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Review

A Review of Additive Manufacturing Techniques and Post-Processing for High-Temperature Titanium Alloys

by
Binquan Jin
1,
Qing Wang
1,
Lizhong Zhao
1,
Anjian Pan
1,
Xuefeng Ding
2,*,
Wei Gao
2,
Yufeng Song
2 and
Xuefeng Zhang
1,*
1
Institute of Advanced Magnetic Materials, College of Materials and Environmental Engineering, Hangzhou Dianzi University, Hangzhou 310018, China
2
College of Mechanical and Electrical Engineering, Hunan University of Science and Technology, Xiangtan 411201, China
*
Authors to whom correspondence should be addressed.
Metals 2023, 13(8), 1327; https://doi.org/10.3390/met13081327
Submission received: 15 June 2023 / Revised: 19 July 2023 / Accepted: 20 July 2023 / Published: 25 July 2023
(This article belongs to the Special Issue Additive Manufacturing of Non-ferrous Alloys)
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Abstract

:
Owing to excellent high-temperature mechanical properties, i.e., high heat resistance, high strength, and high corrosion resistance, Ti alloys can be widely used as structural components, such as blades and wafers, in aero-engines. Due to the complex shapes, however, it is difficult to fabricate these components via traditional casting or plastic forming. It has been proved that additive manufacturing (AM) is an effective method of manufacturing such complex components. In this study, four main additive manufacturing processes for Ti alloy components were reviewed, including laser powder bed melting (SLM), electron beam powder bed melting (EBM), wire arc additive manufacturing (WAAM), and cold spraying additive manufacturing (CSAM). Meanwhile, the technological process and mechanical properties at high temperature were summarized. It is proposed that the additive manufacturing of titanium alloys follows a progressive path comprising four key developmental stages and research directions: investigating printing mechanisms, optimizing process parameters, in situ addition of trace elements, and layered material design. It is crucial to consider the development stage of each specific additive manufacturing process in order to select appropriate research directions. Moreover, the corresponding post-treatment was also analyzed to tailor the microstructure and high-temperature mechanical properties of AMed Ti alloys. Thereafter, to improve the mechanical properties of the product, it is necessary to match the post-treatment method with an appropriate additive manufacturing process. The additive manufacturing and the following post-treatment are expected to gradually meet the high-temperature mechanical requirements of all kinds of high-temperature structural components of Ti alloys.

1. Introduction

High-temperature titanium alloys can be widely used as structural components in the field of aerospace by virtue of their excellent mechanical properties such as specific strength, low density, corrosion resistance, and high-temperature resistance [1,2]. However, traditional casting and forging methods are disadvantaged in manufacturing complex-shaped titanium alloy components. Additive manufacturing, which uses accumulated powder to produce complex-shaped titanium products directly from digital models, can simplify manufacturing, reduce costs, and shorten development cycles [3]. Moreover, compared with the traditional processing method, the AM technique can achieve better mechanical properties [4,5,6,7,8,9,10,11], thereby opening up a new way for the development and production of high-temperature titanium alloys.
Additive manufacturing processes are diverse, in which powder bed fusion (PBF) and directed energy deposition (DED) are the two mainstream methods for printing titanium alloys [12]. As an emerging technology with great potential, CSAM can reduce the oxidation of sprayed powders due to the low-temperature and high-speed characteristics of sprayed particles. PBF uses a focused high-energy laser beam or electron beam to selectively melt the pre-laid powder layer in the powder bed [13,14,15]. Selective laser melting (SLM) and electron beam melting (EBM) are two typical PBF metal printing processes. DED uses lasers, electron beams or electric arcs to melt metal powders or wires deposited along printed paths [16,17]. CSAM realizes titanium alloy printing via deformation stacking of the substrate impacted by high-speed powders. AMed titanium alloy’s microstructure varies according to the printing principle, printing environment, and heat source, which lead to the differences in the mechanical properties and applicable environment of printing titanium alloys [18]. The defense sector has a large demand for high-performance titanium alloys. The lack of a clear understanding and positioning of the method in the additive manufacturing of titanium alloys may lead to defects in the mechanical properties and heat resistance of materials, affecting the safety and reliability of national defense equipment, which may lead to equipment failure or incompetence for their design purposes. The improper use of additive manufacturing methods can lead to excessive costs and prolonged production cycles, which can lead to project cost overruns and delayed project completion, negatively impacting the defense industry and national security [19,20,21]. Therefore, the four main additive manufacturing processes for Ti alloy components were reviewed in this study, including SLM, EBM, WAAM, and CSAM. Meanwhile, the technological process and mechanical properties at high temperatures were summarized. Moreover, the corresponding post-treatment was also analyzed to tailor the microstructure and high-temperature mechanical properties of Amed Ti alloys.
The mechanical properties of as-printed printed titanium alloys can be effectively improved by the study of the printing mechanism, optimization of process parameters, in situ addition of trace elements, and layered material design, which are the four main research directions regarding the performance improvement of as-printed titanium alloys. As a follow-up process to further enhance product performance, post-processing has become the key to further improving the quality of AMed titanium alloy products with its continuous development and the process connection between post-processing and additive manufacturing technology.

2. Additive Manufacturing

To print titanium alloy components with complex structures and meet the requirements of different operating environments, it is necessary to select an appropriate additive manufacturing method with net-shape printing capability. With regard to this, representative AM methods were successively selected from mature technologies to emerging technologies: selective laser melting (SLM), electron beam melting (EBM), wire arc additive manufacturing (WAAM), and cold spraying additive manufacturing (CSAM). Figure 1 shows the working principle of each method. SLM and EBM belong to PBF, which selectively scans and melts metal powder on the powder bed according to the stratified CAD data with a high-energy beam as the heat source. WAAM belongs to directed energy deposition, using the layer-by-layer cladding principle and an arc as the heat source to melt metal wire and deposit metal parts layer by layer in the designed path. Based on aerodynamics, CS uses compressed gas to give the powder a supersonic speed, which leads to plastic deformation and deposition when colliding with the base and forms metal parts through the design of the nozzle path.

