Lattice structures have excellent mechanical properties such as high specific strength and high specific rigidity as well as outstanding functional characteristics such as vibration reduction, heat dissipation, sound absorption, and electromagnetic shielding. They are widely used in aerospace, biomedicine, and transportation fields. However, the materials used to form lattice structures are mostly stainless steel, Ti6Al4V and AlSi10Mg, which cannot meet the requirements of some complex intelligent components for shapes, performances and functions that can change over time or space. The NiTi shape memory alloy, a new type of smart materials, has excellent super-elasticity, shape memory effect, excellent corrosion resistance and wear resistance, and other functional characteristics. It can achieve a shape recovery through certain external stimuli after deformation, so that it meets the requirements of controllable deformation and regulable performance. Compared with NiTi bulk materials, NiTi lattice structures have low elastic modulus and density, large deformation ability, and can adjust the mechanical properties by designing the size, shape and distribution of holes. The unique performance of NiTi lattice structures makes them have a wide range of application prospects in the aerospace field. For aerospace components, weight reduction is an eternal theme, but it is not clear how to further improve the lightweight characteristics of NiTi lattice structures. In this paper, the NiTi alloy body-centered tetragonal (BCT) hollow lattice structure is proposed, which possesses the advantages of high load-bearing capacity of the BCT lattice structure and the excellent super-elasticity and shape memory effect of NiTi alloys. That is, keeping the outer diameter of the strut constant and hollowing out the inner part of the strut are used to achieve the purpose of reducing the weight of the structure.

In this paper, NiTi pre-alloyed powder is used as the raw material to prepare the BCT hollow lattice structures by laser powder bed fusion (LPBF). The surface morphologies and microstructures of the formed samples are observed by scanning electron microscope (SEM), the phase transition behavior of BCT-100 is characterized by differential scanning calorimeter (DSC), and the phase compositions of BCT-100 are determined by X-ray diffractometer (XRD). The finite element simulation method and the uniaxial compression experiments are applied to analyze the influence of mass coefficient on the compression performance of structures. Cyclic compression-thermal recovery experiments are carried out to reveal the influence mechanism of mass coefficient on the shape memory effect of NiTi lattice structures.

The lattice structures manufactured by LPBF have high forming accuracy and relative density, and no defects such as cracks and irregular large-size holes are found (Fig. 4). The lattice structure with a mass coefficient of 100% has the best bearing capacity, the first maximum compressive force can reach 191.73 kN, and the corresponding deformation rate is 0.22. When the mass coefficient of the structure is reduced to 75% from 100%, the first maximum compressive force is 89.80 kN, and the bearing capacity is reduced by 53.16%. At that time, the compression deformation capacity is not weakened, and the deformation rate is still up to 0.21 (Fig. 7). Therefore, in the non-primary load-bearing members, the struts of lattice structures can be hollowed to reduce the mass coefficient of structures by 25% while maintaining the deformation capacity of components. However, further reducing structural mass coefficient weakens the bearing and deformation capacity. The shape memory effect of the lattice structure with a mass coefficient of 75% is the best, and the recovery rate can reach 98.92% in the first cycle (Fig. 9).

The forming quality of lattice structures fabricated by LPBF is high, but there are still dimensional deviations caused by solidification shrinkage, powder sticking, and staircase effect. The M_s (Martensite transformation start temperature) and A_f (Austenite transformation finish temperature) of BCT-100 are 11.02 ℃ and 33.72 ℃, respectively. The phases of components are mainly composed of B2 and B19′ at room temperature, and the B2 phase occupies the dominant position. The appearance of B19′ phase in components is related to the phase transition of B2 phase induced by thermal stress produced by LPBF. The experimental compression force-deformation rate curves of the four structures can be divided into five stages: elastic deformation of austenite phase, stress induced transformation of austenite phase into martensite phase, elastic deformation of martensite phase, plastic deformation of martensite phase, and fracture stage. The lattice structure with a mass coefficient of 100% has the best bearing capacity, the first maximum compressive force can reach 191.73 kN, and the corresponding deformation rate is 0.22. When the mass coefficient of the structure is reduced to 75% from 100%, the first maximum compressive force is 89.80 kN, and the bearing capacity is reduced by 53.16%. At that time, the compression deformation capacity is not weakened, and the deformation rate is still up to 0.21. When the mass coefficient is further reduced to 50% (BCT-50), the first maximum compressive load and deformation capacity are reduced by 81.52% and 36.36%, respectively, that are 35.43 kN and 0.14. The shape memory effect of lattice structures formed by LPBF is good. In the first cycle, BCT-75 has the best shape memory effect and the highest recovery rate can reach 98.92%. The recoverable rates of BCT-93 and BCT-100 are slightly low, which are 97.71% and 96.77%, respectively. The shape memory effect of BCT-50 is the worst and the recovery rate is only 94.94%. In the last two cycles, all components can achieve a fully recovery.

In the present study, the A131 EH36/AISI 1045 bimetallic structure is successfully fabricated by the hybrid DED and conventional processes. At the interface of the bimetallic structure, a transition zone of about 0.5 mm wide with good metallurgical quality is obtained without large cracks and unfused defects. The interface consists of microstructural refinement zones, coarsening zones, dual heat-affected zones, and heat-affected zones. The hardness in the interfacial region increases gradually along the building direction. The tensile strength, yield strength, and elongation of the as-built bimetallic structure are (629.0±1.1) MPa, (471.4±9.2)MPa, and 17.9%, respectively, which increase slightly to (671.3±5.6) MPa, (572.8±8.4)MPa and 22.1% after heat treatment. The as-built bimetallic structure is easier to cut than the heat-treated counterpart. When cutting from the AISI 1045 to the A131 EH36 regions, the cutting force decreases significantly with the maximum reduction of 64.1%. In addition, the surface roughness of the ultra-precision machining face decreases from(111.8±13.6) nm in the AISI 1045 region to (107.0±10.4) nm in the A131 EH36 region.

Considering data quantity, reliability, cost and other factors, it is an effective method to monitor the additive process of powder bed fusion-laser (PBF-L) by using photodiode to collect radiation information of molten pools in engineering practice. The on-line radiation monitoring of molten pools based on photodiode can collect a large amount of information in real time reflecting the internal conditions of molten pools in real time, and thus plays a very important role in predicting and controlling the quality of molten pools. The key is to continuously collect the radiation information of molten pools and to extract and analyze the data characteristics in a wide time scale (from several seconds to tens of hours) by the statistical methods, so as to realize the stable state analysis and quality prediction of the forming process. In PBF-L, all process factors (laser parameters, scanning strategy, powder state, and air flow protection) eventually influence the change of radiation signals of the molten pool, among which as for the main effects of laser power, scanning speed, preheating temperature and others on temperature and radiation intensity of a molten pool, a lot of research has been done. However, in addition to radiation intensity, the radiation signal change of the molten pool over time also contains a lot of other technological information and process stability information and it is worth further digging and studying. In the actual production process, PBF-L is a very complex process and the forming quality is still uneven even when the same equipment and process parameters are monitored and controlled. In order to ensure the repeatability and process consistency of additive manufacturing, the influence of process factors such as floor height, substrate position and scan line angle on the radiation signals of the molten pool is analyzed, and it is found that these characteristics can be used as an important indicator of quality control in the future.

In this paper, the forming experiment of K438 superalloy powder is carried out, and the photodiode is used to collect the radiation signal of the molten pool in the process of PBF-L. First it is analyzed and pretreated. Then it is segmented corresponding to the sample one by one. Statistical methods are used to process the segmented data, and the mean and standard deviations are selected as indicators to evaluate the signal characteristics, and the influence of process factors such as floor height, substrate position and scan line angle on the radiation signals of the molten pool is finally analyzed.

The influence of the process factors of PBF-L on the radiation intensity of the molten pool can reflect some classical laws. With the increase of layer height, the mean radiation intensity of the molten pool shows an overall increase trend, and the mean radiation intensity of the printing layer No.190 increases by about 6% (Fig. 4). In the wind field, the intensity waveform of the upwind molten pool shows the feature of "right deviation" , while that of the downwind molten pool shows the feature of "left deviation" (Fig. 6). The influence of relative position in the wind field on the radiation intensity of the molten pool is that the farther away from the air outlet, the greater the molten pool strength, and the closer to the air outlet, the smaller the molten pool strength (Fig. 8). In the scanning strategy, the scan line angle is exactly consistent with the incline angle of the sample (based on the horizontal direction), in which the mean intensity is the maximum (Fig. 11).

Through the processing and analysis of the intensity signal data of the molten pool, the influence law of some process factors on the radiation intensity of the molten pool is obtained, and the relationship between the radiation intensity law of the molten pool and the physical mechanism is established. The analysis shows that the effect of layer height on the radiation intensity of the molten pool is mainly due to heat accumulation among layers. In the wind field, the influence of wind direction and relative position on the radiation intensity of the molten pool is mainly the smoke masking effect. The effect of scan line angle on the radiation intensity of the molten pool is mainly due to the laser duty cycle. Through establishing the correlation between law and physical mechanism, it is clearer that in the practical engineering application and monitoring closed-loop control, it is necessary to consider the influence of layer height, scanning strategy, wind field conditions and other factors on the radiation intensity of molten pools. And its typical rules show that these characteristics can be used as an important indicator of quality control in the future.

The process characteristics of directed energy deposition melting point by point and stacking layer by layer determining the core process parameters of laser additive manufacturing, such as laser power, scanning velocity, powder feeding rate, layer thickness, overlap rate, and scanning strategy, have a decisive impact on the microstructures and mechanical properties of the titanium alloy components fabricated by laser additive manufacturing, which finally decide whether the components could meet the performance requirements of engineering applications. This work takes the near-β TC17 high strength titanium alloy widely used in advanced aero-engine blisk as the research object. In the process of laser additive manufacturing, different positions of components inevitably experience different interlayer cooling time, which results in the difference of the temperature field distribution and the thermal cycle history. All these complex factors affect the structures and property characteristics of as-deposited components, even the microstructural evolution undergoes the subsequent heat treatment. However, as a key process parameter of laser additive manufacturing, the effect of interlayer cooling on the microstructures and mechanical properties of near-β high strength titanium alloys fabricated by laser additive manufacturing is not clear. This paper is committed to investigate the difference of microstructures and tensile properties of as-deposited TC17 titanium alloys experiencing different interlayer cooling time, and explore the internal variable rules contributing to the massive applications of laser additive manufacturing of TC17 blisk.

The LMD-V coaxial powder feeding laser forming system developed independently by our research group is used to melt the TC17 powder in which the surrounding argon inside is taken as the protective atmosphere and the rolled TC17 plate is used as the substrate. The relatively mature deposition process of high strength titanium alloys is selected to fabricate two thick plates with a geometric size of 200 mm(Y)×40 mm(X)×200 mm(Z), both of which are subjected to anneal and release stress. The as-deposited and heat-treated samples used to observe the microstructures and measure the tensile properties are obtained from the steady-state region of the plate. Three rods are first removed in parallel along the deposition increasing direction (L-direction) and the laser moving direction (T-direction) of the samples, respectively, and then they are processed into the standard room-temperature tensile samples. All the heat treatment tests are carried out in the same box furnace. The microstructures of the samples corroded by Kroll reagent after polishing are observed by optical microscope and scanning electron microscope, and the volume fraction and size of α phases are counted and measured by ImageJ software.

Along the deposition increasing direction, sample A presents the morphology of alternating arrangement of the columnar grain region and the equiaxed grain region, and the grains of sample B also show the morphology of periodic arrangement, but the middle area of the molten pool is a "bamboo" grain morphology composed of a row of elongated small columnar grains and a row of fine equiaxed grains (Fig. 3). With the increase of interlayer cooling time from 0 min to 3 min, the preferred orientation of columnar grains is more obvious. The continuously formed sample A has a bimodal structure, while the interlayer cooled sample B has a basket structure containing ultra-fine α lamellar (Fig. 5). These results show that the interlayer cooling time has an appreciable effect on the microstructures of as-deposited alloys, and the size and content of the α phase are quite different. However, after the triple heat treatment, the two groups of samples show the bimodal structural characteristics (Fig. 6). The continuously formed sample A displays good comprehensive mechanical properties, including reliable tensile strength and excellent ductility (the ultimate tensile strength could reach up to 1128 MPa along the longitudinal direction of the sample and the elongation is 10.5%), which both meet or exceed the level of forgings undergoing a standard heat treatment (Table 4). The increase of interlayer cooling time leads to the mechanical characteristics performing high strength and low plasticity, and the fracture mechanism of as-deposited alloys changes from a typical ductile fracture to a cleavage brittle fracture (Figs. 8 and 9). After the triple heat treatment, the great differences of the room temperature tensile properties of two samples for different interlayer cooling time have been significantly improved, but it fails to achieve good matching in strength and plasticity.

The solidification structure of TC17 titanium alloy fabricated by laser additive manufacturing under different interlayer cooling time is the mixture of periodically arranged columnar and equiaxed grains. The increase of interlayer cooling time raises the temperature gradient of a molten pool, resulting in stronger columnar grains growth, and the equiaxed grains change concurrently to bamboo grain morphologies. The original β grain sizes of the cross sections of the two samples are almost equal. The microstructures of continuously formed TC17 titanium alloys are bimodal. With the increase of interlayer cooling time, the microstructure changes to basket-weave. Besides that, the volume fraction of α phase increases and the lamellar width decreases significantly. The tensile strength and elongation of continuously formed as-deposited TC17 titanium alloys are 1128 MPa and 10.5%, respectively, and the fracture morphology shows a typical ductile fracture. The increase of interlayer cooling time leads to the increase of strength, the decrease of plasticity, and the enhancement of anisotropy, because its fracture mechanism also changes to a cleavage brittle fracture. Two groups of TC17 titanium alloys form similar bimodal microstructures and the room temperature tensile properties after the triple heat treatment. It can be seen that an appropriate subsequent heat treatment can improve the differences of microstructures and properties caused by different forming processes.

Arc additive manufacturing has the advantages such as high stacking rate, high material utilization rate, low equipment cost, and ability to manufacture large components. Aluminum alloy parts with complex structures can be made by wire arc additive manufacturing (WAAM). What is more, WAAM can directly produce a complete component, which greatly simplifies the processing process. However, the aluminum alloy component made by WAAM has the defects of poor forming quality, many defects, and low mechanical properties, while the laser induced arc composite additive (LIACD) technology can effectively improve the forming quality and mechanical properties. Recently, most of research on WAAM of aluminum alloys is still focused on single-channel multi-layer thin-walled parts, but there are few researches on multi-channel and multi-layer arc additive manufacturing. Path strategy is the first step of WAAM. In the case of multi-layer and multi-channel, the heat dissipation conditions of different stacking paths are different, resulting in changes in the microstructures and properties of specimens and thus influencing the application fields. The LIACD manufacturing technology is based on arc and supplemented by a laser, which can further improve the forming quality and mechanical properties of arc additive manufacturing. In this paper, the laser induced MIG composite additive manufacturing technology is used to fabricate 2319 aluminum alloys. The effects of two deposition paths, namely unidirectional linear deposition and crisscross deposition, on the microstructures and properties of 2319 aluminum alloy blocks are studied.

In this study, a low power pulsed laser-MIG composite heat source is used for resurfacing welding of ER2319 welding wires on the 2219 aluminum alloy substrate. The protective gas used in this experiment is Ar with a mass fraction of 99.99% and a flow rate of 20 L/min. The welding current is 140 A, the average laser power is 300 W, the scanning speed is 450 mm/min, the wire feeding speed automatically matches the welding current, the inter-channel overlap rate is 25%, and the wire dry elongation is 11 mm. The interlayer cooling time is 60 s after each layer is stacked, and the inter-channel cooling time is 30 s after each layer is stacked. Two paths of unidirectional linear deposition and crisscross deposition are selected to make the aluminum alloy blocks. The two groups of samples are wire cut. The crisscross sections of these samples are first grinded with sandpaper and then polished and etched successively. The microstructures and fracture morphologies are observed under optical microscope and scanning electron microscope, respectively. The mechanical properties of these two groups of specimens are tested by the micro-hardness tester and the universal tensile test machine. The microstructures and fractography of the two groups of samples are analyzed by energy disperse spectroscopy (EDS).

