin{tabular}{l} $43 /$ \\ $115-155$ \\ \end{tabular} & 50 & $94-98$ & - & No & $\sim 47$ & No & Yes \\ \hline [19] & \begin{tabular}{l} Renishaw \\ AM400 \\ \end{tabular} & \begin{tabular}{l} 250 \\ (pulsed) \\ \end{tabular} & 544-1587 & \begin{tabular}{l} $70 /$ \\ $50-100$ \\ \end{tabular} & $20-35$ & $94.5-98.30$ & $\sim 3.63-4.21$ & 180 & \begin{tabular}{l} $15-45 \mathrm{~W}$ \\ $15-53$ \\ $\mathrm{Y}_{2} \mathrm{O}_{3}$ \\ \end{tabular} & \begin{tabular}{l} No and \\ Yes \\ $\left(\mathrm{Y}_{2} \mathrm{O}_{3}\right)$ \\ \end{tabular} & \begin{tabular}{l} Yes, hundreds of \\ microns and in all \\ directions. \\ Oxides reduce \\ cracking. \\ \end{tabular} \\ \hline $[30]$ & EOS M290 & $150-350$ & $\sim 94-875$ & 100 & 20 & 98.40 & $\sim 4.02-4.47$ & 180 & 15.8 & No & \begin{tabular}{l} Yes, fewer cracks in \\ the bulk. \\ Remelting improves \\ cracking. \\ \end{tabular} \\ \hline [17] & EOS M290 & $200-370$ & $250-1850$ & 50 & 20 & $97.72-98.50$ & $\sim 4.36-4.58$ & 50 & 16.24 & No & Yes \\ \hline $[22]$ & \begin{tabular}{l} SLM® \\ Solution \\ $125 \mathrm{HL}$ \\ \end{tabular} & $200-400$ & 198-905 & 70/105 & 30 & 98.51 & - & 200 & $5-25$ & No & Yes \\ \hline [14] & \begin{tabular}{l} EOSM100 \\ DMLS \\ \end{tabular} & $100-170$ & 125-1062 & \begin{tabular}{l} $40 /$ \\ $40-70$ \\ \end{tabular} & 20 & 99.61 & $\sim 4.12$ & 80 & $10-25$ & No & \begin{tabular}{l} Yes. Longitudinal: \\ straight, $30-100 \mu \mathrm{m}$. \\ Transverse: shorter \\ and S shaped along \\ GBs. \\ \end{tabular} \\ \hline [18] & \begin{tabular}{l} Custom \\ Built \\ \end{tabular} & $200-350$ & $500-1167$ & 50 & 20 & $87.8-89.4$ & $\sim 4.65$ & 200 & 14.41 & No & Yes \\ \hline $[24]$ & EOS M280 & $250-370$ & - & $70-110$ & 30 & $93.3-98.0$ & $6.69-10.31$ & 200 & \begin{tabular}{l} $5-25 \mathrm{~W}$ \\ $1-10 \mathrm{Nb}$ \\ \end{tabular} & Yes, Nb & \begin{tabular}{l} Yes, along HAGBs. Nb \\ alloying partially \\ suppresses \\ cracks. \\ \end{tabular} \\ \hline [29] & \begin{tabular}{l} Aconity 3D \\ GmbH \\ \end{tabular} & $375-400$ & $196-446$ & $100 / 80$ & 40 & $\sim 94.7-98.5$ & - & $600-1000$ & $15-45$ & No & \begin{tabular}{l} Yes (reduced at \\ $1000^{\circ} \mathrm{C}$ ). \\ \end{tabular} \\ \hline $[26]$ & SLM 280 & 400 & - & 100 & - & - & - & - & \begin{tabular}{l} $32 \mathrm{~W}$, \\ $18 \mathrm{Ta}$ \\ \end{tabular} & Yes, Ta & \begin{tabular}{l} Crack density reduced \\ by alloying. \\ \end{tabular} \\ \hline $[27]$ & \begin{tabular}{l} Renishaw \\ AM400 \\ \end{tabular} & \begin{tabular}{l} 400 \\ (pulsed) \\ \end{tabular} & - & 100 & - & - & - & - & - & Yes, Ta & \begin{tabular}{l} Yes, less with Ta \\ (along GBs). \\ \end{tabular} \\ \hline $[16]$ & SLM 125HL & 400 & 238-1667 & \begin{tabular}{l} $80 /$ \\ $100-120$ \\ \end{tabular} & 30 & $\sim 90-97$ & - & $\mathrm{HIP}^{3}$ & 32 & No & Yes, reduced with HIP. \\ \hline \end{tabular} \end{center} ${ }^{1}$ Energy Density $=\frac{P}{l \times v \times D_{\text {beam }}}$, where $\mathrm{P}$ is power, 1 is layer thickness, $\mathrm{D}$ is beam diameter and $\mathrm{v}$ is laser speed. Speed is calculated using point distance and exposure time for pseudo-pulsed laser machines. 2 These values are not calculated according to the equation in note 1 but are reported as found in the respective papers. ${ }^{3}$ HIP: Hot Isotactic Pressing. Note that HIP is not preheating. ${ }^{4}$ A compressive strength in the range of $900-1523$ MPa has been reported for L-PBF W [17,18,22,30,32]. machine-to-machine variation comparison [34-38]. Regretfully, such laser processing maps for pure tungsten do not exist yet. \subsection*{2.5. Mechanical properties} Due to the poor sample quality, few studies have been conducted on documenting the mechanical properties of L-PBF W. The available data are limited to hardness and compressive mechanical properties, Table 2. Depending on the processing parameters, a hardness value of 3.63-4.47 GPa [14, 17-19, 30, 32] was reported for L-PBF W, which was superior to samples made by conventional powder metallurgy and spark plasma sintering (the hardness values range between 3.14-3.92 GPa; i.e., $320-400 \mathrm{HV}$ [17]). Alloying with $5 \mathrm{wt} \% \mathrm{Nb}$ was found to elevate the hardness from $6.69 \mathrm{GPa}$ to $8.01 \mathrm{GPa}$ [24]. Similarly, the addition of nano-yttrium oxides increased the hardness to $\sim 4.51 \mathrm{GPa}(\sim 460 \mathrm{HV})$. The effect was attributed to dispersion strengthening. In the same study, it was also shown that introducing micro-yttrium oxides (instead of nano-sized oxides) resulted in lower hardness than that of conventionally manufactured tungsten. The reduction