The TC17 alloy has excellent mechanical strength, fracture toughness, and corrosion resistance and is primarily used in the manufacturing of aero engine disk components. Directed energy deposition (DED) can achieve specific location repairs of damaged parts and 3D printing of complex and large parts. DED of TC17 alloy has been successfully applied to such parts as integral blade disks. However, the low hardness and poor wear resistance of TC17 alloy make it susceptible to fatigue fracture under the harsh use environment of aero engines, which limits its application in the aerospace field. At present, the problems of poor bonding strength, introduction of new defects, and inability to meet various performance requirements in the spraying of hard coatings on titanium alloy surfaces present potential risks and limitations in aerospace applications. Optimized heat treatment has unique advantages in solving these problems. Optimized heat treatment can adjust material properties by adjusting the microstructure. Several studies have been conducted on improving the tensile properties of TC17 by optimizing the heat treatment process, but further research is needed on whether optimized heat treatments can improve the wear resistance of deposited TC17 alloy. Therefore, in this study, the effects of annealing, post-annealing solid solution, and post-annealing solid solution aging treatments on the microstructure, hardness, and tribological properties of deposited TC17 alloy are investigated. The evolution of the structure during heat treatment and the comprehensive wear behavior under 20 N dry sliding wear are analyzed. The results provide a reference for optimizing the tribological properties and heat treatment process of TC17 alloy.

The experimental material is TC17 spherical powder, with an average particle size of 66.6 μm (Fig. 1). The experiment uses the semiconductor laser to generate lasers, and the robot, equipped with coaxial powder feeding, deposits the TC17 powder on the polished TC4 substrate in the argon environment (Fig. 2). The process parameters are optimized: laser power of 1600 W, scanning speed of 10 mm/s, powder feeding rate of 11 g/min, overlap rate of 45%, and center protection gas flow rate of 11 L/min. A TC17 alloy sample (size of 75 mm×35 mm×12 mm) is obtained using an N-type scanning path. Samples are taken along the direction of laser deposition, with a sample size of 6 mm×6 mm×6 mm. The TC17 deposition samples are subjected to pre-annealing, annealing solid solution treatment, and annealing solid solution aging treatment (Fig. 3), and the microstructure and wear properties of the TC17 deposition and heat-treated samples are characterized by the X-ray diffraction (XRD), energy-dispersive X-ray spectroscope (EDS), scanning electron microscope (SEM), hardness tester, and pin-disk-type friction wear testing machine.

The experimental results indicate that the deposited TC17 alloy consists of α and β phases, displaying a basketweave structure (Fig. 6). After annealing at 840 ℃, some of the fine α phase dissolves because of the high temperature. After solid solution treatment at 800 ℃, the primary α phase (α_P) in the grain interior gradually flattens. During the annealing and post-annealing solid solution treatment stages (Fig. 7), affected by the diffusion rates of different elements, the grain boundary α phase (α_GB) is divided into continuous α_GB and discontinuous α_GB. A phase-free zone (PFZ) appears around the continuous α_GB because of the insufficient concentration of α stabilizing elements at the low-angle grain boundary. After aging at 580 ℃ based on the post-annealing solution treatment, a large amount of fine needle-like secondary α phase (α_S) precipitates, and the PFZ disappears. When the temperature rises to 630 ℃, some of the ultrafine α_S redissolves in the β matrix. After aging at 680 ℃, PFZ reappears with only partial coarsening α_S interspersed between α_P (Fig. 9). The α_S precipitates inside the β grains, and the randomness of orientation makes the size and quantity of α_S very sensitive to changes in aging temperature. After solution treatment, the average microhardness reaches 425.45 HV, which is higher than the hardness in the as-deposited state (Fig. 12). This is attributed to the precipitation and growth of α_P, which increase the volume fraction of the phase, thereby improving the microhardness. After aging treatment, the hardness is further increased, reaching its highest value (~486.93 HV) after aging at 580 ℃. This is caused by the large amount of α_S precipitation, which achieves the strongest dispersion-strengthening effect. The wear test results show that the wear properties of the heat treatment state are superior to those of the deposition state (Fig. 13). Table 3 shows the maximum wear widths and depths in different states. The tribological properties are the best after aging at 580 ℃, which is attributed to the significant increase in hardness and a large number of secondary phases inhibiting dislocation movement and crack expansion. After heat treatment, a variety of wear mechanisms coexist, and the wear rate and wear morphology depend on the changes in the microstructure and oxide layer.

A TC17 sample is prepared by DED and then heat-treated. The changes in phase composition, microstructure, microhardness, and tribological properties during heat treatment are studied. The results show that the main microstructure evolution of TC17 alloy during heat treatment includes the growth and coarsening of α_P; the continuous and discontinuous growth of α_GB, where the width of PFZ is positively related to the continuous growth of α_GB; and the precipitation and growth of α_S, where, as the aging temperature increases, some α_S dissolves and some grows to have a clear phase boundary with the β phase. The hardness after heat treatment is higher than that of the deposited state. In the solid solution stage after annealing, the hardness increases with the increase in α phase volume fraction. After further aging, the precipitation of α_S achieves dispersion strengthening, and the strengthening effect weakens with the dissolution of α_S. The hardening effect is higher at 580 ℃, and the hardness is increased by 20.1% compared with that of deposition state. The tribological performance after heat treatment is better than that in the deposition state. The increase in hardness, secondary phase precipitation, and hard oxide formation are the main reasons for the improvement of the tribological properties. Optimal wear resistance is achieved in the heat treatment systems of 840 ℃/1 h, air cooling+800 ℃/4 h, water quenching +580 ℃/8 h, and air cooling, with a friction coefficient (wear rate) of 0.422 (0.0451 mg/m).

Compared with traditional manufacturing methods, selective laser melting (SLM) can form complex components. As 316 L stainless steel has excellent mechanical properties, several studies on additive manufacturing of 316L stainless steel have been conducted. However, current research on the impact toughness of 316L prepared by SLM based on the heat treatment temperature is insufficient, and the effect on anisotropy has not been reported. Therefore, in this study, we use the SLM method to prepare 316L impact components with different forming orientations and compare the impact toughness and anisotropy of the SLM-formed 316L stainless steel before and after heat treatment at different temperatures. The study provides technical ideas for regulating the microstructure and properties of 316 L stainless steel parts prepared by SLM.

We use 316 L stainless steel spherical powder. The selected process parameters are as follows: laser power of 280 W, scanning speed of 1150 mm/s, layer thickness of 30 μm, and scanning spacing of 0.1 mm. Impact specimens and micro-characterization specimens are prepared in an SLM equipment. Three impact specimens, namely XZ-X, XY-Z, and XY-X, are prepared according to different printing and notch orientations (Fig.2). They are heated to 1050 ℃ and 1100 ℃ in the vacuum furnace and held for 1 h after air cooling. Three groups of comparison samples with different states are obtained, namely the SLM state, heat treatment state at 1050 ℃, and heat treatment state at 1100 ℃. Finally, the prepared impact samples are subjected to the Charpy impact test at room temperature and characterized by X-ray diffractometer (XRD), scanning electron microscope (SEM), electron back-scattered diffraction (EBSD),and transmission electron microscope (TEM).

The impact toughness of 316 L with different states and orientations shows an obvious trend. For different states, the impact toughness of SLM samples is the highest, followed by that of the heat treated sample at 1100 ℃, and that of the heat treated sample at 1050 ℃ is the lowest. For different orientations, the impact toughness of the XY-Z sample in the SLM state is the best, but it has the worst impact toughness after heat treatment (Fig.4). The phase composition characterized using XRD is found to be the single austenite phase (Fig.3). Using SEM and TEM to characterize the microstructure, it is found that 316L in the SLM state is composed of many fine crystals and cellular subgrains. After the heat treatment, recrystallization occurs, fine crystals transform into coarse grains, and subgrain boundaries gradually disappear. This phenomenon becomes more evident as the heat treatment temperature increases (Fig.7). Simultaneously, oxides rich in Si and Mn are observed in the SLM state (Fig.8). The size of the oxides increases significantly after heat treatment but does not increase after the heat treatment temperature reaches 1100 °C (Fig.9). EBSD characterization reveals that the grain size and proportion of large-angle grain boundaries increase after heat treatment, which becomes more evident as the heat treatment temperature increases (Fig.6 and Fig.10). Therefore, the effects of heat treatment on impact toughness and anisotropy are analyzed. Impact toughness is affected by oxide content, grain size, and large angle grain boundaries. Among these effects, the size of the oxide inclusions is dominant. The coarsening of the oxide after heat treatment significantly worsens the impact toughness. Although the increase in grain size and large-angle grain boundaries can improve the impact toughness, the effect on impact toughness is less than that of the oxide coarsening. It is believed that anisotropy is affected by the multi-layer structure and grain texture. The anisotropy of the SLM state is dominated by the multilayer structure, and the notch cracks of the XZ-X and XY-X samples expand between layers and are not hindered by the print layer. The notch cracks of the XY-Z samples expand perpendicular to the layer and are hindered layer by layer, releasing more energy when the impact is exerted (Fig.11). After heat treatment, the multilayer structure is destroyed, and the grain orientation of the XY plane is <110> (Fig.12), which is not conductive for increasing the impact toughness; therefore, the impact energy released by the XY-Z sample is lower.