As for the macroscopic forming quality, as shown in Fig. 5(a), for the unidirectional linear deposition, there exists a uniform fish scale shape in its cross section and obvious wave lines between layers. Aa shown in Fig. 5(b), for the crisscross deposition, there exist fish scales and strips alternating in cross section. In terms of microstructures (Figs. 6, 7, 8 and 10), the grains for the unidirectional linear deposition are smaller, and the growth direction of the columnar crystals is consistent. The pores are mainly distributed in the transition region between deposition layers, and the eutectic structures at the grain boundary present a chain distribution. The grains for the crisscross deposition are relatively thick and the distribution of columnar crystals is disorderly. The eutectic structures at the grain boundary show two forms of chain and bone, and the distribution of stomatal defects is wide. Because of the addition of a pulsed laser, a layer of fine equiaxed crystal region is generated between layers, and the distribution of the equiaxed crystal region for the unidirectional linear deposition is continuous, while the cross distribution is discontinuous. The difference in mechanical property for two paths can be seen from the cloud maps of hardness distribution (Fig. 11). The overall hardness distribution for two paths shows a trend of first high, then low, and finally high from bottom to top, and the porosity has a great influence on hardness. This phenomenon is related to the change of grain morphology at different positions of the sample. There is a large softening zone in the crisscross, indicating that the porosity defect has a great influence on the hardness distribution.

Observed from the macroscopic morphology, the grain size of the microstructures for the unidirectional linear deposition is smaller and the defects are less than those for the crisscross deposition. In terms of performance, the average hardness for the unidirectional linear deposition is 97.9 HV, and that for the crisscross deposition is 89.2 HV. The strength and plasticity for the unidirectional linear deposition specimens are anisotropic, while those for the crisscross deposition specimens are isotropic. The tensile strength for the unidirectional linear deposition specimens sampled along the X-axis is 233.58 MPa, and the elongation is 6.34%. The tensile strength sampled along the Y-axis is 275.52 MPa, and the elongation is 11.12%. The ultimate tensile strength for the crisscross deposited specimens sampled in the XY plane is 251.33 MPa, and the elongation is 7.68%. From the tensile property diagram (Fig. 12) and the SEM fractography for the tensile test specimens (Fig. 13), it is found that the crisscross plasticity is poor, and the large holes appear at the fracture, which influences the performance. The results show that the overall microstructure and properties of the aluminum alloy blocks in the Y-axis direction produced by laser induced MIG additive manufacturing unidirectional linear deposition are higher than those obtained by crisscross deposition.

During laser powder bed fusion (PBF), the complex thermal cycles and intensive heat accumulation may lead to serious residual stress and deformation, which have a great impact on the forming accuracy of parts. For parts with characteristic structures such as thin-walled structures and cantilever beams, various scanning strategies can cause different residual stresses and deformations. Thus, it is important to study the influence of printing parameters on the characteristic structures to improve forming accuracy. Numerical simulation can be used to well predict the residual stress and deformation of characteristic structures in order to enhance the data integrity and further avoid the deficiencies of experimental measurements. However, it is difficult to simulate the residual stress and deformation of large-scale structures by using the thermal-elastic-plastic model due to the high computational cost. In this work, the inherent strain method is adopted to achieve a fast and accurate prediction of residual stress and deformation, which is based on the improved intrinsic strain theory for different characteristic structures under four scanning strategies.

In order to achieve an efficient and accurate prediction of residual stress and deformation of characteristic structures, the modified inherent strain method is used for simulation. The thermal-elastic-plastic model is first established using the ABAQUS software to simulate the temperature and stress fields during the two-layer laser PBF of Ti-6Al-4V for the 0° line scanning strategy. Subsequently, the elastic and plastic strain vectors of multiple points along the scanning path are extracted. The inherent strain vector of each point and the averaged inherent strain vector for the 0° line scanning strategy are finally calculated. The inherent strain vector for the scanning strategy along the 0° direction without rotation between layers is obtained via averaging the inherent strain vector of each point. Moreover, the inherent strain vector for the 90° rotation scanning strategy is obtained via averaging the inherent strain vectors in the x and y directions. The finite element model for the characteristic structures is established using the approach of equivalent layer containing several individual layers. The direction of the inherent strain vector is updated to realize the simulations along different scanning directions via changing the allocation of material properties. The simulation of residual stress and deformation for the whole structure is completed via taking the inherent strain as the thermal expansion coefficient and sequentially activating the equivalent layer with the increase of temperature.

The modified inherent strain model is successfully used to predict the residual stress and deformation of characteristic structures during laser PBF. It is found that the crossed thin wall has concave shrinkage along the length direction (Fig. 8), and the deformation is asymmetrical along the height direction (Fig. 11) due to the asymmetric constraint during printing. The unsupported cantilever beam has serious warpage deformation and the deformation rate approaches 50% during the printing process (Fig. 13). The warpage deformation rate of the supported cantilever beam after cutting is reduced to 6.7% (Fig. 15). For the suspended circular hole structure, the vertical tensile stress on the outside of the structure increases with the increase of printing height (Fig. 19). After cutting from the substrate, the tensile stress along the z-direction on the outside of the free edge is obviously reduced and the warpage deformation occurs at the free edge (Fig. 20). The minimum deformation occurs in the scanning direction along the short side, which results from the corresponding shortest scanning vector length and deformation.

The finite element model based on the inherent strain method is developed for the fast and accurate prediction of residual stress and deformation of characteristic structures under typical scanning strategies for laser PBF of Ti-6Al-4V. The crossed thin wall has shrinkage deformation along the length direction, and the deformation first increases and then decreases along the height direction as the consequence of asymmetric constraints during the printing process. The warpage deformation of the overhanging part is reduced after cutting from the substrate compared with that of the cantilever beam without support. For the suspended circular hole structure, the vertical tensile stress occurs on the outside of the structure due to the restraint of the substrate. After one side from the substrate is cut, the free edge shows the warping deformation. The reason is that the residual stress along the z-direction on the outside of the suspended circular hole structure is released and the warpage deformation of the free edge occurs. It is found that the thin-walled structure, cantilever beam, and suspended circular hole structure show the same deformation trend for different scanning strategies. The results obtained from this research can provide valuable support for the printing of low stress and high precision additive manufacturing parts.

Generally, single material is difficult to meet the increasingly stringent industrial requirements, but multi-material components with gradient properties have broad application prospects. 316L stainless steel with excellent corrosion resistance and mechanical properties is widely used in the nuclear industry, but its high temperature stability is poor. In contrast, nickel base alloy, Inconel 718, has good thermal stability, creep strength and radiation resistance at high temperature. Because the density, specific heat capacity, thermal conductivity and other thermal physical parameters of the materials mentioned above are similar, it is easy to form a multi-material component with good interfacial bonding. Therefore, Inconel 718 can be combined with 316L to form multi-material 316L/Inconel 718, which is resistant to high temperature and large temperature difference. It can meet the strict environmental requirements of nuclear reactor and combustion wall in engine. Compared with the traditional multi-material component forming method, the laser directed energy deposition technology has obvious advantages, which can flexibly control the powder distribution and has high densification degree.

In this paper, a numerical model of laser directed energy deposition (LDED) forming of heterogeneous materials is proposed based on the finite element method, and the temperature field in the LDED forming of 316L/Inconel 718 multi-material process is established. The influences of laser power and scanning speed on the interface thermal behavior, interface defect evolution, and interface bonding performance are studied. The formation mechanism of interface defects driven by thermal action is revealed. Meanwhile, the 316L/Inconel 718 samples are formed by LDED, and the interface bonding properties of heterogeneous materials under different process parameters are characterized to verify the accuracy of the model.

When 316L material is deposited, the longitudinal section of the molten pool is narrow at the front and wide at the back (Fig. 5). The cometary tail at the end of the molten pool indicates that there is a large temperature gradient at the front of the molten pool, which could be attributed to the fact that the laser scanning speed is higher than the solidification rate of the molten pool. The heat accumulation effect and the change of thermal conductivity caused by the melting and solidification behavior of the molten pool make the laser spot center in front of the maximum temperature value. In addition, when the scanning speed of the Inconel 718 layer increases from 7 mm/s to 20 mm/s (laser power is 1100 W), the maximum temperature gradient of the molten pool increases from 6.02×10⁵ ℃/m to 1.19×10⁶ ℃/m [Fig. 7(a)]. The lifetime of liquid phase decreases from 0.52 s to 0.125 s [Fig. 8(b)], and the remelting depth decreases from 0.45 mm to 0.22 mm (Fig. 10). When the laser power of Inconel 718 layer increases from 900 W to 1500 W (scanning speed is 10 mm/s), the maximum temperature gradient decreases from 8.15×10⁵ ℃/m to 6.93×10⁵ ℃/m [Fig. 7(b)], and the lifetime of liquid phase increases from 0.3 s to 0.4 s [Fig. 8(a)]. The remelting depth increases from 0.28 mm to 0.48 mm (Fig. 9).

With the increase of v_IN718, the temperature of 316L/Inconel 718 in the positive direction of Z axis decreases gradually, which is accompanied by the decrease in the lifetime of the 316L remelting liquid phase and the molten pool size. The maximum temperature gradient is located on the 316L substrate surface and decreases gradually along the positive direction of the Z axis. With the increase of P_IN718, the positive temperature of 316L/Inconel 718 in Z axis gradually increases together with the lifetime of the 316L remelting liquid phase and the size of molten pool. The maximum temperature gradient is located on the 316L substrate surface, and decreases first and then increased gradually along the Z axis. Combined with the microstructural analysis of the LDED 316L/Inconel 718 sample, it is found that the experimental results are in good agreement with the three-dimensional finite element temperature field simulation results, indicating that the model can effectively analyze the thermal behavior of the LDED multi-material forming process.

Laser powder bed fusion (LPBF) technology is a new type of manufacturing method. The three-dimensional (3D) model is sliced using a software, and then, powder is scanned and melted layer-by-layer using laser to obtain a 3D entity. The most studied LPBF alloys mainly include titanium alloys, cobalt-chromium alloys, and nickel-based superalloys. Meanwhile, the most studied LPBF formed aluminum alloys are Al-Si alloys, such as AlSi10Mg. High-strength aluminum alloys represented by 7075 have high hardness and strength and have unique advantages in aerospace and weapon manufacturing industries. However, the wide solidification interval of high-strength aluminum alloys and poor welding performance, defects, such as solidification cracks during the LPBF process, significantly affect the mechanical properties of the LPBF high-strength aluminum alloys, and thus restrict their application in the additive manufacturing field. Current research on the LPBF forming of high-strength aluminum alloys, such as 7075, mainly focuses on material modification and process optimization. This study explores the effect of a Zr-containing amorphous alloy on the forming quality of 7075 high-strength LPBF aluminum alloy. A printed workpiece with stable quality without cracks and defects was obtained. Our study can improve the LPBF printability of high-strength aluminum alloys and promote the expansion of the available engineering material library and process parameter library for additive manufacturing.

The Zr-containing amorphous and 7075 high-strength aluminum alloys powders were uniformly mixed in a mass ratio of 2: 23 through mechanical powder mixing to obtain composite powder. First, a single-track printing experiment was performed to explore the effects of laser power and scanning speed on the morphology of the single-track channel and the size of the molten pool. Then, a single-layer printing experiment was performed to measure the surface roughness of the printed single-layer and study the effects of the hatch distance on surface quality. Finally, the optimal combination of laser power, scanning speed, and scanning distance was explored by designing an orthogonal experiment with three factors and four levels. The Archimedes drainage and metallographic image methods were used to measure the density. Scanning electron microscope and electron backscatter diffraction technology were employed to study the effect of Zr-containing amorphous alloy on the microstructure of 7075 high-strength aluminum alloy LPBF-formed samples. In addition, the microhardness of the LPBF-formed samples under different process parameters was investigated to characterize their mechanical properties.

This study reveals that the addition of Zr-containing amorphous alloys can effectively achieve grain refinement and has a significant inhibitory effect on solidification cracks in the LPBF printing process of 7075 high-strength aluminum alloy. Orthogonal experiment analysis results show that the change of the laser scanning speed significantly affects the density of the bulk sample, which is a crucial factor affecting the density within the range of the parameters studied in this experiment (Table 4). The optimal process parameters are laser power: 300-340 W, scanning speed: 600-800 mm/s, and hatch distance: 50-70 μm. As the laser energy density is increased from 33.8 to 142.9 J/mm³, the density of the formed sample gradually increases, producing a crack-free formed sample with a density of 99.4% (Fig.7). Then, the molten pool grain structure gradually changes, becoming a small equiaxed crystal with a grain size of only 1-2 μm (Fig.9). Zr-containing amorphous alloys can generate Al₃Zr particles during the LPBF process of 7075 high-strength aluminum alloy, promoting the transformation of columnar crystals to equiaxed crystals and inhibiting the generation of solidification cracks (Figs.10 and 11). The introduction of Zr-containing amorphous alloy can optimize the mechanical properties of 7075 high-strength aluminum alloy LPBF prints. The highest average microhardness value of the deposited sample reached 154.4 HV, which is 17.7% higher than that of the unmodified LPBF-printed 7075 high-strength aluminum alloy (Fig. 13).

This study aims to improve the printability and quality control of the LPBF of high-strength aluminum alloys, innovatively introducing Zr-containing amorphous alloys into 7075 high-strength aluminum alloys and study the effect of Zr-containing amorphous alloy on the printing quality of 7075 high-strength aluminum alloy under different process conditions. The analysis of the single-track printing experiment shows that the laser power and scanning speed affect the size and morphology of the single melt pool. Extremely high or low laser power and scanning speeds are not conducive to the formation of stable melt tracks. Analysis of the surface roughness of single-layer printed sample shows that the size of the hatch distance affects the quality of the single-layer surface. When the overlap ratio is 60%-80%, the printing quality of a single-layer surface with a surface roughness of less than 13 μm can be obtained. Orthogonal experimental results show that the highest density and microhardness process overlap, the optimal process parameters are laser power: 300-340 W, laser scanning speed: 600-800 mm/s, and hatch distance: 5070 μm. As the laser energy density increases from 33.8 to 142.9 J/mm³, the molten pool grain structure gradually transforms into fine equiaxed crystals, thereby realizing the complete grain refinement. The crack-free sample with density of 99.4% is obtained, and the hardness of the deposited sample reaches 154.4 HV. The introduction of Zr-containing amorphous alloys can achieve grain refinement in the LPBF printing process of 7075 high-strength aluminum alloy, which has a significant inhibitory effect on the solidification cracks and can improve the mechanical properties of 7075 high-strength aluminum alloy LPBF-formed samples.

Ni-based superalloys are the ideal high-temperature materials due to their excellent oxidation resistance and microstructure stability at the elevated temperature. But the technical characteristic and limitation of traditional manufacturing techniques restrict the development and product of superalloy components. Recently, the laser additive manufacturing (LAM) technology provides an effective tool to fabricate the integral and complex components. However, the compositions of heritage alloys are designed based on the traditional techniques without considering the specifications of the LAM process. During the LAM process, the non-equilibrium characteristics of multiple thermal cycles, rapid heating and cooling rate and the localized microstructural evolution result in the metallurgical defects such as cracks, pores, and lack of fusion, which are difficult to be completely eliminated by optimizing the process parameters. Note that the Ni-based superalloys are developed from the Ni-20%Cr alloy, and the Ni-Cr-Al system can be regarded as the basic composition of the Ni-based superalloys. The basic composition plays a significant role in the further design of multicomponent alloys owing to the correlation to the microstructural stability, mechanical properties and weldability of Ni-based superalloys. Therefore, it is necessary to optimize the compositions by investigating the influence of Cr and Al contents on microstructures and properties of the basic alloys. In this paper, five representative basic alloys are designed based on alloying of binary Ni-20%Cr alloy with Al element, and the influence of composition on microstructures and properties of as-deposited alloys is systematically investigated. This research can be helpful to design the Ni-based superalloys which are fit for the LAM process.

Five basic alloys are first designed through the "cluster-plus-glue-atom" model, and then fabricated by the LAM process. The pure Ni plate is chosen as the substrate. The elemental powders of Ni, Cr and Al with purity (mass fraction) of 99.90%-99.99% and particle size of 50-150 μm are chosen as feedstock materials, and the powders are blended by a ball grinder for 10 h. The specimens are built on the LDM-800 additive manufacturing system using a strategy of bidirectional scanning, and the process parameters are optimized as laser power of 2 kW, scanning speed of 5 mm/min, laser beam diameter of 2 mm, overlapping rate of 50%, and powder feeding rate of 6.8 g/min. The LAM specimens are cut along the build direction for the microstructural and mechanical property analysis. The crystal structures of as-deposited alloys are identified through X-ray diffraction. The microstructural evolution and elemental distribution are analyzed by scanning electron microscope (SEM) and electron probe microanalyzer (EPMA), respectively. The precipitated phase is identified and investigated by transmission electron microscope (TEM) equipped with the selected-area electron diffraction. The microhardness is measured by the hardness tester, and the room compressive test is tested on a material testing machine. The continuous variable-temperature oxidation is performed on the thermal analyzer. In order to evaluate the weldability of alloys, three cross-sections of each alloy along the build direction are observed by optical microscopy and the solidification temperature range of alloy is measured by differential scanning calorimetry (DSC).