in hardness was rationalized by the agglomeration of micro-sized yttrium oxides, which weakened the material [19]. In terms of compressive properties, a wide range of compressive strength (900-1523 MPa) was reported [17,18,30,32,33], whereas no strength/ductility data were found in tension. With the ubiquitous existence of cracks, it is not surprising to see rather poor mechanical property data on L-PBF W. \section*{3. Laser directed-energy-deposition} \subsection*{3.1. Method} L-DED is an AM technique in which metal powder is fed into a melt pool created by a laser. After the first layer is deposited the powder feeder moves upward and deposition of the second layer begins. L-DED is usually conducted under an argon atmosphere that utilizes an argon blower. L-DED is suitable for the manufacture of large parts relatively quickly and offers excellent design freedom due to the additional parameters involved in the process. Some unique features of L-DED are powder feed rate and the option to change input powder composition using multiple hoppers to manufacture composite materials or grade composition throughout AM parts. L-DED can also be used to repair parts due to its ability to accurately deposit material anywhere in the build chamber. These features allow L-DED to process functionally graded structures, which can mitigate the challenges associated with joining dissimilar materials such as $\mathrm{W}$ and ferritic-martensitic steels. Fig. 6 displays a schematic of the L-DED process [39] showing the laser power source, powder feeder, and build platform. In this section, we will assess the current state-of-the-art in L-DED W and tungsten alloys and the effects of L-DED processing on their structure and properties. To understand these processes and how to successfully implement them, it is critical to recognize the processing challenges and defects associated with L-DED of tungsten and its alloys, and possible mitigation strategies. \subsection*{3.2. L-DED tungsten and tungsten alloys} \subsection*{3.2.1. Deposition of pure tungsten} Several studies have demonstrated that tungsten can be printed with moderate success utilizing L-DED. However, these studies also reported difficulties in fully melting tungsten powder during deposition. Polygonal tungsten powder was printed on a reduced activation ferritic/ martensitic (RAFM) steel substrate in [40]. Single tracks were printed using the processing parameter combinations in the range of $\mathrm{P}=2000-$ $4000 \mathrm{~W}$ and $\mathrm{v}=200-600 \mathrm{~mm} / \mathrm{min}$ at a constant powder feed rate $(\dot{\mathrm{m}}=$ $29.3 \mathrm{~g} / \mathrm{min}$ ). Significant compositional mixing was observed between the steel substrate and powder, with tungsten content ranging from 12 to $55 \mathrm{wt} \%$ at various locations within the melt pools. Fig. 7 [40] displays single track melt pool cross sections at various parameter sets. Unmelted tungsten particles can be observed throughout the melt pools at each of the parameter sets displayed in Fig. 7. These unmelted tungsten particles were also observed in single laser clads of tungsten and tungsten-nickel alloys printed on a mild steel substrate studied in [41]. Mixing between substrate material and the deposited tungsten reportedly increased with laser power, which corresponded to a decrease in overall tungsten content and hardness of the single tracks [40]. This is due to increased laser penetration into the steel substrate, increasing the relative amounts of $\mathrm{Fe}$ and $\mathrm{Cr}$ in the melt pool. Microstructural evaluation of multi-layer tungsten prints was also conducted [40]. Intermetallic precipitates were observed in the scanning electron micrograph (SEM) in Fig. 8a and were identified to be $\mathrm{Fe}_{7} \mathrm{~W}_{6}$ from transmission electron microscopy (TEM) analysis (Fig. 8b and c). Energy dispersive spectroscopy (EDS) analysis of the precipitates was consistent with the TEM observations of $\mathrm{Fe}_{7} \mathrm{~W}_{6}$. X-ray diffraction peaks of single- and double-layer prints on the RAFM steel substrate (Fig. 8e) identified the existence of W, $\mathrm{Fe}, \mathrm{Fe}_{7} \mathrm{~W}_{6}$, and $\mathrm{Fe}-\mathrm{Cr}$ phases. In a 9-layer print conducted at $3000 \mathrm{~W}$ and $3000 \mathrm{~mm} / \mathrm{min}$, the authors observed a significant compositional gradient from $25.23 \mathrm{wt} \% \mathrm{Fe}$, $2.73 \mathrm{wt} \% \mathrm{Cr}$, and $72.04 \mathrm{wt} \% \mathrm{~W}$ near the substrate to $3.91 \mathrm{wt} \% \mathrm{Fe}$, $0.28 \mathrm{wt} \% \mathrm{Cr}$, and $95.82 \mathrm{wt} \% \mathrm{~W}$ at the top of the build, as can be