The SLM sample exhibits the highest impact toughness, which decreases after heat treatment owing to an increase in inclusions. As the heat treatment temperature increases, the grain size and large-angle grain boundaries also increase, along with the impact toughness. The XY-Z sample in the SLM state has many impact obstacles and exhibits good toughness. After heat treatment, the obstacles are weakened, and the texture dominates. The <110> crystal orientation is not conducive to impact performance, and the toughness of the XY-Z sample decreases.

Laser powder bed fusion (L-PBF) is a representative technology in metal/alloy additive manufacturing. It utilizes a laser as the heat source with a small beam size, enabling the production of fine and intricate parts. During L-PBF additive manufacturing of complex structure parts, the forming direction is variable which increases the complexity of heat transfer and solidification. As a result, the thermal efficiency differs among different forming directions, which can impact the performance of the printed parts. However, the current research on the influence of forming direction on the microstructure and properties of L-PBF 316L stainless steel is not sufficiently systematic. Therefore, it is further necessary to deeply investigate the influence of formation direction. We investigated the correlation between the forming direction and the microstructure and properties of L-PBF 316L stainless steel and explored the variations in mechanical properties to provide valuable insights for the development and applications of L-PBF 316L stainless steel.

First, the 316L stainless steel parts were fabricated using L-PBF from two different forming directions of 0° and 60°. The microstructure and mechanical properties of the parts were investigated through metallographic and tensile tests. The phases of the 316L stainless steel powder and as-printed samples were determined using X-ray diffraction (XRD). The grain orientation distributions, grain sizes, and grain boundary angles of these samples at different deformation strains were characterized using electron backscattered diffraction (EBSD) and scanning electron microscope (SEM).

The results show that some porosity defects occur in the L-PBF 316L stainless steel (Fig.5). During the printing process, the rapid cooling rate leads to the retention of the α-Fe ferrite phase (Fig.6). Tensile testing shows that the samples printed from the forming direction of 60° exhibit higher tensile strength than those from the forming direction of 0°, while the samples printed from the forming direction of 0° demonstrate better elongation than those from the forming direction of 60° (Fig.7 and Table 3). In-situ tensile testing results indicate that there exist significant differences in grain boundary angles, phase contents, surface morphologies, grain orientations, and grain sizes among the L-PBF 316L stainless steel samples printed from different forming directions during the tensile deformation process. In both 0° and 60°, the samples exhibit predominantly high-angle grain boundaries before tensile deformation. However, as the deformation strains increase, the proportion (volume fraction) of low-angle grain boundaries gradually increases and finally surpasses that of high-angle grain boundaries (Fig.9). In terms of phase composition, the γ-Fe face-centered cubic (FCC) phase account for over 98% in the sample before deformation, but its proportion (volume fraction) decreases while the α-Fe body-centered cubic (BCC) phase increases with increasing deformation strains (Fig.10 and Table 5). In terms of surface morphology, the samples underwent dislocation slip and twinning during the tensile process. The sample built from the forming direction of 0° exhibits much more slip bands as well as a large amount of deformation twinning compared with the part printed from the forming direction of 60°, which improves the tensile properties of the parts (Fig.11). In terms of grain orientation, the samples manufactured using L-PBF exhibit anisotropy. For the sample built from the forming direction of 0°, the initial <101>∥Z1 grain orientation gradually transforms to <001>∥X1 and <111>∥X1 during the tensile process. In contrast, for the sample built from the forming direction of 60°, the initial <111>∥Z1 grain orientation gradually transforms to <111>∥X1 during tensile deformation. This difference in grain orientation is related to the formation of deformation twinning within the grains during tensile deformation, which induces grain orientation rotation (Fig.12). In terms of grain size, the L-PBF 316L stainless steel undergoes grain refinement with increasing deformation strain. The coarse columnar grains in the as-printed state are progressively fractured under external forces, leading to a reduction in grain size with increasing strain. The sample built from the forming direction of 0° exhibits a higher degree of grain refinement and smaller grain size than the sample printed from the forming direction of 60° (Fig.13).

The influence of two forming directions of 0° and 60° on the microstructure and mechanical properties of L-PBF 316L stainless steel was studied. The evolution of microstructure and grain orientation during tensile deformation of L-PBF 316L stainless steel was studied using in-situ tensile testing. There are some defects in the L-PBF 316L stainless steel, and a fish-scale-like melt pool occurred in the part printed from the forming direction of 60°. The sample built from the forming direction of 60° exhibits a high tensile strength, while the part printed from the forming direction of 0° shows good elongation and plasticity. During in-situ tensile deformation, the proportion of low-angle grain boundaries and the α-Fe-BCC phase content increase, the grain size decreases, and the slip bands appear within the grains. Compared with the sample built from the forming direction of 60°, these changes are much more significant in the part printed from the forming direction of 0°. In the forming direction of 0°, the initial <101>∥Z1 grain orientation gradually transforms to <001>∥X1 and <111>∥X1 during tensile deformation, while in the forming direction of 60°, the initial <111>∥Z1 grain orientation gradually transforms to <111>∥X1.

To achieve carbon neutrality, lithium batteries, as a new generation of green energy products, are poised to enter the terawatt-hour (TW·h, equivalent to 1000 GW·h) era. Currently, mainstream manufacturers of state-of-the-art energy batteries have reached a production capacity of 200 pieces per minute (PPM), and there are plans to increase the production capacity of cylindrical batteries to 300 PPM. Consequently, extremely high-speed production lines present substantial challenges to the welding process. Therefore, the development of highly efficient and reliable battery-welding technologies and processes has become an urgent concern for the automobile manufacturing industry. Laser welding, with its small laser spot, high energy density, efficient welding, precise energy control, automation capabilities, and safety features, has been widely used in the field of new energy battery welding, including vehicle manufacturing. Recent research has primarily focused on enhancing the welding quality of aluminum alloys by optimizing laser welding process parameters and beam shaping. However, as the demand for higher welding efficiency in power battery welding increases, the scope for ensuring both welding quality and speed becomes constrained, making it increasingly challenging to identify suitable process parameters. Considering that lasers, as a new type of welding light source, exhibit characteristics distinct from those of the arc light sources generated by arc welding machines, research has primarily focused on laser power. The impact of laser light source characteristics on welding quality and efficiency, particularly the influence of laser beam quality, has received limited attention. To meet the demands of high-speed production lines for new energy power batteries, the effect of the beam quality from fiber lasers in aluminum alloy laser welding is systematically analyzed in this study based on the energy distribution characteristics within the laser welding process. The quantitative relationship between beam quality and welding stability, as well as welding efficiency, is also explored.

An industrial-grade 3-kW continuous fiber laser is used in the experiment. The laser employs a circular swing path for welding with a swing amplitude of 0.6 mm and spacing of 0.25 mm. The upper-layer material consists of a 1.5-mm-thick 3003 aluminum alloy plate, while the lower-layer material is a 3-mm-thick 3003 aluminum alloy plate.

With a decrease in the laser beam quality factor (M²), the welding speed increases, corresponding to the same weld penetration (2.7 mm), and the complex process capability index (CPK) value of weld penetration also increases. When the M² is reduced from 11.6 to 1.25, the welding speed increases by 5.5 times, and the CPK value of the weld penetration increases by 2.3 times, corresponding to the same weld penetration (2.7 mm) (Fig. 4). Analyzing the energy distribution of postlaser welding oscillations on the YOZ surface reveals that, as the M² improves, there is a consistent downward trend in the maximum value of the laser energy density. In addition, the area between the two maximum values decreases. For instance, the maximum energy density at M²=1.18 is 1.4 times higher than that at M²=11.6 (Fig. 5). To gauge the influence of the laser energy density on the welding efficiency, a factor derived from the product of the maximum laser energy density and the difference in the maximum value of the energy density is introduced. Calculations demonstrate that, at M²=1.18, this influence factor for laser energy density welding efficiency is 5.2 times higher compared with that at M²=11.6 (Fig. 6). To facilitate the assessment of the area disparities between the maximum peak and intermediate minimum at both ends of the weld, the region between these points on the energy distribution map is defined as the laser energy occupied space line density. Remarkably, when M² is 1.18, the laser energy occupied space line density is 10.7 times greater than that when M² is 11.6 (Fig. 6).