The matrices of Ni_75.0Cr_25.0 and Ni_75.0Cr_18.75Al_6.25 alloys are composed of the γ-Ni solid solution, while the γ′ phase begins to precipitate in γ-matrix when the Al content (atomic fraction) is higher than 6.25% (Fig. 1). Additionally, the α-Cr solid solution distributes between the columnar grains of the Ni_75.0Cr_25.0 alloy, and the amount of α-Cr solid solution increases with the increase of Al content. In the Ni_75.0Al_25.0 alloy, the α-Cr solid solution is replaced by γ′-Ni₃Al+ γ-Ni divorced eutectic (Fig. 2). The microhardness and strength of as-deposited alloys first slightly increase with the increase of Al content, and then sharply increase when the Al content is higher than 6.25%, owing to the precipitation of γ′ phase (Figs .4 and 5). But the ductility significantly decreases with the increase of Al content (Table 4). The continuous variable-temperature oxidation curve and the oxidation kinetics data show that the initial temperature of vigorous oxidation overall increases with the increase of Al content, while the total mass gain and mass gain rate change in an opposite trend (Fig. 7 and Table 5), and the improved high-temperature oxidation resistance can be attributed to the formation of Al₂O₃ oxide scale. However, the excessive Al content enlarges the solidification temperature range and deteriorates the weldability of alloys, and the large amount of pores and lack of fusions are formed in Ni_75.0Al_25.0 alloy (Fig. 10 and Table 7).

In this paper, five representative basic alloys are designed using the cluster model based on alloying of the binary Ni-20%Cr alloy with Al element. The influence of composition on microstructures and properties of the as-deposited alloys is investigated. The results show that the matrix structures of as-deposited alloys evolve from γ-Ni solid solution to γ′-Ni₃Al strengthening phase with the increase of Al content. Meanwhile, α-Cr solid solution distributing along grain boundaries changes from granule to long-chain in morphology while increasing its fraction, and is replaced by γ′-Ni₃Al+ γ-Ni divorced eutectic at 25.0% Al. The hardness and strength of as-deposited alloys increase with the increase of Al content due to the change in the strengthening mechanism from solid solution strengthening to precipitation strengthening, whereas the ductility decreases. The increase of Al content is beneficial for improving the high temperature oxidation resistance, but the excessive Al is deleterious to the weldability. Therefore, the Al content should be confined in a range of 12.5%-18.75% to make the basic alloys have a good match of mechanical properties, high temperature oxidation resistance and weldability.

As advanced high strength steel, maraging steel is mainly used in the fields of aviation, atomic energy, and high-end tooling. The traditional methods for the preparation of maraging steel including cold rolling, homogenizing treatment, solution treatment and aging are complex, time-consuming and costly. Additive manufacturing (AM) has a wide application in the forming of complex shaped parts, which mainly depends on its flexibility, parameter controllability, and high cooling rate. AM can provide an effective way to promote martensite transformation, grain refinement, and forming of complex internal cavity structural molds.

The combination of AM technology and maraging steel has great advantages in realizing personalized customization and reducing subsequent processing. The AM of maraging steel opens a new direction for the preparation of high-end die steel. However, it has complicated solidification metallurgy and solid phase transformation behaviors, although maraging steel fabricated by AM has excellent strength, toughness, hardness, corrosion resistance, and wear resistance. Therefore, it is of great significance to deeply understand the correlation among process, composition, and microstructure of maraging steel fabricated by AM.

To improve the mechanical properties of maraging steel fabricated by AM, the post-treatment methods mainly including heat treatment and hot isostatic pressing (HIP) are used, which can effectively decrease the internal defects of forming parts and reduce the internal stress, the porosity, and the un-melted powder particles. However, the maraging steel fabricated by AM has anisotropic mechanical properties along the building direction and the scanning direction. Therefore, the process and properties of maraging steel fabricated by AM need to be further studied.

With the maturity of laser AM technology and the improvement of mechanical properties of building parts, the AM of maraging steel is more widely used in industry and daily life. Laser cladding AM technology has also been applied to metal surface coating, but the research of laser cladding AM technology on coating is relatively less. Above all, the review and prospect of the AM of maraging steel are important in guiding the development of maraging steel with high-performance by AM and regulating of microstructures.

The advances in the improvement of mechanical properties, the post-treatment technology, and the microstructural control of maraging steel are reviewed. Table 1 summarizes the typical mechanical properties of maraging steel fabricated by AM in recent years. There are many researches on additive manufacturing of maraging steel. Bai’s research showed that with the increase of solid solution temperature, the hardness and strength of maraging steel were significantly improved with the formed Ni₃Mo, Fe₂Mo and Ni₃Ti precipitation. Tan et al. eliminated the residual pores of parts by HIP. Compared with that of the as-fabricated parts, the porosity of the maraging steel treated by HIP is obviously reduced [Figs. 2(d-e)]. Yin et al. studied the microstructure of LPBF-18Ni300 maraging steel and found that the very fine cellular structure could improve the hardnesses and other properties of parts [Fig. 2(c)]. According to the transfer mode of powder and the thermal history effect of AM, the heterostructural maraging steel can be obtained. Kurnsteiner successfully prepared Damascus-like maraging steel with a composition of Fe-19Ni-5Ti through direct energy deposition (DED), and the ultimate tensile strength of heterostructural steel was 200 MPa, higher than that of conventional homostructural steel [Figs. 3 and 4(a)]. Tan’s research proved that the heterostructural steel could also be obtained by using the alternating layer printing method [Fig. 4(b)]. AM can achieve a variety of comprehensive properties by combining two metal materials together, i.e., functionally graded structural materials [Figs. 4(c) and 5(b)]. The results show that the thermal diffusion interface between two different materials has good interface bonding and mechanical property stability. Particle-reinforced maraging steel composites with high strength and wear resistance can be prepared by adding ceramic powder in maraging steel. It shows that the WC particles can significantly reinforce the maraging steel composites and the excellent comprehensive mechanical properties are obtained.

The AM of maraging steel breaks through the limitation of a traditional process and shows a great prospect in the manufacturing of molds with complex structures. However, further researches on the AM of maraging steel are needed. In terms of composition design, the development of low alloy elements, the reduction of preparation cost, and the development of special powder for the AM of maraging steel become new research contents. On the optimization of process parameters, reducing laser power and improving scanning speed are worth considering for improve the efficiency and reduce the cost. The microstructure of maraging steel fabricated by AM is dominated in columnar grains, resulting in the mechanical anisotropy between the building direction and the scanning direction, which is also a vital problem to be solved. These defects or undesired microstructures mentioned above can be avoided by controlling the scanning mode of AM in the future, i.e., controlling the energy output and properly preheating of the substrates. An artificial neural network plays an important role in the field of materials science. In the future, machine learning can also be used to assist the composition design of maraging steel, the optimization of process parameters, and the prediction of initial temperature for martensite transformation (M_s).

Nickel-based single crystal superalloys have excellent high-temperature properties and are mainly used in the manufacture of aero-engine turbine blades. A high temperature gradient is the key to single crystal formation. Selective laser melting (SLM) is a kind of metal additive manufacturing, which uses a laser as the heat source to form the parts with fine metal powders layer by layer, especially the complex structures. The small laser spot (80-100 μm) has extremely high thermal density, which promotes a temperature gradient of 10³-10⁵ K/cm and a solidification rate of 10⁵-10⁷ K/s in the molten pool. In addition, a laser has obvious directional heat transfer characteristics, and the solidification heat can be directionally dissipated along the path from the molten pool to the substrate. Therefore, to explore the feasibility of single crystals fabricated by SLM, thirty-five single track samples are fabricated under different laser powers and scanning speeds. The geometric features, crystal orientation, and microstructures are first analyzed systematically. Then, the formation mechanisms of cracks and stray grains are discussed, which provides theoretical and technical reference for the preparation of large-scale single crystal structures by SLM.

The particle size distribution of the gas-atomized DD91 nickel-based single crystal superalloy powder is in the range of 8.6-37.5 μm with an average of 20.4 μm. The SRR99 single crystal substrate prepared by directional solidification is tested by electron backscattered diffraction (EBSD) and it is found that the grain grows along the direction of [001], which meets the experimental needs. Before the experiment, the powder is placed in an oven at 80 ℃ for 10 h to remove moisture. The single crystal substrate is polished with sandpaper, and impurities such as abrasive debris are removed in an ultrasonic cleaning machine. DD91 alloy powder with a layer thickness of 30 μm is spread on the (001) crystal surface of the SRR99 single crystal substrate. A self-developed SLM equipment with a laser spot diameter of ~100 μm is used for the printing experiments. Thirty-five single-track samples with different parameters are fabricated on the substrate with a laser power of 245-305 W and a scanning speed of 500-1100 mm/s. Optical microscope (OM) is used to observe the surface morphologies of the single-tracks and evaluate the forming quality. Scanning electron microscope (SEM) is used to observe the microscopic morphologies and microstructures of powder and molten pools. ImageJ software is used to measure and analyze the geometric feature sizes of molten pools. EBSD is used to analyze the crystal orientation with a scanning step of 0.25 μm, and AztecCrystal software is used to process the collected data.

When the line energy density is in the range of 414-580 J·m^-1, excellent single-tracks could be obtained with clear surface contours, continuous flatness, and obvious metallic luster on the surface (Fig. 3). The shape of the molen pool is more regular and stable at a higher power and a lower speed, which could provide good geometric conditions for the epitaxial growth of single crystals. At a laser power of 290 W and scanning speeds of 1000 mm/s and 1100 mm/s, cracks perpendicular to the scanning direction are produced, running through the entire single-track. The dendrites are exposed to the crack section, leaving a wide irregular crack, which is the characteristic of a typical solidification crack, probably due to the excessive Al content of the alloy and the tearing of the liquid film between the dendrites under stress (Fig. 4). Most of the grains in the molten pool exhibit a very clear meritocratic orientation, i.e., they can continue the orientation of the single crystal substrate and grow along [001] epitaxially. However, there are still crystal defects in some molten pools (Fig. 7). They can be divided into three categories: first, poor metallurgy between the molten pool and the single crystal substrate, resulting in disordered grain growth; second, during the growth of the crystal, its growth direction is shifted by a small angle, producing an orientation deviation due to the flow field in the molten pool; third, the columnar to equiaxed transition (CET) occurs at the top and sides of the molten pool due to the reduced temperature gradient and the accelerated solidification rate at the late solidification stage, resulting in equiaxed stray grains (Fig. 8). At the bottom of the molten pool, near-parallel columnar dendrites grow along [001], continuing the orientation of the substrate without secondary dendrite arms. In the middle of the molten pool, the grains also grow and continue the [001] crystal orientation, however, the grain morphology changes to a spindle shape. At the top of the molten pool, the grain growth direction changes from [001] to [100] (Fig. 9). Due to the larger height of this molten pool, the distance of heat transfer to the substrate from top to bottom at the end of solidification becomes farther and the heat is mainly transferred along the scanning direction (parallel to the substrate). Suitable process parameters should be selected to minimize the height of the molten pool, thus to reduce the heat transfer distance from the top to the substrate and reduce the generation of such defects.

Continuous, smooth, and straight single-tracks could be obtained within a fixed layer thickness of 30 μm, a laser power of 290-305 W and a scanning speed of 500-700 mm/s. The direction of the temperature gradient in the flat and regular molten pool is closer to[001], which is more conducive to the unidirectional heat transfer from top to bottom in the molten pool, and provides a guarantee for the stable growth of the single crystal. A good metallurgical bond between the molten pool and the single crystal substrate is the key to the epitaxial growth of crystals. It is difficult to obtain the initial orientation of the substrate for the molten pool under poor metallurgical conditions, resulting in disordered crystal growth. Small angular orientation deviations occur due to the internal flow field. A small number of stray grains occur at the top and sides of individual molten pools, caused by the lower temperature gradient and the increased solidification rate in the later stages of solidification. The grains in different regions within the molten pools have different characteristics. Grains at the bottom and both sides of the molten pool can continue the substrate [001] epitaxial growth, and the primary dendrite arm spacing is only 0.6-0.8 μm, much lower than that of the directional solidification. Spindle-shaped grains occurring in the middle of the molten pool also continue the [001] orientation, but the solidification shrinkage holes with a width of 0.1 μm generate between the dendrites. At the top of the molten pool, the grains show a fine cellular structure, growing along the [100] direction, which is caused by the heat transfer along the scanning direction (parallel to the substrate). This study provides a process basis for the fabrication of large-size single crystals by SLM.

High-entropy alloy (HEA) has the advantages of high strength, high hardness, high temperature resistance, radiation resistance, and strong corrosion resistance. The refractory HEA system composed of high melting point elements such as Nb, Mo, Ta, W, V, Re, and Hf has become a material of great research value in the future development of aviation industry and military. However, the refractory HEA composed of ultra-high melting point elements such as Nb, Mo, Ta, and W has poor plasticity at room temperature, and the formed samples are prone to cracks. Therefore, in this study the forming plasticity of WNbMoTa HEA is increased by adding plastic elements to form a special HEA powder material for additive manufacturing. The crack-free RHEA01 HEA is formed by laser selective melting. The structure, morphology and properties of the RHEA01 HEA are analyzed, and the HEA samples with excellent mechanical properties are obtained.

In the experiment, the Nb, Mo, Ta, Ti, and Ni powders are mixed in an equal or near equal atomic ratio. After the powder proportioning is completed, the selective laser melting (SLM) equipment independently developed by Xi’an Jiaotong University is used for RHEA01 SLM additive manufacturing forming. The GeminiSEM 500 electron microscope is used for the morphological EDS and EBSD analysis. The Bruker D8 ADVANCE XRD equipment is used for the crystal structural analysis. The HX-1000TM equipment is used for the room temperature hardness testing. The mechanical property testing machines are used for the room temperature and high temperature compression performance test. The Hopkinson pressure bar is used to conduct the dynamic compression mechanical performance test experiment.

It can be seen from the BSD image (Fig. 2) that Ti and Ni, which are quite plastic during the solidification of the alloy, supplement the shrinkage of the high melting point elements after solidification. The refractory HEA RHEA01 formed in this study has no cracks and good formability. As shown in Fig. 3, the refractory HEA RHEA01 has a single-phase BCC structure. Note that there are some miscellaneous peaks around the peak of RHEA01. These peaks are a small amount of new phase structures formed after Ti and Ni elements fed in the later stage of solidification. It can be seen from the surface energy spectrum in Fig. 4 that the elements in RHEA01 formed by SLM are uniformly distributed without macro-segregation. As shown in Fig. 5, the average grain size of the RHEA01 alloy is about 8.5 μm. RHEA01 has an obvious grain boundary filling phenomenon due to the small amount of Ni and Ti elements from the BSD diagram. From the perspective of grain size, RHEA01 has a relatively good theoretical strength due to the effect of grain refinement. At the same time, because it has a single-phase BCC structure composed of a large number of high melting point elements, it has good high temperature resistance.

SLM is used to form crack-free RHEA01 with added plastic elements, which solves the crack defects in the process of SLM forming of NbMoTa HEA and greatly increases the formability of NbMoTa HEA. The average grain size of RHEA01 is 8.5 μm and it is a single-phase BCC structure. Compared with NbMoTa HEA formed by SLM, the addition of plastic elements refines the grains. The yield strength and normal temperature compressive strength of RHEA01 are 1277.35 MPa and 1597.62 MPa, respectively, 27.9% and 36.9% higher than those of NbMoTa HEA (VAM). The hardness is 511.76 HV. The high temperature resistance of RHEA01 has good retention. The compressive strength at 1000 ℃ is as high as 993.84 MPa, only 37.8% lower than that at normal temperature. The dynamic compressive strength at 1400 ℃ (temperature)and 2000 s^-1 (strain rate) is as high as 1015 MPa. The comprehensive performance of RHEA01, the high temperature compression resistance, and the normal temperature compression resistance are all at the world’s leading level.