seen in Fig. 9 [40]. These observations show that although tungsten content increases with each additional layer, $\mathrm{Fe}$ and $\mathrm{Cr}$ from the substrate continue to diffuse into the upper layers of the matrix after 9 deposited layers of tungsten. SEM analysis of this multi-layer build indicates the existence of unmelted W-rich particles, dendrite structures, and microcracking within the tungsten particles. Overall, this study highlighted the need for designing a compositionally graded transition from steels to tungsten to avoid producing the undesirable intermetallic phases. Cracking and porosity were observed within tungsten single tracks, single layers, and multi-layer deposits as can be seen in Fig. 9j \& k [40]. Cracks were speculated to be due to liquation cracking and residual stresses from the rapid heating and cooling cycles. Single layer cracking occurred at the top of the deposit and propagated toward the substrate causing $\mathrm{W}$ particles along the crack path to internally fracture and others to de-bond with the matrix. Spherical porosity also occurred within single layers of deposited tungsten, likely due to gas trapped within the melt pool during solidification. Pores are a typical defect observed in materials manufactured via L-DED and can be mitigated by manipulating process parameters [42]. Multi-layer tungsten deposits contained tungsten particles with a large number of microcracks near the top of the deposits, as can be seen in Fig. 9d \& g [40]. These microcracks did not appear as frequently in the lower layers of the build. This may be due to the higher degree of remelting experienced at the bottom of the sample, relative to the top. Many commercially available L-DED machines are not capable of achieving the large laser powers utilized in reference [40]. A study [43] attempted to print $12 \mathrm{~mm}$ high thin-walled vertical tungsten tubes using L-DED using polyhedral $99.7 \%$ purity tungsten powder at lower values of laser power. They printed 35 samples with parameters ranging between $\mathrm{P}=600-1000 \mathrm{~W}, \mathrm{v}=50-350 \mathrm{~mm} / \mathrm{min}, \dot{\mathrm{m}}=5-25 \mathrm{~g} / \mathrm{min}$, and a laser diameter $=750 \mu \mathrm{m}$ to determine optimal parameters for printing the material. With a fixed layer thickness of $100 \mu \mathrm{m}$ and a constant but undisclosed hatch spacing, they reported 7 of the 35 parameter sets reached the targeted build height with the rest of the specimens either \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-07} \end{center} Fig. 6. A schematic of the directed-energy-deposition process showing the laser source, powder feeder or sprayer, and the build platform [39]. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-08} \end{center} Fig. 7. Single tracks fabricated at various parameter sets (a-f), displaying different melt pool dimensions as well as varying levels of unmelted W particles [40]. exceeding or falling short of the target range, as can be seen in Fig. 10 [43]. However, no analysis of microstructural homogeneity or degree of melting the tungsten particles was conducted in the study. \subsection*{3.2.2. Deposition of tungsten alloys} 3.2.2.1. Tungsten-nickel deposition. Due to the various applications of tungsten alloys, many studies attempted to print tungsten alloys. One group successfully printed a $60 \mathrm{~W}-40 \mathrm{Ni}$ collimation component (Fig. 11a) using L-DED with the following parameters [41]: $2000 \mathrm{~W}$ laser power, $300 \mathrm{~mm} / \mathrm{min}$ scan speed, and $8 \mathrm{~g} / \mathrm{min}$ feed rate. Single laser clads of $60 \mathrm{~W}-40 \mathrm{Ni}$ printed on a mild steel substrate were observed to contain unmelted tungsten particles and dendritic structures that were speculated to be $\mathrm{Ni}_{4} \mathrm{~W}$ and $\mathrm{NiW}_{2}$ intermetallic phases. Another group similarly observed unmelted tungsten particles in a LDED W-15Ni alloy along with a $\gamma$-Ni phase containing $15 \mathrm{wt} \% \mathrm{~W}$ [44]. They reported a layered microstructure with unmelted tungsten particles dominating regions of the deposit that were only subjected to initial melting, and $\mathrm{W}$ dendrite structures in regions that were subjected to remelting by the subsequent layer (Fig. 11d and e). These W-15Ni specimens had tensile strengths of $\sim 500 \mathrm{MPa}$ and were prone to brittle fracture ( $\sim 3 \%$ strain to fracture) at room temperature. 