2.1. SLM

The appropriate applications of the SLM process can be determined by analyzing its advantages and disadvantages. SLM is a process of printing metals in an inert gas environment using a high-power laser as the heat source. The high-energy laser layer-by-layer melting method can produce complex parts with high precision, good surface quality, and relatively high density without additional processing. Table 1 shows the mechanical properties of SLMed titanium alloys. The yield strength and ultimate tensile strength of as-printed titanium alloy samples are nearly 50% higher than that of forged samples, while the maximum elongation is only about half of that of forged samples [26]. However, in the SLM process, the local high-energy input of the high-energy laser beam will lead to an uneven internal temperature of the metal, resulting in a thermal gradient, high residual stress, and uneven deformation of the structure [27,28]. Moreover, SLM can only print small-size and small-batch components due to equipment limitations, so SLM is suitable for printing small-size and low-precision titanium alloy components. SLM has emerged as a versatile additive manufacturing technique for fabricating titanium alloys, finding widespread applications in industries such as aerospace, medical, automotive, and energy. SLM-printed titanium components are utilized in producing lightweight aircraft parts, patient-specific medical implants, customized automotive components, and high-performance energy storage devices.
Currently, the research on SLM-printed titanium alloys mainly focuses on process parameters, post-processing techniques, and in situ element addition to improve their performance, among other aspects. This section describes the microstructure, process parameters, and in situ element addition of SLM-printed titanium alloys, while post-treatment techniques are elaborated on in the next chapter. The SLM process generates a high-temperature gradient, which causes the accumulation of thermal stress, and a high cooling rate leads to the formation of a microstructure dominated by columnar β crystals [37] and α’ martensite intracrystalline substructures [38]. However, since different process parameters will impact titanium alloys’ microstructure, and mechanical properties, much research work has been carried out on parameter optimization. Firstly, it is reported that different scanning strategies could affect each layer’s remelting condition and thermal gradient, thus affecting the grain size, porosity, and anisotropy of the printing component. However, an optimal solution has not been obtained due to the diversity of scanning strategies. Currently, the commonly used scanning strategies mainly include dual-channel interlayer rotation and island scanning strategies [39,40,41,42,43,44,45]. Secondly, Cai et al. [36] studied the effect of energy density on the microstructure evolution of SLMed TA15 titanium alloys and found that low energy density would lead to inadequate melting and thus generate pores, while high energy density would cause excessive melting and result in a microscopic sphericity effect, which is consistent with the research results of Liverani et al. [46] and Wei et al. [47]. In the appropriate energy density range, with the increase in energy density, the sample’s microstructure is gradually transformed from coarse and equiaxed particles to columnar particles. Meanwhile, the mechanical properties of titanium alloys are also different, as shown in Table 2. Finally, new findings have been made in the study on the influence of scan spacing parameters. Yao et al. [48] explained the mechanism of a martensitic colony leading to the concentration of dislocation accumulation in the column β grain boundary, resulting in low ductility, and found that moderately increasing the scan spacing could reduce the formation of the martensitic colony, which improved the ductility of Ti64 of SLMed by 194%. It is difficult to confirm multi-parameter values, but the Taguchi method provides an efficient method for parameter confirmation. Liu et al. [49] used the Taguchi method to optimize the parameters of TA15 printed by the SLM method, effectively reducing the number of multi-parameter experimental groups, improving the efficiency of selecting the optimal parameters, and verifying the validity of the optimized parameter results. It has guiding significance for parameter optimization of SLM printing titanium alloys. Parameter optimization has been continuously investigated.
On the other hand, initial efforts have been made to design SLMed unique alloys. For example, Zhang et al. [35] adopted a collaborative alloy design method and printed titanium alloys with layered α + β in each layer by adding pure titanium powder and iron oxide nanoparticles into Ti6Al4V raw material through in situ alloying, as shown in Figure 2. The anisotropy of mechanical properties caused by the thermal gradient and different microstructures between thermal history cambium layers is overcome. Zhang et al. [34] added a small amount of 316L stainless steel to Ti6Al4V by in situ alloying. Micron-scale modulation of the elements contained in 316L of Ti–6Al–4V matrix formed, a fine-scale modulated β + α’ dual-phase microstructure with a high tensile strength of about 1.3 gigapascals, elongation of 9%, and excellent working hardening mechanical properties of >300 MPa. These successes demonstrate the great potential of in situ alloying to modulate the microstructure of unique SLMed titanium alloys, which plays an important guiding role for subsequent research directions.
In summary, in terms of structural properties, SLM-printed titanium alloys form columnar β grains and martensitic phases due to the high cooling rate, columnar β grains will give rise to an anisotropic increase, and a large number of brittle and hard α’ phases lead to the high ultimate tensile strength of SLMed titanium alloys, but the ductility is weak. In terms of performance improvement, the two-channel interlayer rotation strategy and the island scanning strategy are conducive to a more uniform temperature distribution, a more uniform microstructure of components, and less anisotropy of mechanical properties [50]. As for the energy density, the minimum energy density output should be ensured on the premise of full melting of titanium alloy powder to prevent the coarsening of columnar β grain, since this will result in the decline of mechanical properties. Meanwhile, the ductility of titanium alloys can be improved by moderately increasing the scanning spacing. The Taguchi method can effectively improve the optimization efficiency for the optimization of the abovementioned parameters. In terms of development, SLM has achieved results in the design of exclusive titanium alloys by in situ alloying titanium and its alloys with element doping, which carves a path for subsequent research.

2.2. EBM

The characteristics of EBM also determine its application scenarios. EBM prints metals in a vacuum using an electron beam as a heat source. Unlike SLM, EBM’s higher preheating and processing temperatures enable it to print metal components with low residual stress and high density, and it is reported to be able to manufacture complex defect-free products, especially in producing porous structures based on titanium materials with total density and predefined external shapes and internal structures [51,52]. In the meantime, EBM, which exhibits a higher energy utilization rate and a faster processing speed, can realize multi-layer stacking processing and a large number of single printings, but EBM-printed metal’s precision is relatively low, accompanied by poor surface quality and higher EBM equipment costs [53,54,55]. Hence, EBM is more applicable to industrial mass production of small-sized titanium alloys with low precision requirements. EBM has emerged as a prominent additive manufacturing technique for printing titanium alloys, finding extensive applications in the aerospace, medical, and orthopedic fields. EBM-printed titanium components exhibit favorable mechanical properties, excellent biocompatibility, and the ability to create complex geometries, making them suitable for the manufacturing of patient-specific implants, lightweight aerospace structures, and customized medical devices.
Ti64 has been extensively studied because of its excellent performance and biocompatibility, and EBM has become one of the main additive manufacturing methods for printing Ti64. Better printing performance of Ti64 can be achieved by exploring the microstructure of EBMed Ti64. In the process of EBM manufacturing, the preheating environment of 650 °C to 750 °C and the characteristics of slow cooling lead to the decomposition of the martensitic phase, forming α and β phase structures dominated by the α grain boundary and transformed α/β structure, and the grain is filled with the original β grain with a Widmanstätten structure and a lamellar structure (Figure 3 and Figure 4) [56,57]. This microstructure differs entirely from the α’ and β phase microstructure of SLMed titanium alloys. Preheating also plays a role in stress relief, which is conducive to a good match of strength, plasticity, and performance consistency of EBM-manufactured parts [58,59]. Table 3 shows the mechanical properties of EBMed titanium alloys, compared with forged samples, the tensile strength, yield strength, and elongation of EBMed titanium alloy samples increased by 19%, 12%, and 49%, respectively [60]. Because EBMed titanium alloys do not have α’ martensite, the ultimate tensile strength is reduced, but the ductility is improved, with sub-high strength and high ductility. Wang et al. [61] revealed the possibility of overcoming the strength-ductility tradeoff of titanium alloys through additive manufacturing, and the α-strip width in their microstructure is important for the deformation mechanism of Ti64 samples constructed by EBM. Deformation-induced nanoscale twins were observed at room temperature with an average width of 0.6 μm for the α strips, which resulted in high yield strength and ultimate tensile strength as well as remarkable ductility. In contrast, when the average width of α slats was about 0.2–0.3 μm, the fine microstructure specimens with a low working hardening rate only exhibited dislocation plasticity. These findings provide insights into the microstructure-dependent deformation mechanism of manufactured Ti-6Al-4V.
In terms of structural properties, EBM has a higher preheating temperature to avoid a martensitic phase, forming an α/β structure. Because of the reduction of the martensitic phase, the ultimate tensile strength of EBMed titanium alloys decreases, but remains high thanks to the high density and basket structure, and its ductility is improved by the increase of β phase content. With regard to performance improvement, the width of α slats can be controlled by optimizing parameters to improve mechanical properties. In the development of application, EBM printed titanium alloys display high ductility and relatively high ultimate tensile strength, making them more suitable for the high precision requirements of small-scale titanium alloy production.