The influence of the beam quality from fiber lasers on aluminum alloy welding is systematically analyzed by considering the energy distribution characteristics during the welding process. A quantitative relationship between beam quality and welding stability, as well as welding efficiency, is established. The improvement mechanism for the laser beam quality effect is as follows. When the M² decreases while retaining the same penetration depth, the maximum value of the energy density increases, and the energy density between the edge of the weld bead and the middle interval also increases. The higher the energy absorption efficiency of the workpiece material, the higher the welding efficiency of the laser. The product of the maximum value of the laser energy density and the difference in the maximum value of the energy density is used as a factor influencing the laser energy density welding efficiency. The theoretical calculations show that, when M² is 1.18, the influence factor of the laser energy density welding efficiency is 5.2 times higher than that when M² is 11.6, which is basically consistent with the experimental results.

Laser utilization as a heat source to connect nickel-based superalloys has been applied in aviation, aerospace, weapons manufacturing and other fields. Solidification behavior of the laser welding molten pool of nickel-based alloys, including the nucleation, growth, collision and movement of grains, and the diffusion, enrichment and segregation of alloying elements, directly affect nickel-based alloy laser weld performance. Therefore, in-depth research of the solidification behavior of the nickel-based alloy laser welding molten pool is of paramount significance, particularly for laser welding process optimization, welding defect formation control, and the improvement of laser weld mechanical properties.

The macroscopic heat and mass transfer coupling model alongside microstructure evolution is utilized to quantitatively simulate the macroscopic heat and mass transfer, and the microstructure evolution of a laser welded IN718 alloy. The fluid volume (VOF) method is used to simulate the molten pool morphology and temperature field distribution of the macroscopic heat transfer process, and the temperature field distribution replaces the solidification parameter variables relating to the phase field control equation and is brought into the phase field model for microstructure evolution process simulation.

The simulated and experimental microstructures are shown in Fig.6, where it can be observed that all tissues are columnar crystal structures with a consistent morphology. As shown in Fig.7(a), the solidification velocity R and temperature gradient G at different positions were extracted in the macroscopic simulation results. The solidification velocity R gradually decreased from top to bottom with the molten pool boundary, and the temperature gradient G gradually increased. The cooling rate G·R showed a decreasing trend, the dendrite spacing gradually increased, and the results are shown in Fig.7(b). Compared with the Hunt and Kurz et al. numerical models, the phase field model calculated results are more accurate and close to the experimental results, showing consistent regularity, as shown in Fig.8. The distribution of Nb elements perpendicular to the growth direction of columnar crystals in the simulation results is shown in Fig.9 (b), this shows obvious periodic changes, and the change period is closely related to the phase morphology, because the IN718 solute partition coefficient is less than 1. Solute elements tend to be segregated and enriched at the dendrite gap position. In the simulation results, the mass fraction of Nb elements inside the solid-phase dendrite is significantly reduced, and the lowest mass fraction occurs at the columnar center owing to segregation caused by solute redistribution during solidification. The mass fraction of Nb element increased significantly at the liquid phase position of columnar crystal gap, and the mass fraction of Nb element in the position was relatively higher than that closer to the bottom of the columnar crystal. From Fig.10, the Laves phase enriched by the Nb element precipitated from the IN718 molten pool after solidification is distributed in the γ phase matrix in the shape of droplets. The Laves phase morphology and distribution are approximately consistent with the simulated distribution of Nb elements. The SEM results are shown in Fig.11(b), the EDS spot scan analysis is performed on the illustrated position, with the test results shown in Figs.11(a) and (c). The mass fraction of Nb element in the matrix γ is approximately 4.7%. The mass fraction of Nb element in the Laves phase is approximately 9.8%, and the Nb element mass fraction in Laves phase is significantly higher than that in matrix γ. This proves the accuracy and reliability of the simulation results.

The results demonstrate that the simulated microstructure grows in a columnar crystal structure. The solidification rate R of the molten pool gradually decreases from top to bottom, the temperature gradient G gradually increases, the cooling rate G*R decreases continuously, and the primary dendrite arm spacing increases with the decrease in cooling rate, from 4.52 to 7.12 μm, which is consistent with the experimental results. The mass fraction of Nb elements in the columnar crystal spacing increased significantly, and this mass fraction was relatively higher near the bottom of the columnar crystals in the liquid phase. The Nb elements are finally distributed in the shape of droplets and approximately consistent with the Laves phase in morphology and distribution, which is also consistent with the experimental results. The microstructure transformation process and elemental segregation behavior of IN718 in laser welding are examined, and the solidification theory of the laser welding molten pool of nickel-based alloys is enriched. Finally, this research provides a foundation for a numerical solution of the defect formation process for laser welding IN718 cracks and pores.

Laser technology is extensively used in various fields, including additive manufacturing, removal processing, and laser weaponry. This technology has the potential to revolutionize battlefield dynamics through defensive and offensive applications. Research on laser irradiation under high-speed airflow provides a theoretical basis for efficient damage strategies and the protection of aircraft, this is crucial for deploying military laser systems. However, conducting real-scale model tests for large-scale engineering structures is challenging due to equipment limitations and testing environments. Additionally, wind tunnel tests with real-scale models are prohibitively expensive and time-consuming, preventing extensive testing. Consequently, scaled-model tests are often relied upon for regularity studies. Therefore, establishing a similarity relationship in the thermomechanical responses between real and scaled models under laser irradiation and high-speed airflow is a practical approach. Significant efforts have been made to understand the similarity theory of laser-induced thermomechanical behavior under static air conditions. Nonetheless, due to the complex fluid-thermal-structural interactions in high-speed airflow, the similarity criteria for thermomechanical responses in an airflow environment significantly differ from those in static air. In this study, we propose new similarity criteria and scaling laws suitable for the thermomechanical responses of a metal plate subjected to high-speed airflow and laser irradiation.

To clarify the similarity relation of thermomechanical responses for metal plates under coupling conditions, the effects of the tangential airflow were equivalently converted to the structural force and thermal load boundary conditions using the approximate equivalence method, and the dimensionless governing equations of the coupling problem were established. Thus, combined with the analysis of dominant factors, the similarity criteria and scaling laws suitable for the thermomechanical responses of the metal plate under the combined action of a high-speed airflow and laser were determined. According to the similarity criteria, there is a contradiction in the similarity relationship between the thermal boundary condition and force boundary condition under the fluid-thermal-structural coupling effects, which cannot be satisfied simultaneously. Considering that the thermal stress induced by the temperature gradient is much greater than the mechanical stress due to the aerodynamic force under the combined action of high-speed airflow and laser irradiation, this study focused on the similarity of aerodynamic heat transfer, ignored that of the aerodynamic force, and established the corresponding scaling law. Then, a fluid-thermal-structural coupling numerical example of a metal plate irradiated by a high-power laser under tangential flow was conducted to verify the scaling law under different scale ratios and Mach numbers.

The similarity criteria and scaling laws for the fluid-thermal-structural coupling analysis of the metal plate subjected to laser irradiation and high-speed airflow are presented in Tables 1 and 2, respectively. A numerical example of the fluid-thermal-structural coupling of a metal plate irradiated by a high-power laser under a tangential flow is conducted to verify the scaling law. The results show that under different scale ratios and Mach numbers, the predicted response errors between the scaled and original models are within 1%, which proves the reliability and accuracy of the scaling law. Simultaneously, with the increase in scale ratios or Mach numbers, the aerodynamic heat transfer effect is enhanced, making the thermal-mechanical response difference between the scaled model and real model more obvious when the aerodynamic similarity criteria are not considered.