3.2.2.2. Tungsten-nickel-iron deposition. Unmelted W-rich particles embedded in an FCC Ni-Fe matrix were observed in L-DED manufactured $90 \mathrm{~W}-7 \mathrm{Ni}-3 \mathrm{Fe}[45,46]$, similar to what was observed in $\mathrm{W}$ and W-Ni deposits. An alternating microstructure with layers dense in unmelted tungsten and partially melted tungsten particles was also observed [45,46], similar to those observations in W-15Ni, Fig. 11 [41, 44]. Tensile testing of $90 \mathrm{~W}-7 \mathrm{Ni}-3 \mathrm{Fe}$ samples resulted in high strength (1037 MPa) and low ductility (3.5\% elongation). These as-printed materials displayed significantly higher ultimate tensile strength (UTS) and lower ductility than those of traditionally manufactured $90 \mathrm{~W}-7 \mathrm{Ni}-3 \mathrm{Fe}$ via liquid phase sintering (LPS), Fig. 12a [45]. Additionally, large periodic variations in microhardness ( $\sim 76 \mathrm{HV}$ ) were observed, Fig. 12b. These variations are attributed to a periodic sublayer change where $\mathrm{W}$-particle dense regions are observed above and below the regions with lower relative amounts of W particles. L-DED materials have an average hardness of $\sim 415 \mathrm{HV}$, higher than specimens made by LPS. This is due to higher amounts of hard W-particles embedded in the matrix in L-DED specimens. Fig. 13 displays fracture surfaces of $90 \mathrm{~W}-7 \mathrm{Ni}-3 \mathrm{Fe}$ where large pores were observed, indicating that the porosity acts as fracture initiation sites [45]. Lower ductility during tensile testing in the as-deposited specimens compared to traditional LPS is attributed to the residual porosity observed at these fracture surfaces. Tungsten particle cleavage and tearing of the ductile Ni-Fe matrix were also observed at the fracture surfaces. \subsection*{3.3. Challenges in L-DED tungsten} Common challenges associated with L-DED $\mathrm{W}$ on ReducedActivation-Ferritic-Martensitic (RAFM) steels are illustrated in Fig. 14. There are many critical parameters that affect melt pool morphology during L-DED, including laser power (P), scan speed (v), hatch spacing (h), laser focus (f), substrate temperature, powder size distribution and morphology, and powder feed rate ( $\dot{m}$ ). These variables are critical in \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-09(1)} \end{center} (e) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-09} \end{center} Fig. 8. Phase analysis of pure W printed on a RAFM steel substrate using (a) scanning electron microscopy (SEM), (b, c) transmission electron microscopy (TEM), (d) energy dispersive spectroscopy (EDS), and (e) x-ray diffraction (XRD) [40]. achieving successful prints in L-DED. Only a limited number of these variables have been explored in L-DED W and tungsten alloys. Challenges such as attaining targeted build heights and mitigating porosity can be resolved by optimizing parameters such as $\mathrm{P}, \mathrm{v}, \mathrm{h}$, and $\dot{\mathrm{m}}$ as well as improving feedstock quality [42, 47-49]. Additionally, residual stress-induced cracking has been shown to be mitigated by substrate preheating [50-53] which, to our knowledge, has not been attempted on L-DED W. Substrate preheating may also result in increased melting of tungsten particles that would remain unmelted if printed on a room temperature substrate. Mitigation of intermetallic particle formation may be achieved by introducing filler alloys between the steel base plate and tungsten, circumventing regions in the alloys' phase diagrams in which detrimental phases are stable. These strategies may improve the feasibility of additively manufacturing tungsten via directed energy deposition. \section*{4. Electron beam melting} \subsection*{4.1. Method} The EBM or EB-PBF process belongs to the powder bed family of additive manufacturing technologies. Similar to L-PBF, the heat source is selectively moved across the powder bed to melt the regions of interest in a layer-by-layer process. Electron beams, compared to lasers, are high in energy density exceeding several kilowatts focused into spot sizes of several hundred microns in diameter when melting. EB-PBF occurs under controlled vacuum conditions to both maintain the quality of the electron beam spot size and offset fluctuation in pressure associated with \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-10} \end{center} Fig. 9. Scanning electron microscope (SEM) and energy dispersive spectroscopy (EDS) analysis of a 9-layer W sample printed at $3000 \mathrm{~W}$ and $3000 \mathrm{~mm} / \mathrm{min}$. (a, b) Low magnification cross-sectional SEM images of the specimen, indicating where EDS analysis was conducted. (c) The results of EDS analysis conducted on the areas displayed in (b). (d-f) Higher magnification images of the top, middle, and bottom of the