2.3. WAAM

WAAM is an additive manufacturing process particularly suitable for industrial-scale production. As one of the most influential and valuable direct energy deposition (DED) technologies, WAAM has various advantages, including a high deposition rate, short production lead time, and a high material utilization rate [69,70]. WAAM is mainly used for repairing metal parts and printing metal of a clean shape. Compared with the forging process, WAAM-printed Ti-6Al-4V has slightly lower average yield strength, ultimate tensile strength, and similar ductility [71,72], as shown in Table 4. At the same time, it breaks through the limitation that the powder bed melting method can only print small-sized components. However, WAAM has the disadvantages of suboptimal surface quality, inadequate dimensional accuracy, and low density due to melt pool interactions [73]. Moreover, the slow cooling rate and significant heat source caused by the WAAM process may lead to higher tensile residual stress and more significant deformation in fabricated components. WAAM has gained prominence as an additive manufacturing technique for fabricating titanium alloys, finding applications in industries such as aerospace, automotive, and maritime. WAAM-printed titanium components are used in the production of large-scale structural parts, such as aircraft wings, automotive chassis, and ship components.
The study of the WAAM process is mainly divided into three aspects: product microstructure, post-processing, and the unique design of titanium alloys with layer-wise element addition. To improve the mechanical properties of WAAM-printed titanium alloys, a large number of studies focus on the study of the microstructure of WAAM-printed titanium alloys and explore the variables affecting their microstructure. Due to the repeated rapid heating and cooling cycles in the WAAM process, the microstructure of WAAM titanium alloys mainly comprises of large columnar initial β particles. In contrast, the microstructure comprises delicate Widmanstätten α layers in the upper layer and coarse Widmanstätten α structures in the lower layer, as shown in Figure 5. At different cooling rates, the high-temperature β phase transforms into different α phases, as shown in Figure 6, including martensitic α’, martensitic α’’,acicular α, acicular secondary α, grain boundary α, and Widmanstätten structures [71,74,81,82,83,84,85]. The coarse columnar β grains of titanium alloy manufactured by WAAM have a potential influence on the strength and uniformity of the mechanical properties of titanium alloys, which is manifested as obvious anisotropy. Therefore, researchers have extensively probed into the WAAM system for grain refinement. Post-processing is mainly incorporated into the WAAM process to form a new "WAAM+" system to achieve grain refinement, e.g., WAAM + inter-pass rolling (IRWAAM) system [86,87], WAAM + inter-pass machine hammer peening (IMHP) system [88], hot wire arc additive manufacturing (HWAAM) [89], ultrasonic impact treatment (UIT) GMAW-based WAAM system [90], ultrasonic peening treatment (UPT) GMAW-based WAAM system [91], and laser shot peening (LSP) GMAW-based WAAM system [92]. The development of "WAAM+" can realize one-step production, considerably shorten the production time, and further magnify the advantage of the low cost of WAAM equipment, which has great significance for industrial development. Direct energy deposition has attracted increasing attention in the field of layered materials due to its unique advantages in the obtained continuous composition [93,94]. Reichardt et al. [95] reported the preparation of Ti6Al4V to 304L gradient components by laser deposition additive manufacturing. Liu et al. [96,97,98] successfully prepared Ti6Al4V/AlSi10Mg graded materials by the direct laser deposition manufacturing process. Compared with laser deposition, WAAM has the advantages of low equipment costs, a high utilization rate of wires, high production efficiency, and high density of the surfacing layer, which is more suitable for the industrial production of layered materials. The preparation of layered materials by laser deposition provides a practical basis for layered materials by the WAAM method, which is also direct energy deposition.
In summary, the microstructure of WAAM printing titanium alloys conprises columnar β crystals and α phases, and the transformation of β phases into different α phases can be controlled by parameter optimization and cooling rate control. In terms of performance improvement, the columnar β crystals will cause anisotropy of mechanical properties, so the "WAAM+" mode integrating post-treatment and additive manufacturing is introduced to refine the grain and reduce the residual stress, so as to achieve better mechanical properties. In the development and application, the layered material has a high match with the layer-by-layer printing WAAM, and the economy of WAAM is more suitable for the industrial scale production of layered materials, with a greater development potential.