The macrostructure in the Z direction of the TC4 as-deposited state is composed of coarse original β columnar crystals that grow approximately parallel to the deposition direction, and the macrostructure in the XY direction is composed of equiaxed crystals. With the increase of V content, the average width of the bottom of Z directional columnar crystals decreases from 1 mm to 0.24 mm, and the average length of equiaxed crystals in the XY direction at the top of columnar crystals decreases from 600 μm to 250 μm. The microstructure of the TC4+ V as-deposited state is composed of a large amount of α phases with densely packed hexagonal structures and a small amount of β phases with body center cube structures. With the increase of V content, the primary α is gradually refined with the gradual decrease of the length-width ratio (lowest to 4.9) and width (lowest to 0.2 μm). Meanwhile, the content of β phase gradually increases to 30.91% (TC4+ 10%V). Thus, under the effect of solution strengthening, the Vickers hardness of the samples slightly increases with the increase of V addition. The increase of equiaxed α phase decreases the microhardness slightly. Under the effect of fine grain strengthening and solid solution strengthening, the overall tensile strength of the sample as-deposited with V is higher than that of TC4. However, due to the increase of intragranular component segregation and the thinning of the width of the secondary α photo layer, the elongation decreases gradually. The results show that the TC4+ 6%V as-deposited sample has the best comprehensive mechanical properties, with the maximum tensile strength (about 1018 MPa), yield strength (about 892 MPa), elongation (about 11.2%), and hardness (about 381.7 HV).

Additive manufacturing of magnesium alloys has broad application prospects for customized and complicated structures used in the fields of aerospace, transportation, and medical devices. However, magnesium has the characteristics of high reactive activity, high vaporization, susceptibility, high molten fluidity, and high thermal expansion. Thus, the additive manufacturing of magnesium alloys faces many process challenges. This paper systematically investigates the process optimization during laser powder bed fusion (L-PBF) of WE43 alloys.

WE43 powder has a diameter range of 15-50 μm and an average grain diameter of 28.9 μm (Fig. 1). The BLT S210 equipment for the L-PBF processing of WE43 is equipped with a customized auxiliary gas flow system (Fig. 2). The WE43 bulk parts are built with a set of parameter windows (Table 2), and the relative density of the parts is measured with optical methods. To enhance the forming accuracy of porous structures, the influence of laser parameters on the size of as-built parts is studied, and the thin rods with different diameters are built using different optimized parameters. The four types of porous scaffolds are designed and manufactured: body centered cubic (BCC) lattice, diamond (D), sheet-gyroid (SG) and lattice-gyroid (LG) (Fig. 3). The tensile test for bulk parts, compression test for porous structures, and optical and scanning electron microscope (SEM) morphologies are done.

For the adopted L-PBF equipment and the WE43 powder materials, high fusion quality can be obtained in a large parameter range (60 W≤P≤100 W, 600 mm/s≤v≤ 800 mm/s), and the relative density of the sample exceeds 99.50% (Fig. 4). The micro-hardness of as-built bulk parts is (86.11±3.85 )HV, while the tensile strength and elongation reach (275±2.81) MPa and (15.6±1.3)%, respectively. The tensile fracture surface is rough, no obvious forming defects are found, and a certain number of dimples can be observed, showing that the characteristics of partial plastic fractures show good ductility (Fig. 9).

The accumulation and melting of multiple layers of L-PBF lead to repeated heating, forming a microstructure similar to that of the welding heat affected zone (Fig. 8). Oxides are found in the EDS mapping of WE43 as-built parts, which may come from the oxide shell of powders. During the processing, the oxide shell around the powder is broken to form irregular flakes, which is enriched with elements of O, Y and Zr and mainly composed of Y₂O₃. The behavior of oxides is closely related to powder materials and rare earth elements.

The thin rods built by three optimized parameters all have high relative density, indicating that the change of part feature size in the range of 0.3-2.0 mm has no obvious impact on the consistency of forming quality (Fig. 5). Therefore, the laser energy optimized by the block process test can be used to manufacture the porous support composed of thin rods. However, the laser energy input influences the size of a thin rod. When a higher energy input is applied, the molten pool gets wider and deeper (Fig. 7). If the scanning contour is the same, the part built by high laser energy expands in size perpendicular to the building direction. The above results show that when the porous support with WE43 magnesium alloy is formed by the L-PBF process, the laser energy input should be optimized according to the requirement of relative density (fusion quality). When the laser energy input is determined, the contour scanning indentation should be further determined according to the rod diameter (characteristic size) of the porous support, so as to improve the forming accuracy.

The geometrical error of porosity between the designed and fabricated porous scaffolds is within 10%. For the porous scaffolds with SG and LG units, the powder adhesion and roughness distribution on the surface are relatively uniform, and the spatial distribution of forming error is relatively consistent, which is more conducive to the subsequent process to remove the powder adhesion and improve the surface quality. There is a large deviation in the spatial distribution of the forming error of the porous scaffolds with Diamond and BCC units. The powder sticking is more serious at the connection of the support rod and the surface with a small suspension angle (Fig. 6). The compressive strength and modulus of elasticity of porous scaffolds are 17.6-37.6 MPa and 325-619 MPa, respectively, far lower than those of solid WE43 and equivalent to those of cancellous bone (Fig. 10). By designing the pore units with different shapes, sizes, porosity, and distributions, the mechanical properties of the porous scaffolds with the same appearance and size can be adjusted in a wide range so as to obtain the mechanical behavior most conducive to the application scene. Good process performance and forming quality are the fundamental guarantee to realize the design intention.

The results indicate that the L-PBF process for WE43 alloys has a promising prospect for the industrial application. Laser parameters and scanning strategy are optimized to achieve high formation quality and accuracy. The microstructures of the WE43 as-built parts show a great difference with those of the traditional manufactured ones.

Additive manufacturing (AM) has been widely used in aerospace, military, medical, automotive, nuclear power, and other fields due to its great potential in producing lightweight parts with complex structures and high personalization. As one of the most widely used AM technologies, laser powder bed fusion (LPBF), also known as selective laser melting, is characterized by huge temperature gradients, drastic phase changes, and extremely unstable molten pools. Because of its particularity in the manufacturing process, there easily exist internal defects of parts and harmful influence on forming quality and mechanical properties. Some of the defects cause the reduction of density, and further reduce other mechanical properties. For example, micropores and lack of fusion lead to the reduction of density, resulting in the decrease of its strength, hardness, and fatigue strength. The balling and spattering caused by lack of fusion also influence the surface morphology and phase composition. In addition, the residual stress generated in the machining process also causes cracks and warping deformation. The cracks influence the performance, while warping deformation influence the dimensional accuracy. Therefore, it is of great significance to understand the characteristics, formation mechanism, and influencing factors of defects, so as to explore the control mechanism of defects and control the quality and performance of parts.

This paper introduces the characteristics, formation mechanism, and influence of common defects on density, including unmelting caused by lack of fusion (Fig. 2), balling (Fig. 3), spatter (Fig. 5) and micropores (Fig. 6) , and the crack (Fig. 7) and warping deformation (Fig. 8) caused by residual stress. The effects of control methods of forming process and composite manufacturing on defects are discussed. The control methods of forming process can be classified into five categories: process gas supply, powder bed, laser beam, processing parameters, and scanning strategy (Fig. 9). A suitable processing environment and powder bed are the basis to prevent defects. As one of the most important factors in LPBF, a laser has many controllable aspects (Fig. 10). Proper focus shift, spot size, and intensity distribution can not only enhance the stability of the molten pool, but also improve the morphology of the molten pool and the microstructures of the parts. In addition, choosing an appropriate wavelength of a laser according to the absorptivity of materials can both improve the energy utilization of the laser and effectively inhibit the generation of lack of fusion defects. Process parameters are the most flexible means to control defects. Scanning strategy can change the overlap ratio and stress distribution in the process to inhibit the generation of defects, effectively inhibit the continuation of defects, and even eliminate the defects ever generated (Fig. 12). The composite manufacturing control methods are divided into additive-subtractive hybrid manufacturing and the multi-energy assisted process. Additive-subtractive hybrid manufacturing has the flexibility of AM and the ability of milling to eliminate the internal defects and improve the internal quality of parts (Fig. 13). However, it is inevitable that the alternate processing of adding and reducing materials greatly reduces the processing efficiency of parts, so a more efficient way of additive-subtractive hybrid manufacturing needs to be explored urgently. Magnetic field assisted LPBF has the effects of homogenizing microstructure, refining grains, inhibiting segregation, and reducing density. Ultrasonic-assisted LPBF has a positive effect on reducing residual stress, and improving anisotropy and performance of parts.

Understanding the typical features, formation mechanisms, and influence of defects can discover the relationship between various factors in the forming process and defects or internal quality of parts more effectively, which is helpful for researchers to explore various control methods of the forming process. With the development of the LPBF technology, the adjustment of laser parameters in forming process and the new composite manufacturing control methods have also been investigated, providing some multi-dimensional and more advanced control methods of defects. For example, the problem that LPBF is difficult to process pure copper and other infrared high reflection metal materials is solved by adjusting laser wavelength. Meanwhile, these control methods have a great potential and play a great role in eliminating defects, improving product quality and processing efficiency, and regulating microstructure. Adjusting spot morphology can improve the heat distribution in the process, which has the potential to change the morphology and improve the stability of molten pools, and further improve the stress distribution of parts. The magnetic field assisted LPBF forming process has the function of refining grains and homogenizing structures and compositions. The ability of ultrasound to refine grains and reduce defects has been demonstrated in directional energy deposition. In addition, the on-line inspection technology assisted LPBF is also a trend of defect control technologies. Finding the defects by on-line monitoring and combining with defect characteristics, formation mechanism, and control methods are able to achieve the closed-loop control of defects, which can greatly improve the stability and reliability of part forming quality and performance.

Metallic components have been widely used in aviation, aerospace, marine and other industrial departments for their excellent properties. Over the past decades, metallic components are developing toward to high-performance and multi-function but with low production cost, which pushes new manufacturing techniques to be used. Laser additive manufacture (LAM) is one of the new additive manufacturing techniques, which is widely used to manufacture near-net-shaped metal parts layer-by-layer by melting metal powders using a laser beam. Now, it is also widely used to manufacture large and critical high performance metallic components with advantages in reducing material waste, production time and cost. In the LAM process, the components undergo periodic and unstable thermal cycling, which influences the microstructure and internal stress. More importantly, the inappropriate selection of process parameters leads to the appearance of defects with degradation of mechanical property. For example, if there exists porosity in the material powder, the pores occur during LAM and the liquid condensation speed is faster than the gas escape speed. If the selected laser power is too low, the powder is hardly fully molten and the unfused pores and inclusions are generated.

Therefore, the extensive researches worldwide focus on studying the generation mechanism of defects in the LAM process. In Deng’s research, the 15%(volume fraction) SiC ceramic reinforced steel (MS) metal matrix composites were prepared by selective laser melting (SLM). Facing the compatibility and cracking problems raised between SiC and metal matrix, great efforts on the suppression of defects during the SLM process are taken from various aspects, including laser melting, substrate preheating, and design of support and build directions. Substrate preheating can be used for the significant suppression of cracks. Tillmann et al. treated the IN718 component with the HIP method and the density of components was increased to 99.985%-99.989% within a certain range. However, most of the researches focus on just one kind of defect as well as its corresponding control methods. There are a few summaries of formation mechanisms and control methods of typical defects in LAM of metallic components.

We also summarize and analyze the formation mechanisms and control methods of three types of defects (cracks, inclusions and pores) in LAM of metallic components. The thermal cycle during the LAM process usually causes the generation of large internal stress, which causes the micro-crack formation during or after deposition. According to the formation temperature, cracks can be divided into cold cracks and thermal cracks (Figs. 1 and 2). Pores are always formed due to the insufficient energy input during the deposition process (Fig. 3) or the trapping of the residual gas inside the molten pool, which is related to the flowing behavior of molten metal fluid in the pool (Figs. 4 and 5). There are two main types of inclusion defects: one is oxide inclusions mostly caused by the mixing of oxygen in the production atmosphere, and the other is high-melting metal inclusions caused by the mixing of powders with high-melting metal powders (Fig. 6).

Different control methods are required for different types of alloys and defects. For cold cracks, the currently commonly used control methods include optimizing the compositions of metal powder and adding post-heating treatment. For hot cracks, the currently commonly used control methods include optimizing the compositions of metal powder and the process including selecting appropriate parameters (Fig. 7) and preheating substrates (Fig. 8). The pore defects can be minimized by either improving powder quality (Fig. 9) or adopting post-processing. For high melting-point metallic inclusions, the effective control methods include the selection of reasonable process parameters and scanning schedule. For oxide inclusion defects, the main control method is to control the oxygen content in environment.

The laser melting deposition additive manufacturing technology for metallic components is a high-performance, low-cost, and designable manufacturing technology. However, the degradation of material properties due to the defects is still one of the problems that must be solved during component production. It is still necessary to conduct a more in-depth research in the following aspects. For crack defects, it is necessary to discuss the stress evolution during the deposition as well as the relationship between mechanical behaviors and internal stress distributions. For pore defects, the quantification of the influence of each process parameter on the formation mechanism is required. For inclusion defects, it needs to research the thermodynamic mechanism of the protective gas composition and their flowing behaviors.

Carrying capacity is the core indicator of the value of a rocket. The weight of a rocket and its carrying capacity are trade-offs. Reducing the dead weight of a rocket and improving its structural efficiency are the keys to obtain a high carrying capacity. As an aerospace structural material, aluminum (Al) alloys are widely developed as aviation structural materials with a low melting point and high specific strength. Al alloys used as fuselage materials have become an important direction for future civil passenger aircraft material selection, and the laser melting deposition (LMD) technology needs to be developed urgently. The alloys have problems such as large hot cracking tendency and microstructural inhomogeneity that influence the properties of Al2024. In order to solve the problem of low strength of Al2024 alloys in additive manufacturing, the emerging friction stirring technology has been widely studied in which the Al2024 alloys are used. First, a two-layer gradient structure covered with a laser melting deposition layer and a stirring layer is fabricated on the substrate surface by combining the friction stirring technique. The resulting surface is then laser deposited. Friction stirring combining with laser cladding results in ultrafast laser hybrid fabrication. This gradient structure exhibits high strength, high toughness, and is resistant to fracture risk.

The diameter of the stirring head is 10 mm, and the friction stirring process parameters are the rotation speed of 800 r/min and the feeding speed of 50 mm/min. First, using the friction stirring technology, laser fusion deposited 2024 aluminum alloys repaired by different processes are prepared, and the microstructures and mechanical properties of Al alloy samples prepared by this composite process are studied. Next, the microstructures and mechanical properties of the composite-fabricated Al2024 alloys are studied with an optical microscope and a Vickers hardness tester to analyze the relationship between the process and their microstructures and properties.

The results show that the microstructures of Al2024 alloys stirred with an overlap ratio of 80% are refined and present a gradient distribution compared with those by LMD. Through the orthogonal experiment of laser fusion deposition and the single-layer multi-channel experiment, the evolution law of penetration depth and width of an Al alloy under different process conditions and the evolution law of microstructures are obtained. The spot diameter of 2.0 mm, the scanning speed of 600 mm/s, and the overlap rate of 50% are chosen as the final composite process parameters (Figs. 4 and 5). Through the microstructural analysis, it is found that the friction stirring treatment can significantly refine the grains, homogenize the structure, and reduce metallurgical defects. On the other hand, crystal grains with a directional gradient are deposited in the stirring zone, the growth direction of which is perpendicular to the substrate, and the size of the crystal grains tends to decrease. As the number of deposited layers increases, the microstructure presents fine and uniform grains without obvious directionality. Along the gradient direction, the transition layers are well combined, and the obvious gradient structure characteristics are observed as a whole (Fig. 6). After the aging treatment, the hardness of the substrate region is in the range of 135-140 HV, which is significantly improved. And with the addition of the stirring zone, the hardness after secondary deposition can reach 160 HV. In addition, as the number of depositions increases, a gradient distribution of hardness appears.

In the present study, friction stirring welding is carried out by the heat generated by the metal through friction. The difference from traditional friction welding is that it not only rubs the metal surface, but also uses a stirring rod to insert into the metal for friction (stirring), so that the depth of welding and the depth of bonding are greatly improved. Friction stirring welding does not need welding wire and has a small heat-affected range, which is especially suitable for occasions that are prone to cracking and deformation. At present, the scope of application of this technology is limited, mainly in the field of aerospace exploration and application. In this study, a novel gradient-structured aluminum alloy is successfully fabricated by the laser + stirring technology, and the surface is composed of periodic deposition layers covered with a stirring layer. After aging treatment, the hardness after secondary deposition can reach 160 HV, which indicates that LMD can also achieve ultra-high hardness. Our study shows that gradient alloys with comprehensive mechanical properties can be obtained through a rational design.