specimen, respectively. (g-i) High magnification images displaying W particles, dendrite structures, and microcracks [40]. (j, k) Low and high magnification images displaying cracks and porosity in W deposited on RAFM steel [40]. vaporization of the metal in the liquid state as the beam is melting. Additionally, in EB-PBF the powder bed is heated to elevated temperatures through defocusing the electron beam and rapidly rastering it across the powder bed surface to allow for the powder particles to loosely sinter to one another and conduct the negative charge of the electron beam away. If the negative electrical charge of the imparted electrons is not conducted away, the powder bed will build-up a negative charge which results in the repulsion of the powder particles from one another, i.e., a "smoking" event. In the instance of tungsten, the powder bed is heated to between 1000 and $1400{ }^{\circ} \mathrm{C}$ [54]. As a result of heating the powder bed, materials processed through EB-PBF often have lower levels of residual stress than corresponding materials processed through L-PBF [54]. One of the advantages of the EB-PBF process over that of L-PBF is the ability to rapidly manipulate the electron beam heat source over the entirety of the build area to locally control thermal conditions of the material. This has been shown as beneficial for controlling the microstructure [55] as well as stress states of the material to suppress defects such as cracks in non-weldable materials [56].\\ (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-11(2)} \end{center} $10 \mathrm{~mm}$ (b) \section*{Proper build domain} \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-11(1)} \end{center} $10 \mathrm{~mm}$\\ Over build domain \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-11} \end{center} $10 \mathrm{~mm}$ (c)\\ \includegraphics[max width=\textwidth, center]{2024_04_13_b812db9e8fc840c3b338g-11(3)} Fig. 10. Tungsten specimen fabricated at various parameter sets displaying heights that are either below, matching, or exceeding the targeted build height [43]. (a) Images of specimens that are under built, over built, or matching the targeted height. (b) The measured height increment for each layer in samples printed at different speeds plotted against laser power. (c) Cubes printed at various parameters in the under build, over build, or target printing regimes. \subsection*{4.2. Porosity} Various levels of success in the processing of tungsten through EBM (EBM W) have been reported. Key to the processing of tungsten is the ability to obtain materials that approach the theoretical density of pure tungsten at $19.3 \mathrm{~g} / \mathrm{cc}$. In literature, four states for EBM $W$ based on density and porosity of the material have been identified, as depicted in Fig. 15 [57]. In the figure, volumetric energy densities ranging from 208 to $3840 \mathrm{~J} / \mathrm{mm}^{3}$ with a substrate temperature of approximately $850{ }^{\circ} \mathrm{C}$ were used. These states are (a) limited fusion, (b) insufficient fusion, (c) proper fusion, and (d) excessive fusion. Limited fusion is characterized by resultant relative densities of $<70 \%$ with excessive balling of the tungsten observed within the melt layers. Insufficient fusion exhibits relative densities between $70 \%$ and $90 \%$, however, significant interconnected porosity exists within the material including chimney porosity [58]. Proper fusion is the optimal processing state where densities greater than $90 \%$ are achievable and interconnected porosity is mitigated through full melting and wetting of the tungsten. The last state, excessive fusion, can be defined as fully dense material that exhibits swelling due to too much energy being imparted into the material. Similar trends were also identified in work that varied the linear energy used to melt tungsten from 333 to $5000 \mathrm{~J} / \mathrm{m}$ with a powder bed temperature of $1000{ }^{\circ} \mathrm{C}$ [59]. Various levels of porosity have been reported in the literature, with many studies achieving success for high density tungsten. Densities as high as $\mathbf{9 9 . 