2.4. CSAM

As a new emerging process, CSAM exhibits great potentials; however, thee CSAM printed products still suffer from quality problems. CS was originally used for the metal coating process. Recently, it has been proposed as a new additive method for metal manufacturing. The above methods are based on the fusion-solidification principle for printing titanium alloys. The high-temperature process is usually associated with some defects, such as metallurgical defects, tensile residual stresses, cracks, and deformation. Compared with the melting process, CSAM has the advantages of shortening production time, an unlimited product size, and high flexibility. Meanwhile, the sprayed particles of CSAM are characterized by low temperature, high speed, and low oxidation, thus having attracted increasing attention in the industry [99,100]. When applied to print titanium alloys, however, this method still faces significant challenges. Because of its high melting point and high strength, it is challenging for titanium base particles to produce thermoplastic deformation in the cold spraying process. As a result, the critical rate of successful deposition of Ti and its alloys is generally higher than that of most metals (e.g., over 1000 m/s for Ti6Al4V) [101,102], which also results in cold spray titania-based deposits that are often porous, as shown in Figure 7. The current state of cold spray additive manufacturing (CSAM) for printing titanium alloys shows promising advancements in various industries, such as aerospace, automotive, and healthcare. CSAM has demonstrated its capability to produce high-quality titanium alloy components with improved mechanical properties and complex geometries, paving the way for potential applications in lightweight structures, wear-resistant parts, and medical implants. However, further research and development are required to optimize process parameters, enhance material properties, and ensure reliable and cost-effective production in industrial settings.
To solve the poor mechanical properties caused by the porosity of titanium-base sediments, many scholars have begun to study the mechanism of cold spray and post-treatment from two aspects, and the post-treatment research will be elaborated on in the next chapter. Brit et al. [104] used FIB+TEM technology to study the cohesive region in the cold spray Ti6Al4V sediments with nitrogen and nitrogen −73% helium as propulsion gases. As shown in Figure 8, it was found that two regions could be identified in the sedimentary particles. Smooth areas are transformations near particle-particle boundaries, while textured areas have broken martensite lathes and elongated or twisted particles. Lek et al. [105] observed the effect of Ti6Al4V particles on Ti6Al4V substrate using FIB + TEM technology. Observed under a transmission electron microscope at a high power, the narrow gap (or void) at the interface formed by the rebound of particles at the impact center is about 60 nm wide. Near the binding zone, α-phase nanoparticles (less than 50 nm) were formed under the strong thermal-mechanical coupling action during the impact, while the particles near the central zone were about 50–200 nm. At the same time, it was found that the subsequent particles will lead to a certain deformation of the deposited particles, which has a certain tamping effect on the impact center.
It can be inferred that in the process of CSAM manufacturing, as shown in Figure 9, the impact zone, bonding zone, and particle edge will be formed when particles impact the substrate. The bonding starts from the edge of the contact zone between the deformed particles and the substrate, and the impact center (impact zone) of titanium alloy particles will rebound to form pores. The void generated by the particle impact center is also smaller, but the center of the contact area will not be combined, and no void can be achieved [106,107]. On this basis, it can be concluded that the porosity of CSAM printing titanium alloys can be effectively reduced by increasing the particle injection velocity. This conclusion lays a foundation for the subsequent research on the performance improvement of CSAM printing titanium alloys.
SLM, EBM, and WAAM have shown promise in printing titanium alloys that meet industrial standards. Figure 10 provides an overview of these three printing methods. It is noteworthy that SLM and EBM stand out in achieving high density through the localized melting and solidification of metal powders. On the other hand, WAAM demonstrates a relatively lower density owing to its wire-based deposition approach. When it comes to print size, SLM and EBM are well-suited for creating small to medium-sized parts, whereas WAAM offers an advantage in producing large-scale components. Precision in print accuracy is a strength of SLM and EBM, making them suitable for fabricating intricate and detailed parts. Nevertheless, WAAM exhibits lower accuracy due to factors such as wire deposition and heat transfer. Comparatively, SLM and EBM ensure superior surface quality, benefiting from their ability to control energy density and melting. Conversely, surfaces produced by WAAM tend to be rougher due to larger deposition sizes. Residual stress, a key factor influencing mechanical properties, is lower in SLM and EBM owing to their localized melting and rapid solidification. Conversely, WAAM often experiences higher residual stress levels, necessitating additional post-processing procedures. Finally, it is important to consider the costs associated with each technique. SLM and EBM generally involve higher material costs due to the use of specific powders and advanced equipment. In contrast, WAAM is more cost-effective because it utilizes wire feedstock. To summarize, SLM, EBM, and WAAM each possess distinct characteristics. SLM and EBM excel in achieving high density, precision, and surface quality, albeit at a higher cost. On the other hand, WAAM offers advantages in terms of print size capability and cost-effectiveness, albeit with sacrifices in accuracy and surface quality. The decision regarding the appropriate printing technique should be based on the specific requirements and constraints of the application at hand.
Figure 11 presents the mechanical properties of the three methods and cast-forged titanium alloys, including ultimate tensile strength (UTS) and elongation (EL). As can be seen from Figure 11, SLM can print titanium alloys with the highest ultimate tensile strength, but the lowest ductility. In addition, the unique titanium alloys designed by in situ alloying have great potentials for achieving better mechanical properties; EBM-printed titanium alloys harvest the highest ductility while taking into account the high ultimate tensile strength; The titanium alloys printed by WAAM can achieve higher ductility but lower ultimate tensile strength after parameter optimization. At the same time, the layered design and microelement doping of WAAMed titanium alloy show the possibility of improving ductility and ultimate tensile strength, respectively. As shown in Table 5, the CSAM method is not yet able to print titanium alloys with performance up to industry standards, and there is still a lot of room for improvement.
For SLM and EBM, the current research direction should focus on the use of in situ alloying to modulate microstructures and the design method of titanium alloys to achieve further breakthroughs in mechanical properties. The WAAM process can further exert the advantage of low costs to realize industrial one-step production and the design of layered materials doped with trace elements as the future development direction; the printing mechanism of CSAM has not been investigated thoroughly enough, thus remaining to be further explored.

3. Post-Processing

For the layer-by-layer printing method, the surface quality of printed parts is generally low due to the layer-processing and step effect [111,112]. During the manufacturing process, porosity, cracks, powder agglomeration, thermal stress, and other internal defects occur between different printing layers. These defects seriously affect the internal microstructure and mechanical properties of the final components [113,114,115,116,117]. However, CSAM printed titanium alloys do not easily form pore-free microstructures with almost no residual stress [118]. AM technology alone cannot simultaneously manufacture parts that meet the requirements of mechanical properties and surface roughness [119,120]. Therefore, post-processing operations are usually required to improve the surface quality and mechanical properties of parts after the completion of manufacturing [121,122]. In this study, the most common post-heat treatment, hot isostatic pressing, inter-pass cooling, and shot peening technologies were selected, their effects were summarized, and the AM technology suitable for this method was analyzed and matched.

3.1. Post-Heat Treatment

Post-heat treatment (PSHT) is usually used to change the composition and microstructure of elements by heating, holding, and cooling in the solid-state range. It can improve ductility, eliminate brittle phases, and refine grains.
The effect of PSHT on the microstructure and mechanical properties of additive titanium alloys has been investigated. Cai et al. [36] studied the effect of heat treatment temperature on the microstructure and mechanical properties of SLMed TA15 alloys. As shown in Figure 12, with the increase of heat treatment temperature, the α’ phase is gradually decomposed into strip α + β phase, the β phase content gradually increases, and the grains are coarse, forming a coarse-layered structure. The column β grains contribute to good ductility but may lead to strength reduction [55]. The α’ martensite phase is a hard and brittle non-equilibrium phase, which can improve the tensile strength of the component, but is harmful to the ductility. The decomposition of α’ martensite and the increase of β phase content reduce the ultimate tensile strength and improve the ductility of the sample. Such results are consistent with the conclusion of Wang et al. [123] using the double annealing process. In addition, Figure 13 summarizes the effect of post-heat treatment on residual stress, and it can be seen that PSHT can effectively reduce residual stress. Despite the better mechanical properties of titanium alloys, multiple heat treatments show low economic benefits and consume more production time and costs, so single-stage post-heat treatment is generally adopted in the industry.
For the CSAM process, unlike SLM and WAAM, post-heat treatment can promote adhesion in the sediment, resulting in increased tensile strength and ductility. After PSHT treatment at 600 °C for 2 h, the tensile strength of cold-sprayed Ti6Al4V sediments increased by 72% to 765 MPa when helium was used [126]. However, the tensile strength for sediments produced using nitrogen remained relatively low (462 MPa) even after 4h of PSHT treatment at 1000 °C due to the large porosity. These results indicate that PSHT can improve the tensile strength of cold-sprayed titanium alloy deposits with low porosity [126,127]. Therefore, PSHT reduces residual stress, eliminates martensitic SLM and WAAM, and promotes the adhesion of CSAM-printed titanium alloys.