Ti-6Al-4V alloy is a typical α+ β-type titanium alloy combining the advantages of α and β titanium alloys. It has excellent mechanical properties and is the most widely used titanium alloy in marine engineering. Generally, corrosion occurs on the surfaces of titanium alloys due to the severe ocean environments. Thus, the corrosion resistance of repaired surfaces is essential for the service safety of the structural parts composed of titanium alloys. Recently, the in-air directed energy deposition (in-air DED) technique has been rapidly developed to repair damaged structures with high performances. Thus, we propose a novel technique, namely, underwater directed energy deposition (UDED), which introduces in-air DED into underwater environments. UDED has many advantages such as concentrated heat input and small heat-affected zones. Previous studies show that most of the current research has focused on the corrosion resistance of titanium alloys fabricated using in-air additive manufacturing methods. A few research studies have reported the electrochemical corrosion behavior of Ti-6Al-4V repaired using the powder-blown UDED technique. Therefore, studying the corrosion behavior of the UDED-repaired samples is essential. The results of this study are useful for the performance assessment of components repaired using UDED and optimizing the UDED process. Furthermore, this work provides an experimental basis and theoretical foundation for advancing the on-site repair of titanium alloys using UDED.

Herein, the Ti-6Al-4V alloy was employed as the testing material. First, a trapezoidal groove was machined on the as-received Ti-6Al-4V plates using wire-cut electrical discharge machining before UDED. Then, the preprepared Ti-6Al-4V was repaired in a water tank using UDED. For comparison, the preprepared Ti-6Al-4V was also repaired using the in-air DED in a homemade protective bag filled with high-purity argon. The as-deposited samples were subjected to standard metallographic preparation for OM, SEM, TEM, and EBSD observations. Then, we conducted microhardness measurements on the repaired Ti-6Al-4V zone and determined the average microhardness. Next, we conducted the electrochemical tests on the samples in 3.5% NaCl solution at room temperature using a PARSTAT 2273 electrochemical workstation with a conventional three-electrode cell. Before the electrochemical tests, the test specimens were immersed in the 3.5% NaCl solution for 24 h. Finally, the open circuit potential, potentiodynamic polarization, and electrochemical impedance spectroscopy (EIS) measurements were performed on the samples using the electrochemical workstation.

Several key findings are reported in this study. 1) The microstructure of Ti-6Al-4V fabricated using UDED is significantly finer than that using in-air DED (Fig. 3). This is attributed to the water quenching effect caused by the underwater environment, which increasing the cooling rates of the underwater melt pool and reducing the heat accumulation of the as-deposited material. This contributes to the rapid solidification of the melt pool and fast solid-state phase transformation, leading to the formation of an acicular martensite microstructure. 2) The results of EBSD show that the microstructures of Ti-6Al-4V repaired by UDED and in-air DED have multiple orientations (Fig. 4). The volume fraction of the β-phase in the sample repaired using in-air DED was slightly higher than that in the sample repaired by UDED. 3) The TEM results show that the dislocation density in the sample repaired using UDED is high due to the steep temperature gradient and rapid cooling rates in the UDED process (Fig. 5). However, the dislocation density in the sample repaired using in-air DED is low due to the significantly intrinsic heat treatment during the in-air DED process. 4) The microhardness of the sample repaired using UDED is relatively high (Fig. 6). The microstructure is strengthened by three factors, namely, the supersaturation of Al and V elements within the Ti matrix, small size of the acicular martensite, and high dislocation density within the lath martensite. 5) The Tafel polarization curves show that the UDED-repaired samples are the least susceptible to corrosion. However, when corrosion starts, their corrosion rate is relatively fast (Fig. 7). Additionally, pitting corrosion occurs on the surface of UDED-repaired samples. 6) The EIS results show that the Ti-6Al-4V repaired using UDED has the largest arc radius (Fig. 8), indicating that it has better corrosion resistance. 7) The influencing factors of microstructural features on the corrosion behavior of the Ti-6Al-4V are the grain size of the as-deposited microstructure, distribution of alloy elements, and surface activity state.

Herein, we repaired the preprepared trapezoidal grooves on the Ti-6Al-4V substrate using UDED and in-air DED. The mechanisms of different deposition processes influencing the microstructural formation/evolution, microhardness, and electrochemical corrosion behavior were revealed. The main conclusions are as follows. 1) The water quenching effect increases the cooling rates of the melt pool and decreases the thermal accumulation of the as-deposited metal. This results in the formation of acicular martensite with a high dislocation density. 2) The supersaturation of Al and V elements in the Ti matrix, fine acicular martensite, and high dislocation density increase the microhardness of UDED-repaired samples. 3) The corrosion resistance of UDED-repaired samples is better than that of in-air DED-repaired samples. The corrosion resistance of the former is mainly determined by three factors: the grain size, distribution of alloy elements, and surface state of the microstructure. The satisfactory corrosion resistance of titanium alloy repaired using UDED is important for their service safety. Additionally, the results of this work provide an experimental and theoretical basis for improving the corrosion resistance of titanium alloy structural parts repaired using UDED.

With the rapid development of metal additive manufacturing technologies, the use of selective laser melting (SLM) technology to rapidly manufacture nickel-based superalloy components has made a major breakthrough, which has greatly improved the manufacturing efficiency of high-performance complex components in the aerospace field and promoted optimized and upgraded component structures. IN625 nickel-based superalloy is maturely used in the SLM technology. It has high high-temperature mechanical properties, good high-temperature corrosion resistance and high-temperature oxidation resistance. It is used in nuclear power, industrial gas turbines and key materials for hot-end components in aerospace and other fields. The unique microstructural characteristics of SLM IN625 alloys cause their solid-state phase transition characteristics under long-term high temperature conditions to be obviously different from traditional solid-state phase transitions. In this paper the evolution of the structures and properties of the SLM IN625 nickel-based superalloys during long-term thermal exposure at 700 ℃ are investigated with a view to revealing the evolution of the microstructures and mechanical properties of the additively manufactured nickel-based superalloys.

IN625 powder with chemical compositions shown in Table 1 is used. Samples with dimension of 20 mm×20 mm×200 mm are prepared by the EP-M250 SLM system in nitrogen atmosphere. The processing parameters are chosen as follows: laser power of 200 W, scanning speed of 1000 mm/s, hatch spacing of 17 μm, layer thickness of 30 μm, and spot diameter of 100 μm. The scanning strategy involves rotation of 67°of the laser between two adjacent layers. All the samples for a mechanical property test are cut from the as-built samples using wire cutting machining as shown in Fig. 1. The heat treatment schemes used in the experiment are listed in Table 2. In order to compare the influence of the non-equilibrium microstructure in the as-built alloys on the evolution of the aging microstructures, a part of the samples are treated at 1200 ℃ for 1 h and followed by water quenching to eliminate the non-equilibrium microstructures. Subsequently, the as-built samples and the solution annealed samples are subjected to thermal exposure at 700 ℃ for 500, 1000, and 3000 h. Tensile tests are performed at room temperature under quasistatic loading (strain rate of 1 mm·min^-1). To observe the microstructures, all samples are first ground and mechanically polished. Then, electrolytic etching is employed at 10 V for 5-10 s in an electrolyte containing 10 mL HNO₃+ 30 mL HCl + 50 mL C₃H₈O₃. The microstructures are analyzed by optical microscope (OM) and scanning electron microscope (SEM) with energy dispersive spectroscopy (EDS). The average size of the precipitated phases is calculated use Image-Pro Plus 6.0 analysis software. JMatPro software is used to calculate the balanced phase diagram of the IN625 alloy. The temperature range is 600-1400 ℃ and the cooling rate is set to 10 ℃/s. The isothermal transformation phase diagrams at 700 ℃ and 750 ℃ are calculated.

The microstructural morphology of the SLM deposited IN625 alloy is shown in Fig. 3. From Fig. 3(a), we can see the traces of the U-shaped molten pool on the X-Z surface of the SLM forming part. The structure is mainly composed of columnar dendrites and cellular dendrites. These are the typical non-equilibrium structural characteristics of nickel-based superalloys formed by SLM. The EDS composition analysis result in the inset shows that the inter-dendritic region [zone 1 in Fig. 3(c)] has become a Laves phase rich in Nb and Mo elements. Fig. 3(b) shows the microstructural morphology of the SLM deposited IN625 alloy after solution treatment at 1200 ℃. After a high temperature solution treatment, the traces of the molten pool, the dendritic structure, and the Laves phase completely disappear, and the structure has undergone significant recrystallization, forming a uniform equiaxed structure and a large number of annealing twins. By comparing Figs. 5(a) and 5(d), it can be seen that after the initial thermal exposure of 500 h, dense needle-like δ phases are precipitated in the interdendritic regions of the deposited alloy and the original Laves phases are significantly reduced. While no δ phase is found in the solid solution alloy, a film-like precipitated phase is formed on the grain boundary, and a large number of γ″ phase particles are precipitated in the crystal. Until thermal exposure for 1000 h, needle-like δ phases are preferentially precipitated on both sides of the grain boundary in the solid solution alloy, and the δ phase grows from the grain boundary nucleation to the intragranular growth [Fig. 5(e)]. At this time, the δ phases in the deposited alloy interlace each other at an angle of about 60° to form a network structure. It can be seen from Fig. 7 that before the thermal exposure treatment, the ultimate tensile strength (UTS) of the SLM deposited alloy is 890 MPa, the yield strength (YS) is 620 MPa, and the elongation (EL) can reach 52%. After a solution treatment, the UTS value of the alloy is 887 MPa, the YS value is 390 MPa, and the EL value is as high as 64%. After an aging treatment, the tensile strength and yield strength of SLM deposited and solid solution alloys have been improved to varying degrees, while the elongation has shown a downward trend. After aging for 3000 h, the UTS and YS of the deposited alloy are increased by 36% and 51%, and the EL is decreased by 21%. The tensile strength and yield strength of the solid solution alloy are increased by 27% and 87% and the elongation is decreased by about 28%.

In the long-term thermal exposure process, the δ phase in the SLM deposited IN625 alloy preferentially nucleates in the interdendritic region, and after a solution treatment, the δ phase in the alloy gradually grows from the grain boundary nucleation to the intragranular growth. Comparing with solid solution alloy, the δ phase nucleation rate of the SLM deposited alloy is high but the coarsening rate is low, and the γ″ to δ phase transformation speed is fast. When thermal exposure for 1000 h, the transition from γ″ phase to δ phase is basically completed, the δ phase is concentrated on both sides of the grain boundary in the alloy after a solution treatment, and a large amount of γ″ phases are still distributed in the crystal. Due to the segregation of Si element at the grain boundary, a large number of Laves phases are formed at the grain boundary in the SLM deposited alloy, which causes the depletion of the δ phases near the grain boundary, and the grain boundary precipitated phase of the alloy after the solution treatment is mainly M₂₃C₆. After long-term aging, the strengths of SLM deposited and solid solution alloys are significantly increased, while the plasticity is reduced. But the tensile strength and yield strength of the SLM deposited alloys are significantly higher than those of the alloy after a solution treatment, and the elongation rate is still relatively high.

Nickel-titanium (NiTi) alloy is considered the most crucial shape memory alloy owing to its excellent superelasticity and shape memory effect. It can widely be used in aerospace, automobile manufacture, and biomedical fields. Compared with traditional metal materials, NiTi shape memory alloy has high damping capability owing to its martensitic transformation characteristics. However, fabricating NiTi alloy into structures with complex geometric configurations is difficult owing to its high wear resistance and superelastic properties. Nonetheless, as an emerging additive manufacturing technology, selective laser melting (SLM) has outstanding advantages in forming complex lattice structures with high geometric accuracy and surface finish. As one of the widely used structures in SLM structure design, the periodic lattice structure is often used as a buffer absorber given its light weight and high strength. Presently, most studies on forming NiTi alloy lattice structure using SLM focus on the elastic-plastic behavior under quasistatic conditions (tensile/compression). However, only few researchers have focused on the dynamic damping behavior of the lattice structure, especially the coupling between the damping characteristics of the material and the structural damping of the lattice structure. Therefore, the influence mechanism of the dynamic damping characteristics on the material-structure coupling must be investigated.

First, body-centered cubic (BCC) lattice structures with different rod diameters were modeled using the UG12.0 modeling software. BCC lattice structures with different rod diameters were fabricated using Ni_50.4Ti_49.6 shape memory alloy powder by SLM. Moreover, the first six orders of modalities and deformation modes of the BCC lattice structure were predicted using ANSYS finite element simulation. The first-order intrinsic frequencies and damping ratios of the BCC lattice structure with different rod diameters were obtained by shaking the table with sinusoidal sweeping experiments. The influencing factors of damping drop on the BCC lattice structure with a decreasing rod diameter were explored. The phase composition, chemical element content, number of defects, and morphology of the NiTi-BCC lattice structure at different rod diameters were analyzed through differential scanning calorimetry (DSC), oxygen-nitrogen-hydrogen analyzer, and microfocus X-ray computed tomography (Micro-CT).

Simulation and experimental results indicate that the first-order intrinsic frequency of the structure increased linearly as the rod diameter increased from 0.6 to 1.2 mm (Figs. 5 and 6). The enlargement of the rod diameter resulted in increased the volume fraction and elastic modulus of the structure with a certain rod length, thereby increasing the first-order intrinsic frequency of the structure. The decrease in the rod diameter contributed to the deterioration of the structural damping ratio from 0.020 to 0.012 (Fig. 7). To explain the decline in the structural damping ratio as the rod diameter decreases, phase transition temperature and chemical elemental analyses were conducted on samples with different rod diameters (Fig. 8, Table 1). First, the phase transition temperature gradually reduced as the rod diameter decreased. The decline of the laser scanning speed and the deterioration of the rod heat dissipation ability increased the actual laser energy input and absorption. Therefore, the effective oxidation led to the augmentation of the Ti element loss and the reduction of martensite content in the structure. The movable twin interface in the martensite phase is one of the critical damping sources of the NiTi alloy. Thus, the damping property of the SLM-NiTi alloy was reduced by the decreased phase transition temperature. Second, the porosity and number of pores in the BCC structure were characterized and analyzed by using Micro-CT. With the decrease in the rod diameter from 1.2 to 0.6 mm, the type of pores did not considerably change, but the pore number reduced dramatically (Fig. 9). On the one hand, the thermal energy dissipation resulted from the multiaxial stresses around the pores was attenuated by the decrease of pore number, which reduced the damping property. On the other hand, the interface area of NiTi matrix and pores were reduced by the decrease of pores number, thereby reducing the stress-assisted twin grain boundary motion in the martensite. Thus, the material damping was decreased. As mentioned above, the drop in material damping reduces structural damping.

The effects of rod diameter on the first-order intrinsic frequency and structural damping of NiTi alloys prepared by the SLM method were investigated using finite element analysis and experiments. Results show that the rod diameter had an essential effect on the first-order intrinsic frequency and damping of the BCC lattice structure. As the rod diameter increased, the overall stiffness of the structure increased and the first-order mode rose accordingly, which provide a basis to achieve multifrequency damping performance by controlling the rod diameter. The reduction of phase transition temperature and pore number, thereby reducing material damping, was the result of the decreased rod diameter. Therefore, the significant decrease in the structural damping of the NiTi-BCC lattice structure was attributed to the decreased rod diameter.

Metal additive manufacturing can be used to manufacture complex structural components that are difficult or even impossible to be produced using conventional methods. Recent development in constituent technologies has improved the understanding of process parameter-structure-property relationships for as-printed parts; 316L stainless steel (SS) is a face-centered cubic material, and the structure is not transformed when cooled to room temperature. Therefore, it is a good candidate material for analyzing the influence of heterogeneous microstructures on the mechanical performance of additive manufacturing (AM)-processed materials. Several studies have revealed that the strength and ductility of selective-laser-melted (SLM) 316L SS are higher than those of forged SS. This is because SLM parts have unique heterogeneous microstructures. Here, we review the recent SLM 316L SS, considering the process parameters, trans-scale structures, and mechanical properties. We provide a detailed review of SLM 316L SS with high strength and ductility and give insight into the future of this material.