5 \%}$ have been measured; nevertheless, microcracking was found $[57,59]$. SLM and EBM W have also been compared and it was found that comparable densities could be obtained using either technique, although the build temperature greatly influenced the porosity level and defect levels [60]. Lastly, nondestructive techniques such as in-situ near-infrared (IR) defect detection have been used to report defect-free, highly dense samples (>99\%) [61,62]. \subsection*{4.3. Cracking} The occurrence of cracking in EBM $\mathrm{W}$ is a problem akin to that observed in SLM processed tungsten with the debate ongoing for the specific mechanism(s) by which tungsten cracks during processing. It has been suggested that cracks occur in the solid state as a result of significant inelastic deformation along grains neighboring GBs [61]. This was supported through electron backscatter diffraction (EBSD) analysis of the areas surrounding cracks that revealed localized orientation gradients near the edges of cracks. Representative EBSD micrographs showing the cracking in tungsten are shown in Fig. 16 [61]. This is consistent with other studies that observed the cracking phenomena through high-speed in-situ videos of the SLM process with a heated powder bed temperature range above and below the DBTT range of tungsten. This was attributed to the development of significant von Mises stresses when the tungsten cycled below DBTT [12]. It was theorized that thermal stress generated from thermal gradients during SLM processing of tungsten can only be compensated by crack formation along low-strength GBs, particularly those with impurities [19,63]. From the reported studies, the influence of the build substrate temperature and the overall build temperature has a clear influence on the cracking in AM W. A report [29] utilized a SLM system with substrate preheating and showed that increasing build temperature from $200{ }^{\circ} \mathrm{C}$ to $1000^{\circ} \mathrm{C}$ significantly reduced cracking in AM tungsten, though it did\\ \includegraphics[max width=\textwidth, center]{2024_04_13_b812db9e8fc840c3b338g-12(2)} Fig. 11. Laser directed-energy-deposition (LDED) of W-Ni alloys. (a) W60-Ni40 collimation component printed at $2000 \mathrm{~W}$ laser power, $300 \mathrm{~mm} / \mathrm{min}$ scan speed, and $8 \mathrm{~g} / \mathrm{min}$ feed rate [41]. (b, c) Low and high magnification SEM micrographs showing the microstructure of a single laser clad of W60-Ni40 printed on a mild steel substrate [41]. (d) Scanning electron micrographs of a W-15Ni L-DED part showing a layered microstructure containing unmelted $\mathrm{W}$ particles in regions subjected to initial laser melting [44], and (e) dendritic W structures in remelted regions of the deposit [44]. (a) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-12} \end{center} (b) \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-12(1)} \end{center} Fig. 12. Mechanical properties of additively manufactured W-Ni-Fe alloys. (a) Engineering stress-strain curves for $90 \mathrm{~W}-7 \mathrm{Ni}-3 \mathrm{Fe}$ specimens manufactured by laser metal deposition (LMD, alternatively directed-energy-deposition) and liquid phase sintering (LPS). (b) Hardness values along the build direction for 90W-7Ni-3Fe specimens manufactured by LMD and LPS [45]. not eliminate cracking entirely. Multiple studies explored the role of substrate heating: samples have been built with a powder bed temperature of approximately $850{ }^{\circ} \mathrm{C}$, and observed minor levels of cracking [57]. In a comparison between SLM and EBM, significant cracking was found in SLM, with no cracking in EBM. The lack of cracking in EBM W was attributed to a combination of build plate temperature of $1000{ }^{\circ} \mathrm{C}$ and addition of a support structure to raise the tungsten samples off the build plate $[19,60]$. Theoretically, if the substrate is heated above the\\ \includegraphics[max width=\textwidth, center]{2024_04_13_b812db9e8fc840c3b338g-13(1)} Fig. 13. Tensile fracture surfaces of $90 \mathrm{~W}-7 \mathrm{Ni}-3 \mathrm{Fe}$ fabricated via laser directed-energy-deposition (L-DED). Features such as (a) porosity, (b-d) W particle cleavage, and (d) matrix failure are indicated with white arrows [45]. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-13} \end{center} Fig. 14. Graphical representation of challenges typically encountered during directed energy deposition of tungsten on RAFM steel. These challenges include achieving targeted build heights, cracking and porosity, difficulty of melting W particles, formation of intermetallics and dendrites, and layered melted/remelted structures within the build. DBTT of tungsten, dislocations should be significantly more mobile, thus reducing the chance of cracking. Higher substrate temperatures allow for lower temperature gradients during cooling, reducing the stresses generated on the tungsten components. Similarly, by raising the printed components of the build plate by using a support structure, heat is not allowed to dissipate quickly, which in turn reduces the stresses on the components. The use of different metals, such as steel and titanium, as build substrates was also investigated [61]. Titanium build plates have been used due to the high degree of solubility the elements have in one another in an effort to create a metallurgical bond at the interface of part and build plate. Cracking in tungsten was observed to be sensitive to the build preheat temperature. Crack density was drastically lower when build surface temperatures of $1500{ }^{\circ} \mathrm{C}$ were used, compared to $1100{ }^{\circ} \mathrm{C}$.