In this study, similarity criteria and scaling laws suitable for the thermomechanical responses of a metal plate under the combined action of a high-speed airflow and laser are determined. Several numerical examples are conducted and compared to verify the proposed similarity criteria and scaling laws. The main conclusions are as follows: (1) Using the approximate equivalence method and analysis of dominant factors, the effects of the tangential airflow are equivalently converted to structural force and thermal load boundary conditions, and the similarity criteria and scaling laws are determined. Considering that the thermal stress induced by the temperature gradient is significantly greater than the mechanical stress caused by the aerodynamic force, this study focusses on the similarity of the aerodynamic heat transfer and ignores the similarity of the aerodynamic force. (2) A fluid-thermal-structural coupling numerical example of a metal plate irradiated by a high-power laser under tangential flow was conducted to verify the scaling laws under different scale ratios and Mach numbers. The results show that the predicted response errors between the scaled and original models are within 1%, which proves the reliability and accuracy of the scaling laws. (3) However, the scope of application of the proposed similarity criteria should be emphasized in the following aspects: the similarity criteria are applicable for calorically perfect gases. For hypersonic flows, complex chemical reactions occur at high temperatures, and the similarity criteria are no longer applicable. The similarity criteria are applicable for the plate flow condition. However, for the non-plate flow, such as the flow around a blunt-nosed bodies, the similarity criteria are no longer applicable. The similarity criteria are applicable to the thermal-mechanical responses of the metal structure before melting. When melting is involved, similarity criteria are no longer applicable.

With the development of the new energy vehicle industry, laser welding has become increasingly popular in the manufacturing of power batteries because of its high welding speed, small heat-affected zones, and high degree of automation. However, as the laser involves high-energy beam, its interaction with materials is often intense. This can easily lead to defects, such as spatters and explosion points, thus compromising the quality of battery welding. In the field of power batteries, the adjustable ring-mode (ARM) laser has emerged as a high-speed low-spatter laser welding tool, gaining attention from both academia and industry. However, the spatter suppression mechanism of the ARM laser during high-speed welding remains unclear. This limitation hinders theoretical guidance and process optimization for industrial applications. Thus, in this study, the complete welding of an aluminum alloy roof is considered and how the core ring power ratio affects the penetration and width is analyzed. Moreover, how the ARM laser effectively curbs the metal spatter is elucidated by examining the dynamic behavior of the keyhole in the molten pool. Optical coherence tomography (OCT) measurement technology is used to monitor keyhole depth fluctuations in real time, providing a quantitative assessment of welding stability and identifying the optimal process window.

A synchronous-sensing monitoring platform (Fig. 1) is established by integrating high-speed visual shooting with penetration detection. For the visual sensing component, a high-speed camera is utilized to capture sharp keyhole images of the molten pool. For penetration detection, the platform is merged with an OCT-based monitoring module to acquire real-time keyhole depth information during the welding process. Initially, the process window of the ARM laser welding is determined by conducting an orthogonal experiment, as shown in Fig. 3. The keyhole images under different parameters are obtained, and the changes in the keyhole depth are recorded. Comparisons of the keyhole opening and depth reveal the mechanism behind the spatter suppression during ARM laser welding. To identify the best low-spatter process window, keyhole volatility is introduced as a variable. The variance in the keyhole depth, measured by applying OCT in real time, is calculated. This variance is used to assess the depth fluctuations of the keyhole and, consequently, the stability of the welding process. The relationship between the welding process stability and the power ratios of inner ring laser to outer ring laser is then established by using a contour map, resulting in the identification of optimal process window parameters (Fig. 10).

The spatter formation mechanism in the ARM laser high-speed (150 mm/s) welding is analyzed. The spatter formation process and suppression methods are elucidated, demonstrating that the ARM laser can indeed diminish the spatter occurrence rate by enlarging the keyhole opening. The effect of the power ratio of inner ring laser to outer ring laser on the keyhole stability is verified. First, a traditional orthogonal experiment is conducted to determine the process window for melting width when the inner ring laser power ranges from 600 W to 1300 W and the outer ring laser power ranges from 800 W to 1800 W. The process window for the penetration is determined for an inner ring laser power of 500?1150 W and outer ring laser power of 800?1800 W. Subsequently, the optical coherence scanning technology is employed to acquire the keyhole depth information. This information enables a qualitative evaluation of the welding process stability, facilitating the process optimization of the ARM laser welding. The findings suggest that a higher outer ring laser power is better for achieving a suitable penetration. A higher outer laser ring power stabilizes molten pool fluctuations and enlarges the keyhole opening.

This study presents a process optimization scheme combined with real-time monitoring of the laser welding depth. The theory that spatter is mainly caused by keyhole collapse is verified. The laser welding process is further optimized based on the standard deviation of keyhole depth fluctuations. The final process window that satisfies both the traditional process window and keyhole fluctuation stability analysis window is identified: the core laser power ranges from 800 W to 1000 W, ring laser power is between 1200 W and 1600 W, and welding speed is set at 150 mm/s. The optimal power ratio of inner ring laser to outer ring laser for welding aluminum alloys typically lies between 1∶2 and 1∶3. Within this range, the keyhole achieves maximum stability and the defect occurrence rate is the smallest.

The use of low-density lightweight materials, such as aluminum alloys, instead of traditional steel, titanium, and other materials to form a dissimilar material composite structure is an important way to achieve a light weight. Because of the different physical and chemical properties of aluminum alloys and steel, it is difficult to join aluminum and steel by laser welding-brazing. Brittle intermetallic compounds (IMCs) in the interface layer are easily produced owing to the small solid solubility between iron and aluminum. In this research, a rotating laser is applied to improve the temperature distribution and optimize the interface reaction. Based on the analysis of the morphology, type, and thickness of the interface layer of the aluminum/steel laser welding-brazing joint under different rotating parameters, the mechanical properties of the aluminum/steel welded joint are studied by a tensile test, and the fracture morphology and fracture mode of the joint are also investigated.

The test materials are a 304 stainless-steel plate with a size of 100.0 mm×80.0 mm×0.9 mm and a 6061-T6 aluminum alloy sheet with a size of 100.0 mm×80.0 mm×1.2 mm. AlSi12 is used as filler wire. A fiber laser is used as the heat source. High-purity argon (volume fraction of 99.99%) with a gas flow rate of 25 L/min is used as the protective gas. After welding, the cross-sectional morphology of the weld is observed by using a metallographic microscope. A scanning electron microscope (SEM) is used to analyze the morphology and thickness of the IMC layer. The chemical composition of the interface layer is detected using an energy-dispersive X-ray spectroscope (EDS) system integrated with the SEM. The mechanical performance of the joint is represented by the line load. The fracture morphology is observed using the secondary electron detector of the SEM.

After the addition of the rotating laser, the thickness of IMCs composed of θ-(Fe, Ni)(Al, Si)₃ and τ₅-(Fe, Ni)_1.8Al_7.2Si is significantly reduced, and the uniformity of the IMC layers is improved. The line load of the joint without a rotating laser is 215.9 N/mm. The joint with a rotation diameter of 2 mm has the largest line load of 289.1 N/mm, which is 33.9% higher than that without rotation. Compared with nonrotating-laser joints, the joint line load increases because of the thinning of the intermetallic layer and the reduction of the complexity of the IMC. At a rotation diameter of 2 mm and frequency of 30 Hz, a fracture occurs at the weld. Under these parameters, the IMC thickness of the joint is uniform and only composed of the τ₅-(Fe, Ni)_1.8Al_7.2Si phase. Compared with the θ-(Fe, Ni)(Al, Si)₃/τ₅-(Fe, Ni)_1.8Al_7.2Si interface, a single τ₅-(Fe, Ni)_1.8Al_7.2Si/steel interface achieves relatively low interface crystal plane mismatch and better bonding performance, thereby improving the tensile performance of the joint. In the EDS results in Table 4, α-Al and Al-Si eutectic on the fracture can be observed. Additionally, many dents are formed on the fracture, and the fracture mode is ductile fracture.

Compared with the rotating frequency, the rotating diameter has a greater influence on the wetting width of the joint. To obtain a well-formed rotating laser welding-brazing aluminum/steel joint, the welding process parameters should be optimized with a laser rotating diameter of 2 mm. When the laser is not rotating, two layers of IMCs with a thickness of approximately 8.45 μm are formed at the interface. After the rotating laser is applied, the thickness of the intermediate layer is reduced, and the variety is decreased. The rotating laser reduces the welding peak temperature and inhibits the formation of brittle IMCs. At a laser rotation diameter of 2 mm and frequency of 30 Hz, the linear load reaches a maximum value of 289.1 N/mm, which is approximately 33.9% higher than that without the rotating laser. The fracture position of the joint changes from the interface layer without the rotating laser to the weld.

Laser welding has the characteristics of a high energy density, low heat input, and high welding efficiency; however, conventional laser welding has a small focused spot and high requirements for the welding assembly gap. To solve this problem, scholars have developed laser wire filling welding technology. Based on this, some scholars have developed laser hot wire welding technology, which can effectively improve the absorptivity of the welding wire by preheating the welding wire in advance, reduce the requirements for laser power, and improve the welding speed; however, there are still problems such as high requirements for the alignment of the laser focus and the tip of the welding wire, and an uneven weld height. In this study, the process characteristics of Q235 steel by scanning laser hot wire welding are systematically studied, and the mechanism of the influence of the scanning laser on the solidification process of weld metal is clarified, which provides technical guidance for expanding the industrial application of laser welding.