3.2. Hot Isostatic Pressure

Hot isostatic pressing (HIP) is a typical thermomechanical process that combines high-temperature and high-pressure production techniques. In a high-temperature enclosed environment, high-pressure inert gas is used as the pressure medium to squeeze the manufactured parts uniformly in all directions.
HIP is usually chosen for post-processing to reduce printed parts porosity. The high-temperature environment of HIP can make the metal plastic lift, and extrusion under this environment can effectively eliminate the inherent defects and pores inside the metal, reduce the internal porosity, and homogenize the microstructure [128]. Internal defects are common in additive manufacturing titanium alloys, especially in PBF and CS processes. Hausmann et al. [129] reported that the porosity could still be maintained at 0.08% under optimized SLM parameters. Ackelid et al. [130] found that the porosity in EBM-manufactured parts was 0.17%, and the shape and direction of the pores strongly affected the macroscopic ductility [31]. However, the porosity cannot be reduced by heat treatment [131], while HIP positively affects the reduction of porosity (Figure 14 and Figure 15) [129,131,132,133]. For example, the pore volume fraction of Ti6Al4V after HIP decreases from 0.08% to 0.01% [129], and the pore size also decreases significantly [131]. In addition, HIP is usually implemented at very high temperatures, equivalent to a high-temperature annealing process, as shown in Figure 16. Therefore, hot isostatic pressure includes mechanical treatment, which can reduce porosity, and also includes heat treatment, which reduces the residual stress and decomposes the unstable martensitic phase through temperature control, resulting in a decrease in strength and an increase in ductility.

3.3. Inter-Pass Cooling

Cold rolling was initially used to release residual stress of materials through plastic deformation [136,137]. The working principle is shown in Figure 17, and AM was later introduced as a post-treatment process.
The primary purpose of cold rolling between layers is to solve the anisotropy of mechanical properties of printed parts caused by the columnar β grains produced by a thermal gradient. By plastic deformation of the sediments, the initially large columnar β grains are transformed into equiaxial crystals with sizes ranging from 56 to 139 μm by cold rolling technology (Figure 18 and Figure 19), which effectively refines the β grains and significantly reduces the anisotropy of the microstructure [139,140]. The main reason is that cold rolling can introduce stacking faults, point defects, dislocation, and twins and provide greater dislocation density and deformation energy. When subsequent upper layer heating causes the layer to remelt, these defects and twins provide a driving force for recrystallization, and new β orientations can outpace the growth of residual β, thus breaking the columnar β grains and leading to significant β grain refinement [141]. Meanwhile, the residual stress can be reduced or eliminated by cold rolling, which contributes to more uniform mechanical properties. Therefore, SLM, EBM, and WAAM are suitable for layer-by-layer printing and titanium alloys with columnar β grains.

3.4. Shot Peening

Shot peening is a cold-mechanical treatment method in which a high-energy medium impounds the sediment and removes the tensile stress by applying compressive stress to the treated surface.
Similar to the principle of inter-pass cooling, shot peening produces elastic deformation and large plastic deformation in the extrusion part of metal deposits, which can refine the stochastic orientation of surface grains and thus contribute to improving mechanical strength [143]. However, shot peening is limited by the depth of action, about 60 μm below the surface. Though being an excellent post-mechanical treatment, shot peening can only improve the surface material properties for SLM, EBM, and WAAM with a limited range of action, making it not as suitable as inter-pass cooling.
However, the high-speed impact of particles is a very basic and normal process in the CSAM process, which does not have special requirements on the depth of the effect of shot peening technology and avoids its limitations. In addition, the large-particle shot peening technology can effectively play a tamping role in tamping the deposited layer [144]. Based on the findings presented in Figure 20, it is evident that during the CSAM process, the particles undergo deformation upon impact with the substrate. However, under the influence of the shot peening process, the larger particles cause secondary deformation of the deformed particles, resulting in compaction. This compaction effect greatly reduces the porosity of the printed products. Luo et al. [145] first studied the effect of in situ shot peening on cold sprayed Ti and Ti6Al4V sediments and experimentally explored the effect of SP particles of different proportions (0, 10, and 70 vol.%) on porosity and microhardness. As depicted in Figure 21, the density of both Ti and Ti-6Al-4V deposits increases with an increase in the SP (Shot Peening) particle content, reaching its maximum value at the highest proportion of SP particles. This phenomenon can be attributed to the increased content of SP particles, which leads to a greater compaction effect and therefore an enhanced density of the products. The trends in product density and microhardness variations with increasing SP particle content can be observed in Figure 22. It can be seen that the porosity was inversely proportional to the SP particle content, and this effect was more obvious under low SP particle content. When the SP particle content increased to 70 vol.%, the porosity of Ti and Ti6Al4V sediments decreased to 0.3% and 0.7%, respectively. The microhardness of Ti and Ti6Al4V sediments increased to 203 HV and 427 HV, respectively. It can be inferred from these results that customized compact sedimentary beds with low porosity and high microhardness can be prepared by selecting an appropriate proportion of SP particles. In summary, shot peening is more suitable for the CSAM process, which can effectively reduce the porosity of printed titanium alloys and improve their microhardness.
As the most common post-processing method, PSHT has been widely used in the post-processing of AMed titanium alloys. Due to the decomposition of the α’ martensitic phase in the PSHT process and the subsequent increase of β phase content, the ultimate tensile strength of titanium alloys decreases and the ductility increases after PSHT (Figure 23). An appropriate PSHT temperature can optimize the microstructure of titanium alloys and achieve the optimal combination of ultimate tensile strength and ductility. At the same time, reheating can release residual stress and reduce the residual stress accumulated by rapid cooling in SLM, EBM, and WAAM printing processes. The ultimate tensile strength and ductility of CSAMed titanium alloys can be effectively improved by the better adhesion of CSAMed titanium alloys in the deposit after PSHT, as shown in Figure 24. Therefore, PSHT is generally applicable to various AM printing methods for titanium alloys.
The effect of HIP is equivalent to heat treatment and mechanical pressure. While the microstructure of titanium alloys is regulated by heating, the pores of AMed titanium alloys are mainly subjected to mechanical extrusion, and the ultimate tensile strength and ductility of titanium alloys are effectively improved by reducing the porosity of printed titanium alloys. HIP is suitable for a variety of printing methods due to the presence of pores in AMed titanium alloys. As shown in Figure 25, heat treatment reduces the ultimate tensile strength of SLMed Ti64 and improves the ductility, while HIP at 900 °C; further improves the ductility without further reducing the ultimate tensile strength. Therefore, HIP has a better effect than annealing. Due to the high porosity of CSAMed titanium alloys, the performance improvement of CSAMed titanium alloys is the most significant, which can improve the ultimate tensile strength and ductility simultaneously.
Inter-pass cooling and shot peening are matched to SLM, EBM, and WAAM for layer-by-layer printing and the CSAM method for powder jet deposition, respectively. Cold rolling between layers can refine the columnar β grains and reduce the porosity, thus reducing the anisotropy of titanium alloys and improving the mechanical properties of titanium alloys. However, the effect of shot peening on the layer-by-layer printing method is limited, but it is well-matched with the CSAM method. Large-particle shot peening can play a tamping role, tamping the deposited layer and effectively reducing the porosity of CSAMed titanium alloys, so as to improve their mechanical properties.