First, defect formation mechanisms in SLM 316L SS are discussed. To summarize and compare the process-parameter-dependent relative density of as-printed samples, different energy density indices are adopted to calculate the resultant energy density under different processing conditions (i.e., different selective laser melting machines, spot diameters, and materials). Then, the melt pool evolutions with different process parameters are reported. We summarize the relationship between the melt pool geometry and crystallographic texture and present the melt pool morphology predicted through dimensionless analysis. Thereafter, the grain size and morphology, cellular structure, dislocation density, and nanoparticles of SLM 316L samples are discussed, focusing on the formation mechanism of cellular structures, followed by the presentation of the mechanical performance, including hardness, tensile properties, and corrosion behavior, of SLM 316L parts. Additionally, the effects of postdeposition heat treatment on the microstructures and tensile properties are also reviewed.

With an increase in various energy density indices, the relative density of the part increases first, remains constant, and then decreases (Fig. 2). Several dimensionless quantities, including R_HD, η_m, η_v, and KemLd*, are used to determine the minimum threshold for as-printed samples with high relative densities, and their values are 1.2, 2.6, 0.45, and 2.0, respectively (Fig. 3). The energy density imported into a powder bed has a significant impact on melt pool morphology (Figs. 5-8). However, contradictory results are observed when different energy density indices are used (Fig. 5), suggesting that there are limitations in using these indices as design parameters for selective laser melting. Comparing reported data for SLM and forged materials, it is observed that the grain size of SLM 316L SS is relatively large (Table 1). An anisotropic grain structure, namely, a checkerboard-like structure, is formed on the top surface, whereas a columnar grain structure is formed on the side plane of the as-built 316L SS, with the grain-size aspect ratio ranging from 1.4 to 15. Although the formation mechanism of the cellular structure is still not clarified, the structure plays a vital role in determining the mechanical performance of SLM parts. Nevertheless, as-printed materials have massive dislocation networks at the cell boundaries. Such cell structures with dislocations formed in SLM material are similar to the microstructure processed under severe plastic deformation processes. Many studies have reported that high dislocation densities, ranging from 10¹⁴ to 10¹⁵ m^-2, are obtained in as-built 316L SS, which contributes significantly to the enhanced tensile yield strength of SLM samples according to the Taylor hardening law. Recently, it has been reported that SLM 316L samples break through the strength-ductility tradeoff due to their hierarchical microstructure (Table 2 and Fig. 11). By tailoring laser process parameters, the melt pool shape, cell structure size, and size/content of nanoinclusions may change; hence, different strengthening mechanisms will be dominated at a certain printing process (Fig. 11). We infer that the hierarchically heterogeneous microstructure acts as a whole to resist tensile deformation under loading. Our experimental results collectively suggest that melt pool, grain, and cell structure boundaries are relatively weak regions in SLM parts, and the original grains of an SLM part are usually subdivided after tensile deformation (Figs. 12-13). Though the thermal stability of various microstructural features may differ (Fig. 14), independent of the process parameters and printing machines, 873 K is the temperature threshold above which SLM 316L SS exhibits classical strength-ductility tradeoff (Fig. 15).

In this study, we present a comprehensive overview of the evolution of the microstructures of SLM 316L SS from heterogeneous aspects. Unique microstructures, including the presence of crystalline grains, defects, melt pools, cellular structures, very high dislocation density similar to that of a severely plastically deformed material, and nanoinclusions, are formed in SLM 316L SS. Many studies have shown that the mechanical properties of SLM 316L SS are comparable with those of the wrought counterparts, though the mechanical performances may vary with process parameters and change locally within a part. Progress has been made in understanding SLM 316L SS, and the underlying strengthening mechanisms have been sufficiently revealed. Therefore, tailoring the structure and properties of SLM 316L based on scientific principles paves the way to AM metal parts with excellent mechanical properties. This review can serve as a valuable reference for understanding the current state of SLM 316L SS, the scientific gaps, and future research needed to advance this technology.

Additive manufacturing (AM), also known as 3D printing, is a disruptive technique and provides good compensation for conventional manufacturing methods. In AM, 3D parts are processed in a layer-by-layer manner following the designed 3D model and toolpaths. The rapid advancement of AM allows for an unprecedented design freedom for manufacturing complex, composite, and hybrid structures with high precision, which cannot be achieved using traditional fabrication routes. However, the AM process development and optimization usually requires costly and time-consuming trial-and-error experiments, thereby limiting the further application of AM. Machine learning (ML), as a new type of artificial intelligence technology, can accelerate the research and development in many aspects of AM; therefore AM has received extensive attention from academia and industry. With the assistance of ML, AM can be expedited and well optimized. Moreover, the relationship between the process parameters and achievable property of the alloys can be well revealed through ML, which is difficult using conventional methods. The ML technique has exhibited promising potentials in accumulating process optimization and novel alloy design for AM recently. Hence, this work reviews the research progress of the ML-assisted AM in the past decade.

In this paper, first, the ML technology used in AM is described. In general, ML methods can be divided into supervised learning, unsupervised learning, semisupervised learning, and reinforcement learning. According to studies, each ML method has many applications. Therefore, the typical applications for each ML method are introduced (Fig. 1). Second, the application of ML in the control and optimization of the AM metal materials, including the process monitoring and quality control, prediction of the process window, and optimization of the deposition toolpath, is discussed. By combining appropriate ML methods, the AM development processes can be considerably expedited and quality of the deposited parts can be stabilized. Third, the status of research and application of ML in the development of new alloys for AM is introduced. The correlative applications mainly include alloy composition design and prediction of microstructure and property of the deposited alloys. Recent years witnessed the growing research interests in the development of novel alloy materials used for AM (Fig. 8). Because it has been demonstrated that ML is an efficient way to accelerate the development period of novel alloy materials. With more available data accumulation, it can be expected that ML will have a broad prospect in novel alloy development for AM, which could create high-performance alloys for harsh industrial applications.

With the development of artificial intelligence and computer science, ML has been widely used in AM in recent years. The combination of ML and AM avoids a large quantity of trial-and-error costs, thereby reducing the development period of the AM. This work reviews the progress of machine learning-based AM process optimization and the novel alloy materials developments. The application of ML in the control and optimization of the AM includes the process monitoring, quality control, prediction of the process window, and optimization of the deposition toolpath (Fig. 10). The research and application of ML in the development of novel alloy materials based on AM include alloy composition design, microstructure, and property prediction. Finally, the future development trends of ML in the AM were outlined. In studies, the ML method usually focuses on a particular phase of the AM, which considerably limits the application and promotion of machine learning. The development of the generic ML algorithm for AM will further promote the application of ML in AM, which is also the critical research direction of machine learning-assisted AM in the future. For ML-based novel alloy materials developments in AM, several studies have shown that ML can effectively avoid the high costs of the traditional trial-and-error methods. However, ML requires a large number of databases to train the model. Therefore, the construction and development of an effective database is the precondition for ML. In recent years, a large amount of literature regarding AM of metallic materials has been published, which means a large amount of experimental data has been accumulated, and this is the fundamentals for the development of ML technology. With the development of practical data mining technology, the vast database will promote the development of novel metal materials for AM.

Ni-based superalloys are the ideal high-temperature materials due to their excellent oxidation resistance and microstructure stability at the elevated temperature. But the technical characteristic and limitation of traditional manufacturing techniques restrict the development and product of superalloy components. Recently, the laser additive manufacturing (LAM) technology provides an effective tool to fabricate the integral and complex components. However, the compositions of heritage alloys are designed based on the traditional techniques without considering the specifications of the LAM process. During the LAM process, the non-equilibrium characteristics of multiple thermal cycles, rapid heating and cooling rate and the localized microstructural evolution result in the metallurgical defects such as cracks, pores, and lack of fusion, which are difficult to be completely eliminated by optimizing the process parameters. Note that the Ni-based superalloys are developed from the Ni-20%Cr alloy, and the Ni-Cr-Al system can be regarded as the basic composition of the Ni-based superalloys. The basic composition plays a significant role in the further design of multicomponent alloys owing to the correlation to the microstructural stability, mechanical properties and weldability of Ni-based superalloys. Therefore, it is necessary to optimize the compositions by investigating the influence of Cr and Al contents on microstructures and properties of the basic alloys. In this paper, five representative basic alloys are designed based on alloying of binary Ni-20%Cr alloy with Al element, and the influence of composition on microstructures and properties of as-deposited alloys is systematically investigated. This research can be helpful to design the Ni-based superalloys which are fit for the LAM process.

Five basic alloys are first designed through the "cluster-plus-glue-atom" model, and then fabricated by the LAM process. The pure Ni plate is chosen as the substrate. The elemental powders of Ni, Cr and Al with purity (mass fraction) of 99.90%-99.99% and particle size of 50-150 μm are chosen as feedstock materials, and the powders are blended by a ball grinder for 10 h. The specimens are built on the LDM-800 additive manufacturing system using a strategy of bidirectional scanning, and the process parameters are optimized as laser power of 2 kW, scanning speed of 5 mm/min, laser beam diameter of 2 mm, overlapping rate of 50%, and powder feeding rate of 6.8 g/min. The LAM specimens are cut along the build direction for the microstructural and mechanical property analysis. The crystal structures of as-deposited alloys are identified through X-ray diffraction. The microstructural evolution and elemental distribution are analyzed by scanning electron microscope (SEM) and electron probe microanalyzer (EPMA), respectively. The precipitated phase is identified and investigated by transmission electron microscope (TEM) equipped with the selected-area electron diffraction. The microhardness is measured by the hardness tester, and the room compressive test is tested on a material testing machine. The continuous variable-temperature oxidation is performed on the thermal analyzer. In order to evaluate the weldability of alloys, three cross-sections of each alloy along the build direction are observed by optical microscopy and the solidification temperature range of alloy is measured by differential scanning calorimetry (DSC).

The matrices of Ni_75.0Cr_25.0 and Ni_75.0Cr_18.75Al_6.25 alloys are composed of the γ-Ni solid solution, while the γ′ phase begins to precipitate in γ-matrix when the Al content (atomic fraction) is higher than 6.25% (Fig. 1). Additionally, the α-Cr solid solution distributes between the columnar grains of the Ni_75.0Cr_25.0 alloy, and the amount of α-Cr solid solution increases with the increase of Al content. In the Ni_75.0Al_25.0 alloy, the α-Cr solid solution is replaced by γ′-Ni₃Al+ γ-Ni divorced eutectic (Fig. 2). The microhardness and strength of as-deposited alloys first slightly increase with the increase of Al content, and then sharply increase when the Al content is higher than 6.25%, owing to the precipitation of γ′ phase (Figs .4 and 5). But the ductility significantly decreases with the increase of Al content (Table 4). The continuous variable-temperature oxidation curve and the oxidation kinetics data show that the initial temperature of vigorous oxidation overall increases with the increase of Al content, while the total mass gain and mass gain rate change in an opposite trend (Fig. 7 and Table 5), and the improved high-temperature oxidation resistance can be attributed to the formation of Al₂O₃ oxide scale. However, the excessive Al content enlarges the solidification temperature range and deteriorates the weldability of alloys, and the large amount of pores and lack of fusions are formed in Ni_75.0Al_25.0 alloy (Fig. 10 and Table 7).

In this paper, five representative basic alloys are designed using the cluster model based on alloying of the binary Ni-20%Cr alloy with Al element. The influence of composition on microstructures and properties of the as-deposited alloys is investigated. The results show that the matrix structures of as-deposited alloys evolve from γ-Ni solid solution to γ′-Ni₃Al strengthening phase with the increase of Al content. Meanwhile, α-Cr solid solution distributing along grain boundaries changes from granule to long-chain in morphology while increasing its fraction, and is replaced by γ′-Ni₃Al+ γ-Ni divorced eutectic at 25.0% Al. The hardness and strength of as-deposited alloys increase with the increase of Al content due to the change in the strengthening mechanism from solid solution strengthening to precipitation strengthening, whereas the ductility decreases. The increase of Al content is beneficial for improving the high temperature oxidation resistance, but the excessive Al is deleterious to the weldability. Therefore, the Al content should be confined in a range of 12.5%-18.75% to make the basic alloys have a good match of mechanical properties, high temperature oxidation resistance and weldability.

Aiming at the demand of high-energy-absorbing and shock-resistant components for aerospace and transportation, metamaterials with energy absorption have been extensively studied, including truss-lattices, plates-lattices, triply periodic minimal surfaces (TPMS), and bionic metamaterials.

The truss-lattice metamaterials are the spatial structures formed by the multiple connecting rods with lattice or lattice-like arrangement, which possess high mechanical properties and energy absorption. The plate lattice metamaterials are the spatial structures in which the plate vertices replace the lattice nodes and have special plate arrangements. The plates can also generate multiple cavities through specific combinations, thereby achieving the functional effects of sound absorption and noise reduction. The TPMS metamaterials are the spatial structures that possess infinite periodic, continuous, and smooth surfaces in three independent directions. The surfaces have two disjoint regions in space. There are no sharp protrusions and depressions, which can decrease stress concentration. It is the best choice for manufacturing energy-bearing structures. Besides, the researchers have also found that its spatial configuration is similar to the structure of human bone, so it can be used to fabricate bone implants. The bionic metamaterials for energy absorption were first applied in 2000. The structures that exist in living organisms are the result of natural selection and evolution. They are known for their high specific energy absorption efficiency with small mass, which can be used for fabricating energy-absorbing components with impact resistance and energy absorption.

Traditional manufacturing technologies are difficulte to fabricate these metamaterials for their complex structures. The additive manufacturing (AM) technology is based on the principle of discrete slicing and layer-by-layer stacking to rapidly fabricate components, which possesses high manufacturing freedom. The above-mentioned technical characteristics make it an effective way to manufacture energy-absorbing metamaterials with complex structures. Thus, researching and developing metamaterials with energy absorption mean a lot.

The design of truss-lattice metamaterials has first changed from the ordinary regular truss arrangement to the gradient arrangement, and then the solid parts inside the unit cell are reasonably distributed through topological optimization of the computational models to maximize the mechanical properties and energy absorption. However, the design of truss-lattice metamaterials has complex geometric models, diverse mechanical properties, and multidisciplinary. Therefore, the main research direction is to develope efficient and specific mathematical models for the truss-lattice metamaterial design.

The plate-lattice metamaterial evolves based on the truss-lattice metamaterials, whose mechanical properties and energy absorption can be improved through some optimization strategies such as topological optimization. The combined method can realize optimization again, and the effect is remarkable. However, there are process constraints in the AM of plate-lattice metamaterials. The research direction of plate-lattice metamaterials is to optimize the metamaterial and develop a topology suitable for AM for maintaining excellent mechanical properties and energy absorption.

The TPMS metamaterials with heterogeneous and gradient structures are developed. The homogeneous structures is first developed to improve energy absorption, and then the combined TPMS metamaterials to efficiently control the mechanical properties and energy absorption appear. However, diversified TPMS metamaterials have various boundary distributions, so the combination method cannot be simply pieced together. Developing a corresponding mathematical model to achieve the smooth transition of multiple structures needs to be solved urgently.

Compared with the traditional metamaterials, the bionic metamaterials improve the energy absorption, and the bionic metamaterials also change from homogeneous forms to gradient forms to achieve high energy absorption. Later, the design of bionic metamaterials also changes from simply imitating their special macrostructures and microstructures to the combined design of biomimetic and lattice.

At present, with the continuous progress of material design and AM technologies, additive manufacturing of intelligent metamaterials has become a new research direction to ensure that the bionic metamaterials do not break during the process of impact resistance and energy absorption, and the designability and repeatability of bionic metamaterials are greatly improved. Bionic smart metamaterials are developing towards imitating shapes, imitating properties, and imitating functions. However, the development of ultra-high recoverable and deformable smart materials and the design of bionic metamaterials are the current study barriers to the metamaterials with energy absorption, and also the main development direction.

Aiming at the demand of high-energy-absorbing and shock-resistant components for aerospace and transportation, metamaterials with energy absorption have been extensively studied, including truss-lattices, plates-lattices, triply periodic minimal surfaces (TPMS), and bionic metamaterials.

The truss-lattice metamaterials are the spatial structures formed by the multiple connecting rods with lattice or lattice-like arrangement, which possess high mechanical properties and energy absorption. The plate lattice metamaterials are the spatial structures in which the plate vertices replace the lattice nodes and have special plate arrangements. The plates can also generate multiple cavities through specific combinations, thereby achieving the functional effects of sound absorption and noise reduction. The TPMS metamaterials are the spatial structures that possess infinite periodic, continuous, and smooth surfaces in three independent directions. The surfaces have two disjoint regions in space. There are no sharp protrusions and depressions, which can decrease stress concentration. It is the best choice for manufacturing energy-bearing structures. Besides, the researchers have also found that its spatial configuration is similar to the structure of human bone, so it can be used to fabricate bone implants. The bionic metamaterials for energy absorption were first applied in 2000. The structures that exist in living organisms are the result of natural selection and evolution. They are known for their high specific energy absorption efficiency with small mass, which can be used for fabricating energy-absorbing components with impact resistance and energy absorption.