The base material used in this study is the Q235 steel plate. The size is 50 mm×120 mm×2 mm, and the structure is massive ferrite at normal temperature. The flat surfacing welding method is used in the research on the weld surface and section forming. The docking method is adopted in the study of the microstructure and properties of welded joints. According to the previous research and accumulation of this research group, the fixed wire feeding method is front wire feeding, the tilt angle of the welding torch is 45°, and the laser focus is located on the surface of the plate, that is, the defocus quantity is 0 mm. In the welding process, the shielding gas is argon with purity (volume fraction) greater than 99.99%. The gas pipe angle is 60° and the gas flow rate is approximately 20 L/min. In the butt welding experiment, the fixed laser power is 1.8 kW, the welding speed is 1.0 m/min, the preheating current of laser cold wire welding is 0 A, the preheating current of laser hot wire welding is 100 A, the scanning amplitude ranges from 0.6 mm to 1.0 mm, and the scanning frequency ranges from 100 Hz to 200 Hz.

Under different scanning parameters, the distribution of the laser energy is different, which affects the temperature field distribution of the weld pool, and then affects the macro forming, microstructure, and properties of the weld. Compared with that in non-scanning laser hot wire welding, the weld forming in scanning laser hot wire welding is smoother and straighter, and the splash is less (Fig. 5). The weld structure in non-scanning laser hot wire welding is dominated by thick side lath ferrite. Because the scanning laser enhances the flow of the molten pool through the stirring effect and breaks the coarse columnar crystals, the weld structure in scanning laser hot wire welding is dominated by fine crystalline ferrite and acicular ferrite with finer grains (Fig. 7). The tensile strength (578.8 MPa) of the scanning laser hot wire welded joint is basically the same as that (574.7 MPa) of the non- scanning laser hot wire welded joint, but the elongation is increased from 8.4% to 13.1% (Table 3). The dimple size of the tensile fracture surface of the scanning laser hot wire welded joint is more uniform, and the dimple size difference between the laser hot wire welded joint and the laser cold wire welded joint is larger; moreover, there is obvious inclusion precipitation at the bottom of the dimple, indicating that the scanning laser improves the homogeneity of the weld structure (Fig. 10). Simultaneously, the scanning laser improves the gap tolerance during butt welding. In the butt welding experiment of the Q235 steel plate with a thickness of 2 mm, the scanning laser hot wire welding ensures good weld formation without defects when the gap is 1.3 mm (Fig. 12).

In the experiment of scanning laser hot wire welding, by optimizing the process parameters, when the scanning diameter is 0.4?1.0 mm and the scanning frequency is 50?200 Hz, the welds obtained are well formed, smooth, no defects and nearly no splash, which proves that the scanning laser has a good improvement effect on the weld formation. Simultaneously, the scanning laser improves the gap tolerance of laser hot wire welding, which is conducive to achieve stable welding when the gap is uneven and obtain a weld with good fusion with the base metal side wall and no surface collapse. At the microstructure level, the stirring effect of the scanning laser on the weld pool can promote the flow of the weld pool and refine the grain. In terms of mechanical properties, compared with that in non- scanning laser hot wire welding, when the tensile strength is basically unchanged, the fracture elongation increases to 13.1% in scanning laser hot wire welding, indicating that the addition of the scanning laser can effectively improve the toughness of the weld, which is also proved by the deeper dimples in the electron microscope image of the fracture. The hardness of the fusion zone in the laser hot wire welding is the highest, followed by that of the heat affected zone, whereas the hardness of the base metal is the lowest. The hardness of the fusion zone in the scanning laser hot wire welding is lower than that in the non- scanning laser hot wire welding, mainly because the fusion zone in the non- scanning laser hot wire welding is easy to produce segregation, and the generated inclusions increase the microhardness.

Compared with laser welding and arc welding, laser-arc hybrid welding not only inherits the advantages of laser welding and arc welding but also makes up for respective shortcomings. Thus, it is an advanced welding process method with great application prospects. With the continuous development of laser technologies, laser power has exceeded 10 kW or even higher. Therefore, in order to make the development of lasers well meet the need of actual industrial production, the basic theoretical research on high-power laser-arc hybrid welding has been a hot spot in the academic community in recent years. Researchers have carried out a lot of research on the interaction mechanism between laser and arc. However, the laser power involved was mostly below 5 kW. There are few reports on the mechanism regarding the effect of a high-power (higher than 5 kW) laser on the droplet transfer in laser-arc hybrid welding. Therefore, in this study, a high-power (7.5 kW) laser is introduced into the different modes of arc [standard metal active-gas(MAG), cold metal transfer (CMT), and pulsed arc] welding process, and its effects on droplet transition, weld forming and welding efficiency are compared and studied by using high-speed camera, optical microscope, etc.

In this study, a high-power laser-arc hybrid welding platform was built, which mainly consisted of a continuous fiber laser, a welding system, a manipulator arm, and a high-speed camera system. The high-power laser-arc hybrid welding experiments were carried out on 10 mm thick Q345 steel, and the laser used in the test was a fiber laser (maximum output power of 12 kW), with an output laser wavelength of (1080±10) nm and a focused spot diameter of 0.2 mm. Before the welding test, an angle grinder was first used to grind the surface to be welded, and then the ground surface was cleaned with alcohol. The arc-guided laser-arc welding was chosen for obtaining a stable droplet transition process. In order to further understand the influence of a high-power laser on droplet transition in different modes of arc welding, the laser was coupled with three different arc modes (standard MAG, CMT and pulsed arc). The welding shielding gas used in the welding process was the Ar and CO₂ mixture with a flow rate of 20 L/min. During the welding process, a high-speed camera was used to track and monitor the droplet transition behavior with a frame rate of 10000 frame/s. In order to obtain a clear droplet transition image, an infrared filter was added to the camera lens before the experiment began. Image pro plus software was used to process the pictures taken by the high-speed camera, and the droplet transition mode and the number of droplet transitions within 500 ms under each parameter were counted, so as to calculate the droplet transition frequency within 1 s. After welding, the forward and cross-sectional morphologies of the weld were observed by optical microscope.

The high-power laser has a significant effect on the droplet transition mode of arc welding in different arc modes. During standard MAG welding, the high-power laser attracts and compresses the arc, resulting in a significant reduction in arc length. Meanwhile, metal vapor and plasma ejected from the keyhole reduce the droplet transition frequency (Figs. 6 and 7). In the case of CMT welding, the high-power laser extends the single short-circuit transition period, and the resulting molten pool oscillation reduces the stability of the short-circuit transition (Fig. 8). Regarding the pulsed arc welding process, the high-power laser increases the melting rate of the welding wire. In the meantime, the droplet transition mode changes from the droplet transition to the jet transition, and the droplet transition frequency is significantly increased. The air flow at the key hole hinders the droplet transition, so that the droplet transits to the side of the molten pool (Figs. 11 and 12). Compared with that during arc welding, the weld melting width increases during laser-standard MAG and laser-pulsed arc hybrid welding, while no obvious change in weld width is observed in the case of laser-CMT hybrid welding. The residual height of welds in laser-standard MAG and laser-CMT hybrid welding decreases significantly, while the residual height of welds in laser-pulsed arc hybrid welding increases slightly. This is attributed to different degrees of influence of the laser on the droplet diameter and transition frequency in three different modes of arc welding. Furthermore, the melting energy increment value (ψ) of laser-arc interaction varies under different hybrid welding conditions, among which laser-pulse arc welding has the highest ψ value (36%), followed by laser- standard MAG welding (19%), while laser-CMT welding has the smallest ψ value (-12%).

In this study, the effects of laser (7.5 kW power) on droplet transition and weld formation in different modes of arc welding were investigated. The results reveal that the addition of laser has a significant influence on the droplet transition in standard MAG, CMT and pulsed arc welding processes. During standard MAG welding, the high-power laser attracts and compresses the arc, resulting in a significant reduction in arc length, and the metal vapor and plasma ejected from the keyhole reduce the droplet transition frequency. In the CMT welding process, the laser extends the single short-circuit transition cycle, and the melt pool oscillation caused by the high-power laser reduces the stability of the short-circuit transition. Regarding the pulsed arc welding process, the addition of a high-power laser increases the melting rate of welding wires. The droplet transition mode changes from the droplet transition to the jet transition, and the droplet transition frequency increases. Meanwhile, the air flow at the key hole hinders the droplet transition, so that the droplet transits to the side of the molten pool. Compared with arc welding, the weld melting width increases during laser-standard MAG and laser-pulsed arc welding, while no obvious changes in weld width are observed in the case of laser-CMT hybrid welding. The residual height of the welds in laser-standard MAG and laser-CMT hybrid welding decreases significantly, while the residual height of welds in laser-pulsed arc hybrid welding increases slightly. The melting energy increment values of the interaction between laser and arc under three arc modes are: laser-pulsed arc hybrid welding (36%), laser-standard MAG hybrid welding (19%), and laser-CMT hybrid welding (-12%).