\\ Additional studies leveraged an ever-higher surface preheat temperature of $1800^{\circ} \mathrm{C}$ to demonstrate the ability to successfully suppress crack formation in EBM W [62]. Mitigation techniques for suppressing cracking in tungsten aside from the processing science include alloying tungsten with elements such as tantalum. However, for nuclear fusion application tungsten-tantalum alloys are generally considered to be problematic due to tantalum's degradation into the undesirable isotopes during nuclear exposure [64]. \subsection*{4.4. Microstructure} Rather distinct to the pure refractory material systems processed through AM such as tungsten, as well as some common BCC or HCP type\\ (a) Limited fusion \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-14(2)} \end{center} (b) Insufficient fusion \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-14(1)} \end{center} (c) Proper fusion \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-14} \end{center} (d) Excessive fusion \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-14(3)} \end{center} \includegraphics[max width=\textwidth, center]{2024_04_13_b812db9e8fc840c3b338g-14(4)}\\ \includegraphics[max width=\textwidth, center]{2024_04_13_b812db9e8fc840c3b338g-14(5)} Fig. 15. Observed states of pure tungsten fabricated through EBM using different parameters. (a) Limited fusion, (b) insufficient fusion, (c) proper fusion, and (d) excessive fusion. The images are adopted from [57]. \begin{center} \includegraphics[max width=\textwidth]{2024_04_13_b812db9e8fc840c3b338g-15(1)} \end{center} Fig. 16. Scanning electron microscope (SEM) and electron backscatter diffraction (EBSD) data of boundary cracking in EBM pure tungsten. (a) EBSD showing surrounding texture, and (b) an optical image [61]. alloys such as Ti-6V-4Al, are the anomalous textures that form as a result of the AM process [65,66]. Shown in Fig. 17 [62] are representative EBSD micrographs depicting this anomalous texture. While many in the literature for EBM of tungsten and other refractory metals have observed the phenomena, its significance in relation to the processing science of the materials has only been briefly mentioned. Columnar grain structures aligned parallel to the build direction resulting from epitaxial growth of EBM W are reported in literature. Controlling the $\{001\}$ and $\{111\}$ fibrous texture of pure tungsten via $67^{\circ}$ interlayer rotation is also explored. Similar observations regarding the effect of texture on yield strength anisotropy have also been reported for tantalum $[57,60,61,67]$. The ability to obtain a mixed $\{001\}$ and $\{111\}$ fiber texture with the possibility of material having either strong $\{001\}$ or $\{111\}$ build direction fiber textures was also identified. In a similar study of EBM of molybdenum, the role played by area energy density to melt the material on the texture was also discussed [61,63]. This phenomenon was hypothesized to be associated with sensitivities of the melt pool shape to the electron beam energy density coupled with\\ \includegraphics[max width=\textwidth, center]{2024_04_13_b812db9e8fc840c3b338g-15} Fig. 17. Electron backscatter diffraction (EBSD) images showing mixed $\{001\}$ and $\{111\}$ fibers in pure EBM tungsten. (a) Cross-sectional inverse pole figure (IPF) map, and (b) build direction IPF map [62]. The build direction is vertical for both images.\\ formation of networks of LAGBs driven by thermal stresses from solidification. \subsection*{4.5. Performance of EBM Tungsten} Analysis of the performance of EBM $\mathrm{W}$ for thermomechanical behavior is currently limited. The bend strength of an EBM W during three-point bending test was measured to be $340 \mathrm{MPa}$, which is significantly lower than reference wrought tungsten [60]. This has been partially attributed to the porosity of the EBM W samples. The strength of EBM W parallel to the build direction was also evaluated, with fracture being observed to occur along the fibrous GBs via a combination of decohesion and transgranular failure [57,60]. Lastly, the hardness of EBM W as well as its surface deterrence to ITER-like plasma heat load exposures at steady state $\left(10 \mathrm{MW} / \mathrm{m}^{2}\right.