4. Conclusions and Perspectives

High-temperature titanium alloys can be widely used in the field of aerospace heat-resistant parts by virtue of their excellent mechanical properties such as low density, high strength, and good heat resistance. However, traditional casting and forging methods are disadvantaged in manufacturing complex-shaped titanium alloy components. Additive manufacturing has overcome this dilemma and achieved better mechanical properties of AMed titanium alloys, showing great prospects. As modernization proceeds, higher requirements are put forward for the properties of titanium alloys. Therefore, it has become imperative to explore strategies for enhancing the mechanical properties of AMed titanium alloys, which will enable titanium alloys to reach their full potentials in practical applications.
To fully utilize the value of additive manufacturing for high-temperature titanium alloys, it is essential to understand the characteristics and positioning of titanium alloys printed by each process. SLM is particularly suitable for producing small-sized high-precision titanium alloy parts. The microstructure of SLM-printed titanium alloys is typically characterized by numerous brittle and hard α’ phases, resulting in high ultimate tensile strength but low plasticity. EBM is optimal for the mass production of small-sized precision parts. EBMed titanium alloys often exhibit an α + β basket microstructure or a Widmanstätten microstructure. With high β content and low α’ phase content, EBMed titanium alloys display high ductility and ultimate tensile strength. WAAM is suitable for producing large-sized general-purpose titanium alloys with low requirements. The microstructure of WAAMed titanium alloys features diverse morphologies as influenced by the cooling rate. However, a larger grain size of WAAMed products results in lower strength compared to SLM and EBM printed products. CSAM is not yet capable of commercially producing titanium alloys. The microstructure of CSAMed titanium alloys is often characterized by broken martensitic laths and twisted phases. The damaged phase structure and high porosity lead to the unsatisfactory mechanical properties of CSAMed titanium alloys.
To better control the research direction, we need a deep understanding of the performance improvement strategy for titanium alloys printed by each process. Mechanism study, process parameter optimization, in situ micronutrient addition, and layered material design are four different research methods to improve the mechanical properties of post-treated AMed titanium alloys. Currently, SLM is primarily conducted via in situ addition of trace elements and has successfully generated several unique titanium alloys that exhibit enhanced performance. In the future, the scope of element addition can be continuously expanded for the layered material design of the existing titanium alloys. Meanwhile, the research on the optimization of process parameters for EBM has gradually reached saturation. In subsequent research, in situ element addition can be explored to design unique titanium alloys based on the examples of SLM. The titanium alloys produced via the WAAM process fail to achieve excellent performance compared with those of SLM, which may be attributed to the immature optimization of WAAM process parameters. Therefore, it is necessary to further standardize process parameters for industrial production in the future. The CSAM process is still in its infancy, and the research on its printing mechanism should be strengthened to overcome the challenge of high porosity in printed products.
By summarizing the abovementioned methods and comparing post-treatment processes, we can improve the mechanical properties of additive manufacturing for high-temperature titanium alloys and boost the further development of the aerospace industry and other modern industries. PSHT is capable of releasing residual stress and refining the microstructure, while HIP can effectively reduce porosity. Both methods are capable of significantly improving the mechanical properties of AMed titanium alloys and are generally applicable to all four additive manufacturing processes. Inter-pass cooling can refine columnar β grains and improve ductility. Due to the incompatibility with the CSAM printing principle, however, inter-pass cooling is only applicable to SLM, EBM, and WAAM processes. Shot peening can refine the surface grain orientation and improve mechanical strength for SLM, EBM, and WAAM processes, but its effectiveness is limited. In contrast, shot peening in combination with large particle doping and uniform tamping can effectively improve hardness and reduce the porosity in CSAM processes.

Author Contributions

Conceptualization, B.J., Q.W. and Y.S.; investigation, B.J., W.G. and Q.W.; resources, B.J. and Y.S.; wtiting—original draft preparation, B.J.; writing—review and editing; B.J., Y.S., L.Z., A.P., X.D. and X.Z. All authors have read and agreed to the published version of the manuscript.

Funding

This project is supported by the Natural Science Foundation of Zhejiang Province (grant no. 2021C01023).

Data Availability Statement

No new data were created or analyzed in this study. Data sharing is not applicable to this article.

Acknowledgments

This project is supported by the Natural Science Foundation of Zhejiang Province (grant no. 2021C01023). The authors would like to acknowledge support from the Natural Science Foundation of Zhejiang Province.

Conflicts of Interest

The authors declare that there are no conflicts of interest regarding the publication of this paper.