Traditional manufacturing technologies are difficulte to fabricate these metamaterials for their complex structures. The additive manufacturing (AM) technology is based on the principle of discrete slicing and layer-by-layer stacking to rapidly fabricate components, which possesses high manufacturing freedom. The above-mentioned technical characteristics make it an effective way to manufacture energy-absorbing metamaterials with complex structures. Thus, researching and developing metamaterials with energy absorption mean a lot.

The design of truss-lattice metamaterials has first changed from the ordinary regular truss arrangement to the gradient arrangement, and then the solid parts inside the unit cell are reasonably distributed through topological optimization of the computational models to maximize the mechanical properties and energy absorption. However, the design of truss-lattice metamaterials has complex geometric models, diverse mechanical properties, and multidisciplinary. Therefore, the main research direction is to develope efficient and specific mathematical models for the truss-lattice metamaterial design.

The plate-lattice metamaterial evolves based on the truss-lattice metamaterials, whose mechanical properties and energy absorption can be improved through some optimization strategies such as topological optimization. The combined method can realize optimization again, and the effect is remarkable. However, there are process constraints in the AM of plate-lattice metamaterials. The research direction of plate-lattice metamaterials is to optimize the metamaterial and develop a topology suitable for AM for maintaining excellent mechanical properties and energy absorption.

The TPMS metamaterials with heterogeneous and gradient structures are developed. The homogeneous structures is first developed to improve energy absorption, and then the combined TPMS metamaterials to efficiently control the mechanical properties and energy absorption appear. However, diversified TPMS metamaterials have various boundary distributions, so the combination method cannot be simply pieced together. Developing a corresponding mathematical model to achieve the smooth transition of multiple structures needs to be solved urgently.

Compared with the traditional metamaterials, the bionic metamaterials improve the energy absorption, and the bionic metamaterials also change from homogeneous forms to gradient forms to achieve high energy absorption. Later, the design of bionic metamaterials also changes from simply imitating their special macrostructures and microstructures to the combined design of biomimetic and lattice.

At present, with the continuous progress of material design and AM technologies, additive manufacturing of intelligent metamaterials has become a new research direction to ensure that the bionic metamaterials do not break during the process of impact resistance and energy absorption, and the designability and repeatability of bionic metamaterials are greatly improved. Bionic smart metamaterials are developing towards imitating shapes, imitating properties, and imitating functions. However, the development of ultra-high recoverable and deformable smart materials and the design of bionic metamaterials are the current study barriers to the metamaterials with energy absorption, and also the main development direction.

Additive manufacturing (also known as 3D printing) is widely used in automotive, aerospace, shipbuilding, medicine, and other industries. Despite the name, most commercial 3D printing systems currently in use operate in 2.5-D mode, in which materials are accumulated layer upon layer in planes with a fixed direction (usually the opposite direction of gravity). Although this approach has lower hardware complexity and software development costs, it is faced with staircase effects, the need for support structures, and even a lack of manufacturability. Multiaxis 3D printing methods remove the problem of support structure required by traditional methods, showing extremely high manufacturing flexibility. It also offers a wide range of applications in support-free 3D printing of complex structures. This article reviews multiaxis support-free 3D printing process planning methods. These methods are classified as overhang structure decomposition, skeletonization, constraint optimization, curved layer decomposition, and inner or outer volume decomposition based on their capacity to deal with complex models. Following that, the issues and challenges of multiaxis support-free 3D printing are examined in terms of surface quality, overhanging area manufacturability, precision control, path planning, and so on. Finally, in view of the current problems and challenges, the prospects for multiaxis support-free 3D printing are discussed for future development.

The method based on overhang structure decomposition was the first proposed process planning method for multiaxis support-free 3D printing. The core of this method is to distinguish between the core (or buildable) and overhang (or unbuildable) volumes, which requires the overhanging feature to have strong concave edges or loops. To improve the ability to deal with complex geometric parts even further, the researchers obtained the centroid axis, or skeleton, from the geometric information of the input model by following the intrinsic characteristics of the part, and then used the skeleton to guide the decomposition into subvolumes without supports. This method can implement the support-free fabrication of multibranched or tree-like structures (as shown in Fig. 11), but it cannot effectively handle models without obvious skeleton features, implying that depending just on the skeleton for guided slicing is insufficient. Combining constrained optimization methods to minimize the area of the support structure yields a better decomposition result. Following that, a variety of methods, including the ant colony algorithm, beam-guided search algorithm, and downward flooding search algorithm, were proposed to find the optimal solution. In Table 2, the characteristics of constraint-based optimization methods are compared. However, to manufacture on a 3+ 2 axis platform, most of the above methods use a plane as both a dividing plane and a base plane. This planner layer-based constraint imposes constraints on the fabrication of more complex parts, and new approaches are attempted to fully utilize the flexibility provided by multiple degrees of freedom. The curved layer decomposition method attempts to divide the volume into a series of roughly equal-thickness surfaces while satisfying support-free and manufacturability constraints. There is a possibility of nozzle collision and local embedding if the resulting surface layer is concave. As a result, obtaining the convexly curved layers becomes the focal point of this method. Recently, an ellipsoid-based curved layer decomposition algorithm was proposed. Due to the ellipsoid’s convexity, the problem of nozzles’ local embedding was successfully avoided. However, this method is overly convoluted. To achieve a balance between algorithm, control complexity, and manufacturing efficiency, an inner or outer volume decomposition method was proposed, in which 5- and 2.5-axis depositions were applied externally and internally, respectively, to obtain a denser internal entity, which was important for metal parts. Table 4 summarizes the characteristics of the proposed multiaxis support-free 3D printing process planning method and their impact on the manufacturing process.

This paper also summarizes some current issues and challenges. First, previous work has focused on path generation with suboptimal surface quality under support-free and collision-free constraints. To improve surface quality, methods such as microcurved slicing, helical slicing, conformal printing, remelting, and hybrid fabrication have recently been proposed. Second, to improve the overhang area’s manufacturability, methods, such as changing Z-direction increments and tilting the nozzle, are used to expand the maximum permitted overhang angle. Third, while most of the current multiaxis 3D printing systems are in an open-loop state, the control quantities in the printing process are usually coupled with each other. Researchers try to smooth the uneven top surface caused by the variable height deposition strategy with the PI controllers. Finally, as illustrated in Fig. 24, the path generation is further optimized to meet the collision requirements while avoiding singularities as much as possible.

To summarize, the process planning methods of multiaxis support-free 3D printing for complex structures has yielded intriguing results, including the ability to decompose and print volumes with nonsharp edges, but it is still in its early stages. The following aspects are expected to be prioritized in the development of multiaxis support-free 3D printing methods. The first is more widespread process applications. At the moment, most methods are concentrated on the FDM process. To realize support-free printing of complex metal components, the process requirements and metal additive manufacturing characteristics must be further combined. The second is improved precision and performance printing. Feedback control is used to improve the robot’s positioning accuracy and to further optimize the motion trajectory to improve manufacturing accuracy. Simultaneously, by introducing requirements, such as mechanical properties into process planning, the mechanical properties of the parts will be improved even further. Additionally, functionally gradient structure printing of complex structures is a potential research topic. Another option is to use multiaxis hybrid manufacturing. The CNC machine has a high level of precision and efficiency in manufacturing. Combining with a CNC machine can improve manufacturing flexibility even further. There will also be a more open community. Hardware and software compatibility, as well as device openness, have an impact on multiaxis support-free process planning. However, some open source motion and process planning methods have emerged, and more opportunities are expected in the future.

Lattice structures have excellent mechanical properties such as high specific strength and high specific rigidity as well as outstanding functional characteristics such as vibration reduction, heat dissipation, sound absorption, and electromagnetic shielding. They are widely used in aerospace, biomedicine, and transportation fields. However, the materials used to form lattice structures are mostly stainless steel, Ti6Al4V and AlSi10Mg, which cannot meet the requirements of some complex intelligent components for shapes, performances and functions that can change over time or space. The NiTi shape memory alloy, a new type of smart materials, has excellent super-elasticity, shape memory effect, excellent corrosion resistance and wear resistance, and other functional characteristics. It can achieve a shape recovery through certain external stimuli after deformation, so that it meets the requirements of controllable deformation and regulable performance. Compared with NiTi bulk materials, NiTi lattice structures have low elastic modulus and density, large deformation ability, and can adjust the mechanical properties by designing the size, shape and distribution of holes. The unique performance of NiTi lattice structures makes them have a wide range of application prospects in the aerospace field. For aerospace components, weight reduction is an eternal theme, but it is not clear how to further improve the lightweight characteristics of NiTi lattice structures. In this paper, the NiTi alloy body-centered tetragonal (BCT) hollow lattice structure is proposed, which possesses the advantages of high load-bearing capacity of the BCT lattice structure and the excellent super-elasticity and shape memory effect of NiTi alloys. That is, keeping the outer diameter of the strut constant and hollowing out the inner part of the strut are used to achieve the purpose of reducing the weight of the structure.

In this paper, NiTi pre-alloyed powder is used as the raw material to prepare the BCT hollow lattice structures by laser powder bed fusion (LPBF). The surface morphologies and microstructures of the formed samples are observed by scanning electron microscope (SEM), the phase transition behavior of BCT-100 is characterized by differential scanning calorimeter (DSC), and the phase compositions of BCT-100 are determined by X-ray diffractometer (XRD). The finite element simulation method and the uniaxial compression experiments are applied to analyze the influence of mass coefficient on the compression performance of structures. Cyclic compression-thermal recovery experiments are carried out to reveal the influence mechanism of mass coefficient on the shape memory effect of NiTi lattice structures.

The lattice structures manufactured by LPBF have high forming accuracy and relative density, and no defects such as cracks and irregular large-size holes are found (Fig. 4). The lattice structure with a mass coefficient of 100% has the best bearing capacity, the first maximum compressive force can reach 191.73 kN, and the corresponding deformation rate is 0.22. When the mass coefficient of the structure is reduced to 75% from 100%, the first maximum compressive force is 89.80 kN, and the bearing capacity is reduced by 53.16%. At that time, the compression deformation capacity is not weakened, and the deformation rate is still up to 0.21 (Fig. 7). Therefore, in the non-primary load-bearing members, the struts of lattice structures can be hollowed to reduce the mass coefficient of structures by 25% while maintaining the deformation capacity of components. However, further reducing structural mass coefficient weakens the bearing and deformation capacity. The shape memory effect of the lattice structure with a mass coefficient of 75% is the best, and the recovery rate can reach 98.92% in the first cycle (Fig. 9).

The forming quality of lattice structures fabricated by LPBF is high, but there are still dimensional deviations caused by solidification shrinkage, powder sticking, and staircase effect. The M_s (Martensite transformation start temperature) and A_f (Austenite transformation finish temperature) of BCT-100 are 11.02 ℃ and 33.72 ℃, respectively. The phases of components are mainly composed of B2 and B19′ at room temperature, and the B2 phase occupies the dominant position. The appearance of B19′ phase in components is related to the phase transition of B2 phase induced by thermal stress produced by LPBF. The experimental compression force-deformation rate curves of the four structures can be divided into five stages: elastic deformation of austenite phase, stress induced transformation of austenite phase into martensite phase, elastic deformation of martensite phase, plastic deformation of martensite phase, and fracture stage. The lattice structure with a mass coefficient of 100% has the best bearing capacity, the first maximum compressive force can reach 191.73 kN, and the corresponding deformation rate is 0.22. When the mass coefficient of the structure is reduced to 75% from 100%, the first maximum compressive force is 89.80 kN, and the bearing capacity is reduced by 53.16%. At that time, the compression deformation capacity is not weakened, and the deformation rate is still up to 0.21. When the mass coefficient is further reduced to 50% (BCT-50), the first maximum compressive load and deformation capacity are reduced by 81.52% and 36.36%, respectively, that are 35.43 kN and 0.14. The shape memory effect of lattice structures formed by LPBF is good. In the first cycle, BCT-75 has the best shape memory effect and the highest recovery rate can reach 98.92%. The recoverable rates of BCT-93 and BCT-100 are slightly low, which are 97.71% and 96.77%, respectively. The shape memory effect of BCT-50 is the worst and the recovery rate is only 94.94%. In the last two cycles, all components can achieve a fully recovery.

Metal additive manufacturing technology has several advantages such as efficient formation, short processing cycle, and cost effectiveness. Components with complex spatial structures can be produced via additive manufacturing, thereby overcoming the limitations of traditional manufacturing. Therefore, this technology is favored by the automotive, aerospace, and medical equipment industries. However, possible internal defects, such as lack of fusion, cracks, and holes during the forming process of additive manufacturing, limit its promotion and wide application in the industry. Furthermore, the microstructure of the components changes significantly with the variation in the laser power, process approaches, and scanning parameters during the additive manufacturing process. Furthermore, the stability of the phase and characteristic microstructures are affected by the protective gas, while controlling the surface topography. Therefore, the quality control of metal additive manufacturing products, particularly online monitoring, is of great strategic significance. Several approaches of nondestructive evaluation of flaw inspection and material characterization, such as X-ray computed tomography, fluorescent penetrant inspection, and ultrasonic testing, have attracted much interest. Particularly, ultrasonic testing is one of the most commonly used nondestructive methods for detecting internal defects. Compared to traditional ultrasonic nondestructive testing technology, laser ultrasonic nondestructive testing has the advantages of no-contact, high sensitivity, and suitability for harsh environment, which can realize rapid online monitoring.

In this paper, the characteristics of metal additive manufacturing and nondestructive testing on additive manufacturing are briefly introduced, highlighting the fact that applying laser ultrasonic testing on metal additive manufacturing has great strategic significance. Then, two kinds of laser ultrasonic mechanisms are analyzed: thermoelastic mechanism and ablation mechanism. Under the laser ultrasonic simulations and experimentations, the thermoelastic mechanism is chosen without destroying the integrity or performance of the additive manufacturing components. Next, the finite element simulation studies on laser ultrasonic detection are introduced. Based on the finite element method (FEM), the complex models can be processed and the global numerical solution can be obtained by solving heat conduction and thermoelastic equations. Afterward, the principle of laser ultrasonic nondestructive testing and the testing system are introduced, in which common detection methods on laser ultrasonic are listed and briefly explained. Some improving methods on laser ultrasonic testing systems are also discussed. Yan et al. proposed an experimental method of a no-contact all-optical laser ultrasonic detection and built the optical differential detection system using the beam deflection technique, which improved the antinoise ability of the optical path. Finally, the application progress of laser ultrasonic nondestructive testing of metal traditional and additive manufacturing material at domestic and foreign industries is systematically summarized. Moreover, we have analyzed the research progress—both home and abroad—for reference. Studies on laser ultrasonic began earlier abroad. One example is the study by Pierce et al. (1993), who successfully used a pulsed Nd∶YAG laser to excite ultrasonic waves in metal aluminum blocks and increased the laser ultrasonic signals by modulating the frequency of laser source, thus excavating the great prospect of laser ultrasonic in nondestructive testing. In comparison, relevant research in China only started in 2006. For instance, Shen et al. detected rectangular metal aluminum blocks with an artificial surface defect with depth of 0.71 mm and width of 2.00 mm by constructing an optical differential detection system based on the beam deflection method, accurately locating the surface defect position. In addition, laser ultrasonic can be used for detecting the defects and measuring other significant parameters or monitoring other important processes, such as characterizing the elastic modulus, measuring residual stress of additive manufacturing alloy parts, and monitoring the changes in additive manufacturing process like recrystallization.

Several studies have shown the feasibility of laser ultrasonic nondestructive testing on metal additive manufacturing. Given that the effect of additive manufacturing process parameters on quality has been widely studied and reported, developing a link between controllable process parameters and the required process characteristics to support feedforward and feedback control is the best way to achieve the goal of its application in future control systems. However, various challenges still exist. For example, the relationships among the parameters must be identified, including the parameters measured from experiments and those important parameters for monitoring and characterizing but cannot be obtained directly. Furthermore, an algorithm for quickly identifying different defects should be explored, and a corresponding feedback control scheme must be established to improve the quality of the additive manufacturing components. Moreover, the feedback data obtained from the entire online monitoring of the additive manufacturing process are expected to be massive and difficult to deal with. Thus, we should find an algorithm that can efficiently process these data, detect the anomalies in real time, and provide corresponding feedback to the closed-loop detection system, thus allowing us to control the manufacturing engineering of the workpiece in real time and ultimately improve the quality of finished products.