$ ) and transient (105 pulses with $0.14 \mathrm{GW} / \mathrm{m} 2$ ) was also investigated [59]. Ultimately, it was found that the EBM W performed similarly to baseline wrought recrystallized tungsten product as surface deterrence to plasma heat load exposures. \section*{5. Summary, outlook, and recommendations} Although tungsten and tungsten alloys are notoriously difficult to print due to their high melting temperatures, high thermal conductivities, and brittleness, encouraging progress has been made in the last decade to additively manufacture this unique class of materials. Cracking has been and remains to be one of the dominantly challenging issues in the field. Nevertheless, advance has been made to overcome this issue. For example, EBM has shown promises to manufacture crack free samples. Another important issue, which has not been investigated to a large extent, is the pore formation and control mechanisms under keyhole mode processing conditions. In addition to cracks, porosities inevitably influence the mechanical properties of additively manufactured tungsten and tungsten alloys. Up to date, limited mechanical property data are available (especially those related to elevated temperature properties) that are of critical relevance to practical applications. With the rapid progress of additive manufacturing techniques and processing conditions, we expect to witness a rising amount of data in this direction. In addition, meticulous microstructure control and new alloy design strategies are expected in near future for this class of high temperature alloys. We further contend that many issues encountered during refractory metals additive manufacturing are likely applicable to numerous fracture-prone metals such as multi-principal alloys. Strategies are needed to enable us to "print these unprintable" alloys. \subsection*{5.1. Recommendations} Substrate heating has been proved to be an effective strategy to overcome cracking in the EBM process. Similar approaches have not been demonstrated successfully for L-PBF or L-DED. This could be partially due to the higher oxygen contents in the laser processes. Further studies are needed in this direction to make the laser AM processes feasible to manufacture crack-free W components. Microstructure control such as grain shape manipulation has been reported to be effective in reducing the residual stresses and thus cracks in brittle materials. This approach has not been well studied for tungsten and tungsten alloys yet, or any other refractory alloys. Inoculation via the addition of nanoparticles could be another rewarding strategy to overcome the cracking issue for tungsten. In addition to experimental endeavors, computer modeling of thermal history, microstructure, and resultant residual stresses is likely to further advance this field. The above recommended research directions are likely applicable to all three AM techniques reviewed in this work. Revision note: while this paper was under review, a parallel overview paper was published [68]. \section*{CRediT authorship contribution statement} Talignani Alberico: Writing - original draft, Investigation, Conceptualization. Seede Raiyan: Writing - original draft, Investigation, Conceptualization. Whitt Austin: Writing - original draft, Investigation. Zheng Shiqi: Investigation, Conceptualization. Katoh Yutai: Writing - review \& editing, Supervision, Project administration, Investigation, Funding acquisition, Conceptualization. Wang Y. Morris: Writing - review \& editing, Writing - original draft, Supervision, Project administration, Investigation, Funding acquisition, Conceptualization. Kirka Michael M: Writing - review \& editing, Writing - original draft, Supervision, Investigation, Conceptualization. Ye Jianchao: Supervision, Investigation. Karaman Ibrahim: Writing - review \& editing, Supervision, Investigation, Funding acquisition, Conceptualization. \section*{Declaration of Competing Interest} The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. \section*{Data availability} No data was used for the research described in the article. \section*{Acknowledgments} This research was sponsored by the US Department of Energy, Office of Fusion Energy Sciences and Advanced Research Projects AgencyEnergy (ARPA-E) under contract DE-AC05-00OR22725 with UTBattelle LLC. The work at LLNL was performed under the auspices of the US Department of Energy under contract no. DE-AC52-07NA27344. \section*{References} [1] M. Rieth, S.L. Dudarev, S.M.G. de Vicente, J. 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