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Figure 1. Working principle of: (a) SLM [22] (b) EBM [23] (c) WAAM [24], and (d) CSAM [25].
Figure 1. Working principle of: (a) SLM [22] (b) EBM [23] (c) WAAM [24], and (d) CSAM [25].
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Figure 2. Comparison of microstructures and tensile properties of Ti–6Al–4V and the newly developed alloy (25Ti–0.25O) fabricated via laser power bed fusion (L-PBF). (a) Schematic of the L-PBF process and the intrinsic thermal cycles that different locations of the fabricated part undergo. (b) Scanning electron microscopy (SEM)-backscattered electrons (BSE) micrographs showing spatially dependent phases in Ti–6Al–4V along the building direction (BD). Note that the top surface is predominantly composed of acicular α′ martensite. The lower region shows a partial decomposition of α′ martensite due to more thermal cycles. It can be seen that α′ martensite, α phase, and thin β film are presented, as marked with white arrows. The bottom region exhibits a well-defined lamellar (α + β) microstructure. (c) Tensile engineering stress-strain curves of Ti–6Al–4V along the vertical and horizontal directions. Inset, schematic of preparation of vertical and horizontal tensile specimens from the as-built parts. The horizontal tensile specimens are marked from H1 to H6 along the building direction. (d) SEM-BSE micrographs showing homogeneous lamellar (α + β) microstructure in the newly developed 25Ti–0.25O alloy. The well-defined lamellar (α + β) microstructure can be observed from the top surface to the bottom region, and (e) Tensile engineering stressstrain curves of 25Ti–0.25O alloy along the vertical and horizontal directions. Inset, the preparation of the tensile specimens is the same as that of Ti–6Al–4V [35].
Figure 2. Comparison of microstructures and tensile properties of Ti–6Al–4V and the newly developed alloy (25Ti–0.25O) fabricated via laser power bed fusion (L-PBF). (a) Schematic of the L-PBF process and the intrinsic thermal cycles that different locations of the fabricated part undergo. (b) Scanning electron microscopy (SEM)-backscattered electrons (BSE) micrographs showing spatially dependent phases in Ti–6Al–4V along the building direction (BD). Note that the top surface is predominantly composed of acicular α′ martensite. The lower region shows a partial decomposition of α′ martensite due to more thermal cycles. It can be seen that α′ martensite, α phase, and thin β film are presented, as marked with white arrows. The bottom region exhibits a well-defined lamellar (α + β) microstructure. (c) Tensile engineering stress-strain curves of Ti–6Al–4V along the vertical and horizontal directions. Inset, schematic of preparation of vertical and horizontal tensile specimens from the as-built parts. The horizontal tensile specimens are marked from H1 to H6 along the building direction. (d) SEM-BSE micrographs showing homogeneous lamellar (α + β) microstructure in the newly developed 25Ti–0.25O alloy. The well-defined lamellar (α + β) microstructure can be observed from the top surface to the bottom region, and (e) Tensile engineering stressstrain curves of 25Ti–0.25O alloy along the vertical and horizontal directions. Inset, the preparation of the tensile specimens is the same as that of Ti–6Al–4V [35].
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Figure 3. The microstructures of EBM Ti–6Al–4V solid samples: (a–c) α + β lamellar and Widmanstätten structures [62]; (d) columnar β microstructure [63].
Figure 3. The microstructures of EBM Ti–6Al–4V solid samples: (a–c) α + β lamellar and Widmanstätten structures [62]; (d) columnar β microstructure [63].
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Figure 4. EBSD maps for the as-fabricated EBAM Ti-6Al-4V samples: (a) image quality maps, where the light values indicate high image quality; (b) EBSD orientation maps from different samples; (c) the grain boundary map [64].
Figure 4. EBSD maps for the as-fabricated EBAM Ti-6Al-4V samples: (a) image quality maps, where the light values indicate high image quality; (b) EBSD orientation maps from different samples; (c) the grain boundary map [64].
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Figure 5. (a–d) Microstructure of the as-deposited Ti-6Al-4V by WAAM [76].
Figure 5. (a–d) Microstructure of the as-deposited Ti-6Al-4V by WAAM [76].
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Figure 6. The microstructure under different cooling processes: (a) martensite α’[85], (b) needle-like secondary α [82], (c) martensite α” [83], (d) grain boundary α + fine basket-weave structure (Widmanstätten α) [81], and (e) acicular α[76].
Figure 6. The microstructure under different cooling processes: (a) martensite α’[85], (b) needle-like secondary α [82], (c) martensite α” [83], (d) grain boundary α + fine basket-weave structure (Widmanstätten α) [81], and (e) acicular α[76].
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Figure 7. Typical SEM microstructures of Ti (a,c) and Ti6Al4V (b,d) deposits in the etched state [103].
Figure 7. Typical SEM microstructures of Ti (a,c) and Ti6Al4V (b,d) deposits in the etched state [103].
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Figure 8. Microstructure characterization of the cold-sprayed Ti6Al4V deposits with nitrogen at 790 °C and 3.6 MPa: (a) OM image of an etched cross-section showing small bands of the smooth transformation region; (b) SEM-BEI of a cross-section showing in detail the smooth region and the textured region; (c) TEM image of sample S1 taken from the particle interface zone showing a disorganized and broken structure (R1), an organized and stacked lamellar structure (R2), and a fine-grained structure (R3) [103].
Figure 8. Microstructure characterization of the cold-sprayed Ti6Al4V deposits with nitrogen at 790 °C and 3.6 MPa: (a) OM image of an etched cross-section showing small bands of the smooth transformation region; (b) SEM-BEI of a cross-section showing in detail the smooth region and the textured region; (c) TEM image of sample S1 taken from the particle interface zone showing a disorganized and broken structure (R1), an organized and stacked lamellar structure (R2), and a fine-grained structure (R3) [103].
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Figure 9. Cross-sectional TEM image of a Ti splat and steel substrate comprising several low magnification scans. The boxed area and the arrow indicate the bonded region and the non-bonded region, respectively [103].
Figure 9. Cross-sectional TEM image of a Ti splat and steel substrate comprising several low magnification scans. The boxed area and the arrow indicate the bonded region and the non-bonded region, respectively [103].
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Figure 10. The characteristics of SLM, EBM, and WAAM processes [108,109].
Figure 10. The characteristics of SLM, EBM, and WAAM processes [108,109].
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Figure 11. Mechanical properties of untreated titanium alloys manufactured by SLM, EBM, and WAAM.
Figure 11. Mechanical properties of untreated titanium alloys manufactured by SLM, EBM, and WAAM.
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Figure 12. Microstructures of untreated samples and heat-treated TA15 in the partial martensitic decomposition temperature range: (a,b) untreated, (c,d) 2 h at 650 °C and (e,f) 2 h at 750 °C. White and yellow arrows indicate needle-like structures and lath-like structures, respectively [36].
Figure 12. Microstructures of untreated samples and heat-treated TA15 in the partial martensitic decomposition temperature range: (a,b) untreated, (c,d) 2 h at 650 °C and (e,f) 2 h at 750 °C. White and yellow arrows indicate needle-like structures and lath-like structures, respectively [36].
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Figure 13. Effect of post-heat treatment on residual stress: (a) SLM, (b) EBM [124,125].
Figure 13. Effect of post-heat treatment on residual stress: (a) SLM, (b) EBM [124,125].
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Figure 14. Optical micrographs presenting defects. (a) EBM parts with obvious defects. (b) Optical microstructure after HIP [134].
Figure 14. Optical micrographs presenting defects. (a) EBM parts with obvious defects. (b) Optical microstructure after HIP [134].