Femtosecond laser direct writing is a new nanofabrication technology. However, achieving high-quality pattern transfer directly on metal surfaces is difficult owing to the photothermal effect between the femtosecond laser and metal, which can result in laser ablation. To address this, we develop a new metal pattern transfer technology that combines femtosecond laser direct writing technology and the lift-off process. This technology enables precise metal pattern transfer at a sub-micron level (0.89 μm). We explore the effects of femtosecond laser direct writing parameters and development methods and parameters on the performance of metal transfer. We observe that the accuracy of metal transfer increases with higher femtosecond laser exposure doses. Meanwhile, the dry development method outperforms the wet development method. Increasing the development time of the sacrificial layer results in a gradual increase in the transfer line width, accompanied by improved roughness of the line. By incorporating a specific undercut angle in the metal transfer process, we successfully address the edge warpage and high roughness of the metal lines. In this study, a chromium-based transfer grating with 2 inch (1 inch=2.54 cm) diameter is fabricated using the new approach. Additionally, the feasibility of using the prepared chromium pattern as a mask plate is verified using projection exposure on SU-8 photoresist, thus demonstrating the potential of the strategy to locally replace the e-beam lithography for mask plate processing. In addition, Au and Pt patterns are successfully transferred using this approach, demonstrating its universality and wide range of potential applications.

The lift-off process based on femtosecond laser processing comprises sample preparation, femtosecond laser exposure, development, sputtering of metal, and stripping. 1) Sample preparation: The quartz substrate is cleaned using an ultrasonic machine and vacuum plasma machine to remove impurities. Thereafter, a sacrificial layer is spin-coated onto the quartz substrate and baked on a hot plate. Next, photoresist is added dropwise onto the sacrificial layer. 2) Femtosecond laser direct writing: A femtosecond laser is used to expose the prepared samples. 3) Development: The exposed samples are developed by immersing them successively in propylene glycol methyl ether acetate (PGMEA) and isopropyl alcohol (IPA) to wash away the unexposed photoresist. The sacrificial layer is then removed by immersion in NMD-3 solution. Dry development is performed on a microwave plasma debonder. 4) Sputtering of metals: Chromium and gold plating is performed using magnetron sputtering equipment, and platinum plating is performed using a magnetron ion sputtering apparatus. 5) Stripping: The residual photoresist and sacrificial layer are removed using the degumming solution, N-methyl pyrrolidone (NMP), through immersion and ultrasonication, followed by rinsing with deionized water. The samples are left to air dry to complete the transfer of the pattern of the metal.

This study explores the effects of sacrificial layer development methods and parameters on chrome transfer. The results in Figs. 3(a)?(c) indicate that the wet developing method has many defects and is ineffective, whereas the dry development method achieves high-quality chrome transfer within an optimal development time of 20 min [Figs. 3(d)?(f)]. Next, we investigate the effect of femtosecond laser direct writing process parameters on chrome transfer properties. The results show that increasing the exposure dose decreases the width of the chrome lines (Fig. 4). At the direct writing speed of 20 mm/s and direct writing power of 47.5 mW, the line width is as small as 890 nm, albeit with warpage at the line edges. To solve these problems, we design and introduce the undercut angle in the photoresist to protect the chrome lines from the negative effects of the lift-off process. Figure 5(c) shows that the transferred chrome line warps severely on both sides without the undercut angle. When the undercut angle is 10°, the edges of the chrome lines are smoother and the chrome lines tightly adhere to the substrate without any warpage [Fig. 5(d)]. Thereafter, the edge roughness of the chrome lines is reduced by optimizing the sacrificial layer dry-developing process (Fig. 6). Finally, we explore the large-area pattern transfer and multi-metal pattern transfer capabilities of this strategy. Combining the femtosecond laser direct-writing technique and lift-off process, we successfully achieve a 2-inch chromium-based transfer grating and verify the feasibility of using the chromium pattern as a mask plate for projected exposure on SU-8 photoresist. Meanwhile, Au and Pt patterns are successfully transferred using this strategy, demonstrating its universality.

A novel metal pattern transfer technique is developed by combining femtosecond laser direct writing technology and the lift-off process. In the lift-off process, a sacrificial layer is introduced and a dry development strategy is used to achieve metal pattern transfer with submicron (0.89 μm) precision. The effects of femtosecond laser direct writing parameters and development methods and parameters on metal transfer are explored. We observe that the accuracy of metal transfer increases with an increase in the femtosecond laser exposure dose. Increasing the development time of the sacrificial layer increases the transfer line width and significantly improves the roughness of the line. Moreover, the problems of edge warpage and high roughness of metal lines are successfully addressed by incorporating a specific undercut angle in the metal transfer process. In addition, a 2-inch chrome-based transfer grating is achieved using the proposed approach. The feasibility of using the prepared chrome pattern as a mask plate for projected exposure on SU-8 photoresist is confirmed. Finally, the successful transfer of Au and Pt patterns demonstrates the universality of this approach.

The imaging phenomena in extreme ultraviolet (EUV) lithography must be elaborated from more than one perspective. Traditionally, previous studies on waveguide methods have considered the cladding electric field distribution of the absorber as an evanescent field, which is similar to a single parallel-plate waveguide. However, these studies ignore the periodicity of the absorbers. In this study, the absorber of the EUV lithography mask is regarded as a grating waveguide. Owing to the periodicity of the absorber, two adjacent periods must affect the field distribution. Therefore, the electric field of the absorber is a linear superposition of the adjacent periodic field distributions. We propose that the electric field distribution in the absorber in the lowest-order transverse electric (TE₀) mode is a hyperbolic cosine function cosh( ·). We provide the zero-order approximation value n_eff,0 of the effective refractive index n_eff for the TE₀ mode. To further decrease the relative error of n_eff,0 according to the boundary conditions, we derive the eigenvalue equation for the grating waveguide. To obtain a good approximation, we derive an iterative formula of n_eff,m and use the iteration method to decrease relative error.

According to the line types in the grating waveguide, we assumed a fitting curve function for the electric field distribution of the TE₀ mode and provided the α value range of the grating waveguide. According to waveguide theory, the field distribution in the core should be cos( ·). Owing to the periodicity of the absorber, the electric field distribution in the cladding must not be in an exponential decay form, which will be cosh( ·). Furthermore, we assumed a zero-order approximation value for the effective refractive index. To verify the accuracy of n_eff,0 and the feasibility of cosh( ·), we used the COMSOL Multiphysics software to simulate the TE₀ mode in the grating waveguide. We selected Au and Ag as absorber materials for the simulation, and the findings indicated that there were small errors between the simulated and theoretical results. To increase the accuracy further, the eigenvalue equation of the grating waveguide was obtained according to the boundary conditions. We also derived an iterative formula for the mth-order effective refractive index (n_eff,m). As a special example, we selected Au as the cladding material to further verify the iterative formula. An iterative relation equation and iteration method were used to decrease the relative error.

The relative errors of n_eff,0 for Au and Ag are 0.75% and 0.96%, respectively (Table 1). The accuracy of n_eff,0 is extremely high, but there is still a slight error between the theoretical and simulated field distributions (Fig.3). To further increase the accuracy, we selected Au to verify the iterative formula. The relative error changes with the number of iterations (Fig.4); with an increase in m, the relative error decreases. When m=23, the relative error decreases to less than 10^-5. In this case, the field distribution also shows very good agreement with the simulated result shown in Fig.5. It can be observed that with the iterative formula, n_eff,23 can describe the TE₀ mode of the grating waveguide accurately.

We consider the absorber of the EUV lithography mask as a grating waveguide and perform a rigorous simulation to describe the TE₀ mode of the absorber. Owing to the periodicity of the grating waveguide, the electric field in the cladding of the absorber is a linear superposition of two adjacent periodic field distributions, which is the cosh( ·) function. We propose a zero-order approximation value for the effective refractive index. The accuracy of the zero-order approximation value is verified by selecting Au and Ag absorbers for the simulation using COMSOL. The relative errors in n_eff,0 for the two materials are 0.75% and 0.96%, respectively. The relative errors were already small initally, and we use an iterative method to further increase the accuracy of n_eff,0. The eigenvalue equation for the grating waveguide is derived based on the boundary conditions. Subsequently, a simple iterative formula with a high accuracy is obtained. As a specific example, the Au absorber material is selected to verify the feasibility of the iterative formula. After seven iterations, the relative error of n_eff,7 decreases to 0.11%. After 23 iterations, n_eff,23 converges to the simulation value, and the relative error decreases to less than 10^-5. The feasibility and accuracy of the zero-order approximation value and iterative formula are verified.