Laser additive manufacturing (LAM) technology is not only a representative technology of intelligent manufacturing but also the most popular technology used for the additive processing of metallic materials at the moment. The metal products fabricated using LAM have many advantages, including high geometric freedom, good dimensional accuracy, and excellent performance and quality. They have been widely used in the aerospace, biomedical, and defense industries. Because LAM processing includes a high-temperature laser-heating process, the processing area is often shielded with inert atmospheres to avoid the air contamination of the metallic materials. Recently, LAM technologies utilizing reactive atmospheres, such as N_2,Ar-O_2,Ar-N_2,and Ar-CH_4,have rapidly emerged with remarkable achievements in improving the mechanical properties of various metals. It solves a long-lasting problem: the difficulty of modifying the properties of internal materials through atmospheric modification, which might be considered a historic breakthrough. Simultaneously, a new technological path has emerged: modifying feedstock materials before LAM processing utilizing active atmospheres. This study covers recent local and international research advances in the technology of reactive atmospheric LAM, which is a promising developing technology. The achievements are the result of two major technical directions: reactive atmospheric protected LAM processing and reactive atmospheric modified LAM feedstock materials. This review is concluded by analyzing the impacts of this new technology on representative metals such as steel, titanium alloy, aluminum alloy, and high-entropy alloy. In addition, it analyzes and compares material advancements in terms of mechanical properties, microstructural changes, and forming quality, as well as underlying mechanisms. At the same time, this paper discusses the current challenges and opportunities of the reactive atmospheric LAM technology.

In Section 2, this review critically examines the literature involving LAM processing shielded via reactive atmospheres. Figs. 2-11 summarizes and illustrates the effects of reactive shielding gases. Powder-bed-based LAM has received the most research interest due to its success in processing atmospheric-reinforced Ti, Ti alloys, and high-entropy alloys. The abundant and stable gas supply in the sealed processing chamber has been identified as a key advantage for achieving successful and homogeneous metal-gas in-situ reactions. The homogeneous distribution of solute atoms as well as compound precipitates can significantly improve the strength of metals and even boost. A number of attempts at direct-deposition LAM using reactive shielding atmospheres have also been made. However, the results are generally inferior to those achieved via powder-bed-based LAM, owing to insufficient gas-metal reactions during direct-deposition LAM processing. Section 3 introduces feedstock material modification using reactive atmospheres. This atmospheric feedstock modification technology can exert a significant influence before LAM processing, such as drastically increasing N concentration in austenitic stainless steels and improving Cu’s laser absorptivity via surface nitridation/oxidation of feedstock powders (Figs. 12 and 13). The method is particularly applicable to less-reactive metals and can avoid the unfavorable disturbance caused by metal-gas in-situ reactions during LAM processing. Finally, Table 2 summarizes representative achievements of reactive atmospheric LAM fabricated metals. Section 4 delves deeply into and elucidates the mechanisms of atmospheric modification, such as solute element diffusion, compound precipitate formation, influence on solid-state phase transformations, and strengthening mechanisms.

Recently, the application of reactive atmospheres in the LAM of metallic materials has emerged and developed rapidly. This technology has shown great potential in modifying the properties of metals during the fabrication of products. N, O, C, and other alloying elements from atmospheres have been successfully added to a variety of metals, including Ti alloys, steels, Al alloys, and high-entropy alloys. Remarkable strengthening (up to 40%-100% increment) has been achieved in Ti, Ti alloys, and austenitic stainless steels. Furthermore, the effects of reactive atmospheres may be accurately regulated by altering the contents of the atmosphere as well as laser parameters. This review has also proposed several prospects for further clarifying the mechanisms of atmospheric alteration and extending the applications of this technology. It is possible to learn more about the mechanisms at the subnano/atomic scale using in-situ microscopy and X-ray diffraction analysis. Additionally, the reactive atmospheric LAM has the ability to fabricate gradient materials by adjusting atmosphere or laser during processing. As a result, materials with different properties can be deposited at designated positions.

Selective laser melting (SLM), also known as laser powder bed fusion (LPBF), has broad application prospects due to its excellent performance and high fabrication accuracy. SLM technology is developing toward multi-beam, multi-material, high quality, and high-efficiency manufacturing. However, defects, such as internal metallurgical defects, and residual stress, restrict its process reliability and repeatability. Investigating the laser-matter interaction and its internal relationship with forming defects is expected to provide a scientific theoretical basis for SLM to achieve stable forming with fewer defects. This work studies the interaction between vapor plume and spatter behavior during SLM based on high spatial-temporal resolution in situ imaging systems.

The high spatial-temporal resolution in situ imaging system consists of a high-speed video camera (Phantom 2012, Vision Research, USA), a synchronized pulsed high-power diode laser light source (CAVILUX HF, Cavitar, Finland), and a zoom lens system (12X Zoom, Navitar, USA). In this work, imaging was performed at up to 10⁵ frames/s with a 1 μs exposure time. In front of the camera lens, a Thorlabs narrow bandpass filter with a wavelength of 808 nm was placed. With a full field of view of 2.0 mm×1.8 mm, the camera setup angle between the object plane and processing plane was 45°. The spattering trajectory is the projection on the plane parallel to the CMOS detector. An image filtering algorithm used to increase the sharpness of the spatter imaging. The sizes, numbers, ejection angles, and ejection velocities of the spatter trajectory were quantified using ImageJ 1.53.

(1) Droplet spatter (ejected from the "liquid base" of the molten pool) driven by the metal vapor recoil pressure. When the laser energy density is 27.5 J·mm^-3, the ejection of the liquid column in the molten pool is not noticeable. As the laser energy density increases to 59.0 J·mm^-3, the protrusion and the droplet column are formed at the rear of the molten pool. The super-threshold ejection occurs when the laser energy density E_v=90.4 J·mm^-3 (Fig. 2), and the droplet column ejection process can be divided into three stages (Fig. 3): the rising stage (Section ①); the oscillation stage (Sections ②-④); and the falling stage (Section ⑤). The hemispherical protrusion radius of the droplet column R_dc ranges from 16.2 μm to 52.5 μm. Therefore, 3.1-5.6 m·s^-1 can be estimated as the vertical component of the threshold velocity u_{th_l} (Eq. 2). The peak value of the ejection velocity vertical component u_v (20.0 m·s^-1) is greater than u_{th_l}, indicating that the protrusion ejects completely at approximately t=70 μs. With the high-speed ejection and oscillation, the droplet column broke at least four times in 500 μs, and the oscillation frequency is approximately 8 kHz.

(2) Powder spatter (ejected from the "solid base" of the substrate) driven by the metal vapor-induced entrainment. Under the influence of metal vapor entrainment, the typical spatter particle ejection process can be divided into four stages (Fig. 4): ① agglomerates on the substrate (P1, P2, and P3); ② enters the entrained inert gas flow (P1 and P2); ③ enters the metal vapor plume (P2); and ④ falls back into the powder bed or enters the circulating inert gas flow. The ejection velocities of small-sized spatters are relatively high, and the deviation of ejection angle is small (Fig. 5), implying that their average ejection angle can represent the metal vapor propagation direction (134.1°±4.1°). For particle P1, the average ejection angle is 95.2° from t=550 μs to 1100 μs, indicating that P1 does not enter the vapor plume. On the contrary, for particle P2, the average ejection angle is 130.8° (very close to metal vapor propagation direction) from t=980 μs to 1100 μs, implying that P2 enters the vapor plume. Furthermore, under the influence of the metal vapor plume, the geometry of P2 vibrates, indicating the heat and momentum transfer from the metal vapor plume to P2.

(3) Spattering dynamics driven by the metal vapor recoil pressure. As spatter particle P9 enters the laser beam region, when the temperature steeply increases, the upper parts of the particle begin to boil (incandescent state) at t=350 μs, and the particle’s trajectory changes drastically by the vapor recoil pressure (Fig. 7). Within a very short period (t=360-390 μs), the ejection angle changes from 86° to -86.9°, and u_v changes from 1.9 m·s^-1 to -3.3 m·s^-1 (Fig. 8). Therefore, the effective recoil pressure can be calculated as 2.7×10⁴ Pa, which is consistent with the theoretical results of the Clausius-Clapeyron equation. Moreover, according to the SLM equipment parameters adopted in this work, it is calculated that the laser focal depth is 3.4 mm. Hence, the spatter (h= 400 μm) ejected above the molten pool (h=0 μm) is located at the laser focal depth. This suggests that the metal vapor recoil on the surface of the molten pool can be derived from the metal vapor recoil on the spatter P9 by in-situ measurements with high spatial-temporal resolution.

We reveal the ejection process and the melt escape velocity threshold of the droplet spatter from the "liquid base" of the molten pool. The ejection process of powder spatter from the "solid base" of the substrate is also revealed as well as the gas velocity threshold.

The spattering during SLM was used as a tracer particle to obtain in-situ measurements of vapor recoil pressure through the "vapor-solid" interaction between metal vapor and powder particles, and the experimental measurements were in good agreement with the theoretical calculations. This lays the foundation for future research on the "vapor-liquid" interaction between metal vapor and molten pool (e.g., vapor recoil-induced keyhole and other phenomena).

The physical essence of SLM additive manufacturing is "laser-matter interaction." The research will be expanded in the future to include "laser" and "material, " e.g., the multi-beam laser-matter interaction during SLM and the forming mechanism of multi-material and difficult-to-form materials, which is expected to provide the scientific theoretical basis for SLM technology to achieve stable forming with fewer defects.

Insufficient laser energy input typically leads to the well-known balling effect and results in potential porosities. The basis of selective laser melting (SLM) process is the interaction between the laser and feedstock powder which determines the thermo-fluid dynamics of melt pool and the final quality of SLMed parts. An in-depth insight of laser energy deposition during the SLM process is critical for the process’ optimization and defect elimination. However, due to the difficulty of direct observation of laser reflections, the current understanding of laser irradiation mechanisms is still vague and unclear. Numerical modeling is an effective way to simulate the heat and mass transfer during the SLM process at mesoscopic scale. However, the laser heat source models used in existing literature, such as volumetric heat source and vertical ray heat source, have rarely considered the authenticity of laser-material interaction and have neglected the behavior of multiple reflections and absorptions of a laser in SLM. Therefore, a high-fidelity mesoscopic CFD model coupled with the ray-tracing method has been established in this work using which the correlation between melt pool dynamics and laser irradiation behaviors is well visualized and studied.

The numerical simulation is based on the VOF two-phase flow model that fully considers several physical phenomena such as solidification/melting, surface tension, Marangoni effect, recoil pressure, evaporation heat loss. A ray-tracing algorithm (Fig. 1) is developed to describe the laser-material interaction using user defined function in the commercial software FLUENT. A rain drop method is utilized to generate a layer of randomly packed powder particles based on the commercial software EDEM. The Cu-Cr-Zr single track is fabricated by a commercial SLM system (XDM 250, XDM Co., Ltd, China). The surface track morphology in top view and cross-sectional view of the metallographic diagram is observed with a digital microscope to compare with the simulation results. In-situ measurements of the effective laser absorptivity (Figs. 4 and 5) are carried out based on the calorimetric method to validate the modeled laser absorptivity.

Under the laser power of 430 W and scanning speed of 0.6 m/s, the simulated track width is in the range of 105.4-133.2 μm with an average of 122.4 μm. The experimental result on the other hand is in the range of 91.1-140.9 μm with the average of 110.1 μm. Additionally, the simulated and experimental outcomes also match in terms of the track depth (Fig. 6). A continuous melt track is formed with 430 W laser power and 0.6 m/s scanning speed, while a distorted and broken track is observed with 330 W laser power and 1.2 m/s scanning speed (Fig. 7). This can be attributed to the fact that the insufficient energy input leads to the generation of a smaller volume of melted liquid, and the melted liquid tends to aggregate with surface tension at play. By visualizing the laser ray trajectories, different laser reflection behaviors are observed in both cases. Multiple laser reflections are observed in the depression region for the continuous case, while fewer reflections are observed as the balling effect occurs which implies a reduction in global laser absorptivity (Fig. 8). The effective laser absorptivity measurement has good agreement with the simulated global absorptivity using the ray-tracing method with a low relative error of 7.1% (Table 4). In addition, the vertical reflections on the emerging exposed substrate greatly contribute to the reduction of global absorptivity and lead to an intense oscillation (Figs. 8 and 10).

In the present study, a novel mesoscopic CFD model is established to simulate the melt pool dynamics and laser reflection behaviors during SLM. The implementation of the ray-tracing method can well reproduce the laser reflection and absorption behaviors with the evolution of melt pool thermodynamics which provides an in-depth insight into the energy coupling mechanism during the SLM process. The depression region induced by recoil pressure has a light trapping effect which promotes multiple laser reflection and absorption. The exposed substrate surface due to the occurrence of balling effect greatly weakens the light trapping effect and reduces the laser absorption of the powder bed. The global absorptivity under the balling effect has characteristics of violent fluctuation. The unstable melt pool has an adverse effect on energy coupling between the laser and powder bed. The dynamic balance between the melt pool state and laser absorption is the key to forming a single track of good quality.

Considering data quantity, reliability, cost and other factors, it is an effective method to monitor the additive process of powder bed fusion-laser (PBF-L) by using photodiode to collect radiation information of molten pools in engineering practice. The on-line radiation monitoring of molten pools based on photodiode can collect a large amount of information in real time reflecting the internal conditions of molten pools in real time, and thus plays a very important role in predicting and controlling the quality of molten pools. The key is to continuously collect the radiation information of molten pools and to extract and analyze the data characteristics in a wide time scale (from several seconds to tens of hours) by the statistical methods, so as to realize the stable state analysis and quality prediction of the forming process. In PBF-L, all process factors (laser parameters, scanning strategy, powder state, and air flow protection) eventually influence the change of radiation signals of the molten pool, among which as for the main effects of laser power, scanning speed, preheating temperature and others on temperature and radiation intensity of a molten pool, a lot of research has been done. However, in addition to radiation intensity, the radiation signal change of the molten pool over time also contains a lot of other technological information and process stability information and it is worth further digging and studying. In the actual production process, PBF-L is a very complex process and the forming quality is still uneven even when the same equipment and process parameters are monitored and controlled. In order to ensure the repeatability and process consistency of additive manufacturing, the influence of process factors such as floor height, substrate position and scan line angle on the radiation signals of the molten pool is analyzed, and it is found that these characteristics can be used as an important indicator of quality control in the future.

In this paper, the forming experiment of K438 superalloy powder is carried out, and the photodiode is used to collect the radiation signal of the molten pool in the process of PBF-L. First it is analyzed and pretreated. Then it is segmented corresponding to the sample one by one. Statistical methods are used to process the segmented data, and the mean and standard deviations are selected as indicators to evaluate the signal characteristics, and the influence of process factors such as floor height, substrate position and scan line angle on the radiation signals of the molten pool is finally analyzed.

The influence of the process factors of PBF-L on the radiation intensity of the molten pool can reflect some classical laws. With the increase of layer height, the mean radiation intensity of the molten pool shows an overall increase trend, and the mean radiation intensity of the printing layer No.190 increases by about 6% (Fig. 4). In the wind field, the intensity waveform of the upwind molten pool shows the feature of "right deviation" , while that of the downwind molten pool shows the feature of "left deviation" (Fig. 6). The influence of relative position in the wind field on the radiation intensity of the molten pool is that the farther away from the air outlet, the greater the molten pool strength, and the closer to the air outlet, the smaller the molten pool strength (Fig. 8). In the scanning strategy, the scan line angle is exactly consistent with the incline angle of the sample (based on the horizontal direction), in which the mean intensity is the maximum (Fig. 11).

Through the processing and analysis of the intensity signal data of the molten pool, the influence law of some process factors on the radiation intensity of the molten pool is obtained, and the relationship between the radiation intensity law of the molten pool and the physical mechanism is established. The analysis shows that the effect of layer height on the radiation intensity of the molten pool is mainly due to heat accumulation among layers. In the wind field, the influence of wind direction and relative position on the radiation intensity of the molten pool is mainly the smoke masking effect. The effect of scan line angle on the radiation intensity of the molten pool is mainly due to the laser duty cycle. Through establishing the correlation between law and physical mechanism, it is clearer that in the practical engineering application and monitoring closed-loop control, it is necessary to consider the influence of layer height, scanning strategy, wind field conditions and other factors on the radiation intensity of molten pools. And its typical rules show that these characteristics can be used as an important indicator of quality control in the future.