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Figure 15. 3D-reconstruction of the pores within the CS Ti6Al4V samples under different conditions: (a) N2-AF, (b) N2-HIP, (c) He-AF, and (d) He-HIP; the color scale bar corresponds to the equivalent diameter of the reconstructed pores [135].
Figure 15. 3D-reconstruction of the pores within the CS Ti6Al4V samples under different conditions: (a) N2-AF, (b) N2-HIP, (c) He-AF, and (d) He-HIP; the color scale bar corresponds to the equivalent diameter of the reconstructed pores [135].
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Figure 16. SEM images of (a) structures of TiAl6V4 alloy (b) an effect of 900 ℃ heat (c) HIP treatment on the microstructure of the SLM samples [129].
Figure 16. SEM images of (a) structures of TiAl6V4 alloy (b) an effect of 900 ℃ heat (c) HIP treatment on the microstructure of the SLM samples [129].
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Figure 17. Schematic of the rolling rig used for the application of cold working between successive layers [138].
Figure 17. Schematic of the rolling rig used for the application of cold working between successive layers [138].
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Figure 18. Microstructure of the cruciform intersection (a,b) without rolling and (c,d) with rolling [142].
Figure 18. Microstructure of the cruciform intersection (a,b) without rolling and (c,d) with rolling [142].
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Figure 19. EBSD images of the columnar as-deposited region (a) and the recrystallized grains (b) of the WAAMed sample. (c) Legend with regard to color and orientation [138].
Figure 19. EBSD images of the columnar as-deposited region (a) and the recrystallized grains (b) of the WAAMed sample. (c) Legend with regard to color and orientation [138].
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Figure 20. Schematic diagram of in situ shot peening assisted cold spray for fabricating Ti6Al4V deposits [146].
Figure 20. Schematic diagram of in situ shot peening assisted cold spray for fabricating Ti6Al4V deposits [146].
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Figure 21. Cross-sectional microstructures of Ti and Ti6Al4V deposits cold sprayed with different ratios of shot-peening (SP) particles [145].
Figure 21. Cross-sectional microstructures of Ti and Ti6Al4V deposits cold sprayed with different ratios of shot-peening (SP) particles [145].
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Figure 22. Porosity and microhardness of Ti and Ti6Al4V deposits as a function of SP particles content [145].
Figure 22. Porosity and microhardness of Ti and Ti6Al4V deposits as a function of SP particles content [145].
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Figure 23. Comparison diagram of room-temperature and high-temperature tensile properties for the samples without and with heat treatments. The blue frames show the corresponding tensile testing temperatures. The five-pointed stars and squares represent tensile strength and elongation, respectively [36].
Figure 23. Comparison diagram of room-temperature and high-temperature tensile properties for the samples without and with heat treatments. The blue frames show the corresponding tensile testing temperatures. The five-pointed stars and squares represent tensile strength and elongation, respectively [36].
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Figure 24. Typical stress–strain curves for the Ti6Al4V substrate, helium-sprayed Ti6Al4V deposits (as-sprayed and annealed at 600 °C for 2 h), and nitrogen-sprayed Ti6Al4V deposits (as-sprayed and annealed at 1000 °C for 4 h) [103].
Figure 24. Typical stress–strain curves for the Ti6Al4V substrate, helium-sprayed Ti6Al4V deposits (as-sprayed and annealed at 600 °C for 2 h), and nitrogen-sprayed Ti6Al4V deposits (as-sprayed and annealed at 1000 °C for 4 h) [103].
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Figure 25. (a) Stress–strain curves for some TiAl6V4 specimens produced by SLM with and without further thermomechanical treatment; (b,c) variation of UTS and elongations of CS Ti6Al4V samples under different conditions [129,135].
Figure 25. (a) Stress–strain curves for some TiAl6V4 specimens produced by SLM with and without further thermomechanical treatment; (b,c) variation of UTS and elongations of CS Ti6Al4V samples under different conditions [129,135].
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Table
Table 1. Mechanical properties of titanium alloy fabricated by SLM.
Table 1. Mechanical properties of titanium alloy fabricated by SLM.
MaterialUTS (MPa)YS (MPa)EL (%)Ref.
Ti–6Al–4V1211 ± 311100 ± 126.5 ± 0.6[17]
1267 ± 51110 ± 97.3 ± 1.1
1149 ± 111093 ± 1511.3 ± 0.5[29]
1090 ± 101022 ± 1012.7 ± 2.1
1165 ± 21112 ± 311.6 ± 1.2
1095 ± 10990 ± 58.1 ± 0.3[30]
1421 ± 1201273 ± 533.2 ± 0.5[22]
1166 ± 25962 ± 471.7 ± 0.3[31]
1206 ± 81137 ± 207.6 ± 2
1219 ± 201143 ± 304.89 ± 0.6[32]
1269 ± 91195 ± 195 ± 0.5
125011256[33]
140713334.54
Ti64–(4.5%)316L1297 ± 10984 ± 148.8 ± 0.2[34]
Ti64–25Ti–0.5O\1220.8 ± 6.513.7 ± 0.9[35]
TA151422.2 ± 50.2\9.5 ± 1.5[36]
1362.6 ± 46.2\9.3 ± 0.4
1349.8 ± 26.7\5.6 ± 0.1
1234.2 ± 53.1\7.3 ± 0.7
Table
Table 2. Tensile properties at room temperature and high temperature of printed TA15 alloy samples under different SLM processing parameters [36].
Table 2. Tensile properties at room temperature and high temperature of printed TA15 alloy samples under different SLM processing parameters [36].
S(mm/s)P(W)Room Temperature500 °C
UTS (MPa)EL (%)UTS (MPa)EL (%)
6752301422.1 ± 50.29.5 ± 1.5990.2 ± 2.215.6 ± 1.9
2801362.6 ± 46.29.3 ± 0.4949.6 ± 19.613.6 ± 3.2
3301349.8 ± 26.75.6 ± 0.1909.4 ± 40.112.5 ± 0.4
3801063.6 ± 136.97.5 ± 0.5939.5 ± 20.715.3 ± 1.7
Table
Table 3. Mechanical properties of titanium alloy fabricated by EBM.
Table 3. Mechanical properties of titanium alloy fabricated by EBM.
MaterialUTS (MPa)YS (MPa)EL (%)Ref.
Ti–6Al–4V950–990910–94014–16[65]
1150–12001100–115016–25
994–1029883–93811.6–13.6
93086512–17
944.5–964.5823.4–851.813.2–16.3
≈990–1180≈900–1100≈18–23
91583013.1[30]
1012 ± 3962 ± 48.8 ± 1.6[66]
1011 ± 4947 ± 119 ± 1.1
790 ± 10740 ± 102.2 ± 0.3[67]
833 ± 22783 ± 152.7 ± 0.4
851 ± 19812 ± 123.6 ± 0.9
1029.7 ± 7982.9 ± 5.712.2 ± 0.8[60]
1032.9 ± 12.9984.1 ± 8.59 ± 2.9
978 ± 3.2899 ± 4.79.5 ± 1.2[32]
928 ± 9.8869 ± 7.29.9 ± 1.7
Ti62421059 ± 19957 ± 1613.8 ± 0.02[68]
1103 ± 29896 ± 2115.3 ± 0.01
1048 ± 7937 ± 1916.4 ± 0.03
1000 ± 4887 ± 2116.2 ± 0.01
Table
Table 4. Mechanical properties of titanium alloy fabricated by WAAM.
Table 4. Mechanical properties of titanium alloy fabricated by WAAM.
MaterialUTS (MPa)YS (MPa)EL (%)Ref.
Ti–6Al–4V929 ± 41\9 ± 1.2[74]
939 ± 24\16 ± 3[75]
918 ± 17803 ± 1514.8[71]
937 ± 21861 ± 1416.5 ± 2.7[76]
820 ± 6.23710 ± 47.18 ± 0.93
988 ± 19.2909 ± 13.67 ± 0.5
902–923881–9066.4[77]
1166.3 ± 22.11059.1 ± 18.45.97 ± 0.65[78]
Ti64 + Ti64 − 0.07B988.4\11[79]
980.2\11.7
895.2\12.2
886.4\14.9
Ti64–CoCrNi1336.2 ± 7.81260 ± 4.45.08 ± 0.63[78]
Ti64 + Ti64 − CoCrNi1292.1 ± 9.41172 ± 5.15.44 ± 0.3[78]
Ti64–0.9Ni1350 ± 101309.55 ± 0.2[80]
Table
Table 5. Tensile properties of CSAMed Ti64 alloy samples at different propulsion gases [110].
Table 5. Tensile properties of CSAMed Ti64 alloy samples at different propulsion gases [110].
MaterialPropulsion GasUTS (MPa)EL (%)
Ti-6Al-4VHe4803
N21501.5
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MDPI and ACS Style

Jin, B.; Wang, Q.; Zhao, L.; Pan, A.; Ding, X.; Gao, W.; Song, Y.; Zhang, X. A Review of Additive Manufacturing Techniques and Post-Processing for High-Temperature Titanium Alloys. Metals 2023, 13, 1327. https://doi.org/10.3390/met13081327

AMA Style

Jin B, Wang Q, Zhao L, Pan A, Ding X, Gao W, Song Y, Zhang X. A Review of Additive Manufacturing Techniques and Post-Processing for High-Temperature Titanium Alloys. Metals. 2023; 13(8):1327. https://doi.org/10.3390/met13081327

Chicago/Turabian Style

Jin, Binquan, Qing Wang, Lizhong Zhao, Anjian Pan, Xuefeng Ding, Wei Gao, Yufeng Song, and Xuefeng Zhang. 2023. "A Review of Additive Manufacturing Techniques and Post-Processing for High-Temperature Titanium Alloys" Metals 13, no. 8: 1327. https://doi.org/10.3390/met13081327

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