Resin-based composite material (CFRP) surface coatings have always faced the risk of morphological damage and substrate overheating during laser cleaning. The key for solving these problems lies in the need for sufficient process experiments to establish a reliable relationship between the cleaning parameters and characteristics. Cleaning depth (H), surface roughness (S_a), and cleaning temperature (T) are the three most important cleaning indicators. H represents cleaning efficiency and effectiveness, S_a is related to the quality of re-coating, and T reflects the trend of thermal damage. Therefore, this study uses an infrared nanosecond laser to remove paint from a CFRP surface and uses laser power (P), scanning speed (V), overlap rate (η), and repetition frequency ( f ) as variables to study and statistically analyze the H, S_a, and T of the samples. Infrared thermography and high-speed imaging techniques are used to observe the temperature response of the samples, the state of the plume, and the dynamic behavior of the paint layer to determine the cleaning mechanism of the paint layer. This study is expected to provide a basic reference for improving the efficiency of laser paint removal and the quality of respraying and reducing thermal damage to CFRP substrates.

laser power, scanning speed, repetition frequency , and overlap rate. Five levels are designed under each group of parameters to form an L25 orthogonal matrix. Then, a laser cleaning experiment is conducted to obtain 25 sets of samples ranging from No.1 to No.25. After the cleaning procedure is completed, the macroscopic and microscopic morphologies of the cleaned samples are observed. At the same time, the paint cleaning depth and sample surface roughness are measured via a laser confocal microscope. Finally, the obtained experimental data are analyzed using the analysis of the variance (ANOVA) and signal to noise ratio (S/N) methods. In addition, an infrared thermographic camera is used to record the temperature response of the experimental samples during the cleaning process, and a high-speed camera is used to capture the dynamic behavior of the samples.

A signal-to-noise ratio analysis is performed on the cleaning depth, surface roughness, and cleaning temperature using the expected large, large, and small characteristics, respectively. The analysis results (Table 4) indicate that for the cleaning depth, the influencing factors are ranked from high to low by weight, namely, lap rate, laser power, scanning speed, and repetition frequency. For surface roughness and cleaning temperature, the influencing factors are ranked from high to low by weight, namely, lap rate, scanning speed, laser power, and repetition frequency. The ANOVA results (Table 5) indicate that for cleaning depth, roughness, and cleaning temperature, the critical probability (P') values of the overlap rate, scanning speed, and laser power are all less than 0.05. Therefore, at a 95% confidence level, the overlap rate, scanning speed, and laser power have statistically significant effects on cleaning depth, roughness, and cleaning temperature. In contrast, the contribution rate of repetition frequency is relatively low, with a P' value greater than 0.05, making it a less important process parameter. The detection results (Fig. 8) by the infrared thermal imager indicate that the laser cleaning process causes two high-temperature areas. The first is where the laser acts on the substrate. The minimum cleaning temperature in this area is 244 ℃, and the maximum cleaning temperature is 590.4 ℃. The other high-temperature region is the high-temperature plume region above the sample. The high-speed camera monitoring results (Fig. 11) indicate that the paint layer undergoes drastic changes due to the action of the laser, the most obvious being the generation of bright plasma and the formation of a plume perpendicular to the sample. A large number of turbid particles are observed inside the plume.

This study focuses on the influence of process parameters on the laser cleaning of paint layer on the CFRP aircraft skin. For the cleaning depth, surface roughness, and cleaning temperature, the overlap rate is the most significant influencing parameter, with contribution rates of 50.51%, 59.07%, and 69.09%, respectively. A lower overlap rate is not conducive to the uniform removal of paint, and an increase in the overlap rate will significantly increase the temperature of the substrate. Laser power and scanning speed also have a significant influence on cleaning depth, surface roughness, and cleaning temperature, whereas repetition frequency has no significant effect. The removal of paint is mainly based on the thermal erosion mechanism. During the cleaning process, the surface temperature of the paint layer rapidly increases to the decomposition temperature and the paint transforms into small particles and gases, forming a high-temperature plume above the sample. The above results will provide a reference for improving laser paint removal efficiency and respraying quality and reducing substrate thermal damage.

Ti60 is a near-α titanium alloy with good high-temperature performance that has been identified as an important candidate material for aero-engine compressor blades and integral blades. However, when high-temperature titanium alloys are fabricated using traditional processing technology, it has the disadvantages of difficult formation, low material utilization, and high cost. Laser cladding technology uses a laser with high energy density to melt the powder preset on the surface of the substrate , so as to obtain the expected performance of the cladding layer. There are many parameters of the laser cladding process that have significant influence on the forming quality. At the same time, complex thermal cycling in the laser cladding process leads to differences in the grain size, morphology, and size of the precipitated phase, which makes the differences in the mechanical properties of the laser cladding significant. Therefore, this paper mainly studies the effect of the process parameters on the forming quality of laser cladded Ti60 alloy, and the microstructure evolution and tensile properties of laser cladded Ti60 alloy are analyzed to lay a theoretical foundation for the application of laser cladded high-temperature titanium alloy components in the aerospace field.

The material selected in this experiment is Ti60 powder with a particle size of 50?150 μm, prepared using the plasma rotating electrode process (PREP). TC4 titanium alloy is used as the substrate, and the laser cladding system is used as the laser cladding experiment system. The section of the laser cladded sample along the thickness direction of the cladding layer is machined via electric discharge wire cutting into a flake sample with a thickness of 5 mm for the metallographic sample. The Kroll reagent is then used for etching, and finally, the microstructure is observed using a metallographic microscope and field emission scanning electron microscope (SEM). A field emission transmission electron microscope (TEM) is used to analyze the precipitated phase of the cladding specimen. A microhardness tester is used to test the Vickers hardness of the Ti60 cladded sample from top to bottom. The tensile experiment is performed on the high-temperature tensile test machine at room temperature, 300 ℃, and 600 ℃, with a tensile speed of 1.0 mm/min. The tensile fracture is observed, and the fracture morphology and fracture mode are analyzed.

The influence of different factors on the size of the laser cladding layer is analyzed according to the shape and size of the cladding layer measured by the image scanner. When the width of the molten pool is large, the cladding efficiency can be effectively improved, the material utilization rate can be improved, and the cost can be reduced. The thickness of the cladding layer has a significant influence on deposition along the height of the cladding layer (Fig. 5). The microstructure at the top region of the cladded sample is the thin layer of equiaxed grain, and its grain size gradually increases with increasing laser power. In the central region of the sample, the original β grains can be observed growing in the deposition direction along the epitaxial columnar pattern across multiple cladding layers. Moreover, the larger the laser power, the coarser the columnar grains and the microstructure inside the grains (Fig. 6). The sample of the laser cladded Ti60 block is mainly composed of a netted basket of lath α and interlath β phases. There are white Ti₅Si₃ phases with different shapes on the slat α, and the content of the Ti₅Si₃ phase gradually decreases from the bottom to the top (Fig. 9). At room temperature, the tensile strength and yield strength of the laser cladded Ti60 samples are 1128 MPa and 1035 MPa, respectively, and the elongation and section shrinkage are 8.8% and 14.4%, respectively. At 300 ℃, the tensile strength and yield strength are 932 MPa and 796 MPa, respectively, and at 600 ℃, the tensile strength and yield strength are 739 MPa and 627 MPa, respectively (Fig. 12).

The microstructure at the bottom and top regions of the laser cladded Ti60 alloy sample is composed of β equiaxed grains, and the middle region is composed of β columnar crystals. Its size gradually increases with increasing laser power. The microstructure is mainly composed of lath α and interlath β phases, and there is a large amount of the white precipitated phase in the lath α phase. With an increase in laser power, the microstructure changes from a net basket structure to a Weisberg structure. The micro-hardness distribution of the bulk sample is uniform, and its hardness value fluctuates in the range 420?440 HV. The tensile strength of laser cladded Ti60 alloy at room temperature is 1128 MPa, and the elongation and section shrinkage after fracture are 8.8% and 14.4%, respectively. When the temperature is 300 ℃ and 600 ℃, the tensile strength is 932 MPa and 739 MPa, respectively.

This study explores the utilization of laser technology to alter the surface structure of metals, producing rainbow-colored effects. Beyond aesthetic purposes, this method has potential applications in data storage and anti-counterfeiting measures. Laser modification of metal surfaces presents several advantages over alternative techniques: reduced processing time, user-friendly operations, rapid molding speeds, durable structures, and an environmentally-friendly process that produces no pollutants. However, present strategies for creating rainbow structures suffer from sluggish laser scanning rates, leading to slower molding speeds. Many of these approaches rely on laser polarization and interference, which produce relatively basic patterns. Storing intricate data requires multiple scanning passes. To address these challenges, this study employs nanosecond lasers to inscribe grating patterns directly onto stainless-steel surfaces, generating vibrant rainbow hues. A comprehensive statistical analysis is conducted to examine the influence of factors, such as laser power, scanning speed, scanning interval, and repetition frequency, on the shape transformations of the samples. By melding a range of rainbow effects, insights into the laser interaction with stainless-steel surfaces are obtained. Through systematic parameter adjustments, the mechanisms behind laser-induced rainbow coloration on stainless-steel surfaces are identified. Then, a composite structure, which combines grating diffraction rainbow hues, thin-film interference color, and the metal inherent color, is developed. This composite showcases a dynamic dominant color shift among its three colors, depending on the viewing angle. This investigation offers valuable insights for potential industrial implementations, seeks to enhance molding efficiency, and introduces innovative solutions for custom metal surface coloration.

In this study, the relationship among the scanning speed, scanning interval, and repetition frequency is investigated to improve the molding speed. The variation trends among them are calculated and analyzed. Scanning speed of 400 mm/s, repetition frequency of 20 kHz, interval of 0.02 mm, and energy percentage of 55%?65% are chosen to observe the corresponding structural changes, and the effect of laser energy on the experimental results is analyzed. The scanning speed of 400 mm/s, repetition frequency of 20 kHz, energy percentage of 80%, and scanning interval of 0.04?0.07 mm are selected to observe the trend of structural changes and to analyze the effect of the scanning interval on the results. The experimental data are summarized, and the forming effect of the proposed scheme is verified. The parameter range required to achieve rainbow colors is determined, and the mechanism behind the changes in the stainless steel surface structure is deduced. Finally, a composite structure is formed after a single scan.

We speculate that this rainbow-colored structure forms due to laser-induced effects (Fig.5). Laser pulses superimpose on each other, and the parts that do not overlap, due to the absence of subsequent energy input, cool down. The overlapping parts maintain a high-temperature state because of the relative temperature difference, leading to wavy structures in the melt pool. Meanwhile, the translation mode of the laser causes the relative displacement between the pulses to be minimal, creating horizontal overlapping lines along the continuous upper and lower edges. The second pulse continues to sweep over the previous pulse, partially overlapping it. This portion of the energy wave either disrupts the previously formed wavy structures, creating new structures, or if the energy is too high, prevents the wavy parts from cooling down promptly. At the same time, the upper and lower edges keep forming horizontal intersection lines, and subsequent laser beams continuously superimpose in vertical and parallel directions of the laser, resulting in a rainbow effect (Fig.6). At a scanning interval of 0.1 mm, a composite structure consisting of a grating diffraction iris, film interference structure color, and intrinsic color forms (Fig.7).

Using a standard nanosecond fiber laser, the surface of stainless-steel plate is directly etched to realize a high-speed rainbow structure. The results demonstrate that the rainbow structure is influenced by the laser energy parameters and scanning path. Under optimal conditions, structures are formed quickly. An increase in laser energy is found to compromise the rainbow structure from the outset. The ripple shape of the rainbow is determined by the laser scanning path, repetition frequency, and speed. By adjusting the laser pulses, a melt-pool boundary that adheres to the Bragg interference condition is formed, producing a wavering effect that spans the entire visible spectrum. Simultaneously, a composite structure that includes the grating diffraction iris, film interference structure color, and intrinsic color is constructed using a specific process. In this structure, alternating dominant colors for the three hues at different angles are achieved, offering a new personalized surface color scheme.

The 2024-T351 aluminum alloy, which belongs to the Al-Cu-Mg series and is extensively used in the aerospace industry, exhibits inadequate corrosion resistance that affects its reliability and service life. Laser peening (LP) has emerged as a novel surface treatment technology capable of enhancing both mechanical properties and corrosion resistance of 2024-T351 aluminum alloy. However, conventional LP requires an additional energy protective layer, limiting its practical industrial applications. LP without coating (LPwC) offers a promising alternative to conventional LP. The present study focuses on investigating the 2024-T351 aluminum alloy and employs a Nd∶YAG laser to perform LPwC treatment on the sample surfaces. By utilizing the experimental characterization techniques including surface morphology analysis, chemical composition examination, microstructure observation, phase analysis, and residual stress measurement, a comparative analysis of the corrosion behavior is conducted on the LPwC samples with laser power densities of 1.1 GW/cm² and 2.6 GW/cm² in the NaCl solution with mass fraction of 3.5%. The corrosion resistance mechanism of the LPwC process is also elucidated.

The experimental samples processed by the Nd∶YAG pulsed laser were selected from 10 mm×10 mm×2 mm square specimens. First, the surface and cross-sectional morphologies of the samples were observed using a scanning electron microscope (SEM). Second, the element compositions of the selected areas were analyzed by the energy dispersive spectrometer (EDS) attached to the scanning electron microscope. Third, the wetting performance of the samples was determined using a angular contact measuring instrument. Subsequently, the microstructure of the samples was examined with the SEM, and the X-ray diffraction (XRD) analysis provided information on the phase compositions of their surfaces. Furthermore, the X-ray stress tester allowed for measuring the residual stress distribution along the depth direction in these samples. Finally, electrochemical corrosion experiments were conducted on the electrochemical workstation to characterize corrosion resistance properties of these samples after immersing them in an electrolyte solution for 30 min.

Surface morphology observation revealed that after the LPwC treatment, the samples exhibited a multi-level structure with staggered micro-nano bumps and holes (Fig.3). Water contact angle measurements demonstrated the hydrophobic properties of the LPwC-treated samples, with 131° for the L-1 sample and 112° for the L-2 sample, effectively reducing the contact area in corrosive solutions (Fig.4). EDS analysis confirmed the formation of a remelted oxide layer with a thickness of 2?3 μm on the surfaces of the L-1 and L-2 samples (Fig.5). Grain size statistics indicated a 25% reduction in average grain size for the L-1 sample and a 37.5% reduction for the L-2 sample (Fig.6). Residual stress measurement showed maximum residual compressive stress at a depth of 100 μm for both L-1 and L-2 samples, with amplitudes of -100 MPa and -115 MPa, respectively (Fig.8). Furthermore, electrochemical corrosion experiments revealed corrosion inhibition efficiencies of 97.30% for the L-1 sample and 84.63% for the L-2 sample, highlighting significantly improved corrosion resistance (Fig. 9).

The influence of surface morphology, chemical composition, microstructure, phase composition, and residual stress on the LPwC-treated samples with different laser power densities was discussed in this study. The electrochemical corrosion behavior was analyzed and the corrosion resistance mechanism of the LPwC process was summarized. The main conclusions were as follows: 1) The thermodynamic coupling effect induced by LPwC resulted in the formation of a micro-nanoscale bumps and holes staggered multi-level structure and a 2?3 μm thickness dense remelted oxidation layer on the sample surface. This reduced the actual contact area between the matrix and the corrosion solution, improved the chemical activity of the surface material, and enhanced corrosion resistance in corrosive environments. 2) LPwC treatment led to grain refinement and residual compressive stress effects on the surface layers of aluminum alloy samples. The grain sizes of LPwC-treated samples at 1.1 GW/cm² and 2.6 GW/cm² were reduced by 25% and 37.5%, respectively. And the maximum residual compressive stresses of -100 MPa and -115 MPa were formed at a depth of 100 μm. This enabled tensile-compressive transformation within the matrix, inhibiting corrosive ion transmission and corrosion crack propagation within samples while significantly delaying overall corrosion. 3) Laser power density has a significant impact on the corrosion resistance of LPwC-treated samples. Under the 1.1 GW/cm² condition, a micro-nano multilevel structure formed on sample surface exhibited a superior hydrophobic characteristic, with a water contact angle reaching 131°, 17.0% higher than that for the LPwC-treated sample at 2.6 GW/cm². The composite interface created by this multilevel micro-nano structure along with an air micro-cushion layer effectively reduced solid-liquid contact area, resulting in an 82.46% reduction in the corrosion current density of the sample, demonstrating an exceptional corrosion resistance performance.