Titanium matrix composites have attracted considerable attention because of their high modulus of elasticity, high specific strength, high wear resistance, and excellent high-temperature durability. Most studies on titanium matrix composites (TMCs) focus primarily on the in-situ formed TiC reinforced composites. However, few studies have focused on the direct addition of TiC-reinforced titanium matrices. The manners in which the size, morphology, and distribution of TiC evolve during the SLM process and how they affect the microstructure and mechanical properties remain unclear. In this study, TiC/TC4 composites with directly added nanoscale TiC particles are successfully prepared by selective laser melting (SLM), and the microstructure evolution under different volume energy densities is investigated. Further, the TiC evolution during SLM and its influence on the microstructure and microhardness are analyzed. Thus, the findings of this study can provide the support for SLM preparation of titanium composites.

Herein, nanoscale TiC (diameter of 50?150 nm) and TC4 are selected as the reinforced phase and matrix, respectively. The composite powder with TiC uniformly embedded on the surface of the TC4 powder is obtained by low-energy ball milling. Subsequently, the TiC/TC4 composites are prepared via SLM with different volume energy densities (29?97 J/mm³). The forming quality and microstructures at different volume energy densities are observed using optical microscopy (OM) and scanning electron microscopy (SEM) equipped energy disperse spectroscope (EDS). The grain size and crystal orientation are investigated using electron backscattering diffractometer (EBSD), and the phase compositions are measured using X-ray diffraction (XRD). Finally, the microhardness is measured using a digital microhardness tester.

The optimized volume energy densities for the SLM formed TiC/TC4 composites are in the range of 50?70 J/mm³, with a relative density of 99.7% (Fig.3). Owing to the enrichment of TiC in the melt pool boundary zone, the microstructure of the composites exhibits a special double-sized grain distribution in the cross section (Fig.6). Owing to the rapid cooling characteristics of the SLM process, TiC cannot be sufficiently dissolved. Therefore, the SEM and EBSD results reveal three types of reinforcement: undissolved TiC, eutectic TiC, and precipitated TiC. Undissolved TiC is distributed primarily at the boundaries of coarse β equiaxed grains, eutectic TiC is distributed primarily in the boundaries of irregular eutectic β grains, and precipitated TiC is distributed primarily in the grains. With an increase in volume energy density, the chain-like eutectic TiC gradually transforms to rod-like eutectic TiC (Figs.7 and 8), the size of precipitated TiC inside the grain gradually increases, and the sizes of longitudinal and transverse α'-Ti gradually increase.

The optimal volume energy density for the formation of TiC/TC4 composites by SLM is 50?70 /mm³, and the relative density is 99.7% within this parameter range. TiC is enriched in the melt-pool boundary region under a strong temperature gradient and Marangoni convection. The microstructure of the composite has a special double-size grain distribution in the cross section, consisting of primary β equiaxed grains and irregular eutectic regions growing on the periphery. In the longitudinal section, the molten pool is a fish scale, and some chain structures exist in the molten pool that grow from the direction of heat flow to the horizontal direction. With an increase in volume energy density, the size of primary β equiaxed grains decreases, outer-ring irregular eutectic region expands, and morphology of fish scales becomes sharp. The microhardness initially decreases and then increases, essentially reaching 385?392 HV in the optimal molding process window. TiC in the composites is composed primarily of undissolved TiC (distributed near the primary β grain boundaries), eutectic TiC (distributed in the eutectic β grain boundaries in a chain or rod-like network), and precipitated TiC (distributed in the grain in a granular manner). With an increase in volume energy density, the difference in TiC size and quantity inside and outside the molten pool increases, chain distribution of eutectic TiC changes to rod, and the size of TiC in the grains increases. Further, no obvious orientation relationship between eutectic TiC and β-Ti is observed; however, a distinct orientation relationship between eutectic and in-grain TiC and α'-Ti exists: {11?20} α'-Ti∥{110}TiC.

In recent years, significant progress has been made in preparing conformal cooling dies for die casting using additive manufacturing technology. Among these advancements, 18Ni300 maraging steel has been widely applied because of its excellent forming characteristics. Currently, most research on selective laser melting (SLM) manufacturing of 18Ni300 maraging steel has primarily focused on the changes in microstructure after a heat treatment and the influence of precipitate phases on the strength, with limited emphasis on the impact of toughness. However, toughness plays a crucial role in determining the service life and safety of the molds.

Although previous studies have explored reverse-austenite, systematic research on the toughness of 18Ni300 is currently lacking. Therefore, this study aims to systematically investigate the impact of the aging and solution temperatures on the microstructure and mechanical properties of 18Ni300 maraging steel. Additionally, it will specifically analyze the influence of reverse-austenite on the strength and plasticity of 18Ni300 maraging steel prepared using SLM technology. This study clarifies the relationship between the process, structure, and performance of 18Ni300 maraging steel, and proposes an optimal heat-treatment system. These findings offer valuable guidance for the practical application of this steel in various industries.

In this study, 18Ni300 powder was used as the raw material. Experimental samples were obtained through selective laser melting (SLM) using an appropriate method. Following the formation, the samples were subjected to various heat treatments. The bulk samples were ground and polished with sandpaper, followed by etching with a 4% nitric-acid solution in alcohol. The microstructure was examined using optical microscopy (OM) and scanning electron microscopy (SEM). The mechanically polished samples were additionally polished with SiO₂ and the crystal structure of the material was analyzed using electron backscatter diffraction (EBSD). X-ray diffraction (XRD) was utilized to analyze the phase composition and determine its content. Finally, tensile tests were conducted at room temperature using a universal testing machine and the corresponding fracture surfaces were observed.

The morphologies of the tested samples are shown in Figure 3. The printed sample displays distinct fish-scale-like fusion pools and lath martensite structures, whereas the honeycomb-like microstructure is not discernible in the SEM image. Following the aging treatment, the boundaries of the fusion pools in the samples become indistinct, and the boundaries of the honeycomb-like microstructure in the SEM image begin to dissolve. In the solution and aging-treated samples, the boundaries of the fusion pools vanish completely, and the martensite is transformed into a more refined structure. Additionally, the honeycomb-like microstructure observed in the SEM image also completely disappears.

The XRD analysis of the samples reveals that the phase composition of the as-printed sample comprises martensite and residual austenite, whereas the aged sample consists of martensite, residual austenite, and reverse-austenite. Almost the entire microstructure of the solution- and aging-treated sample is composed of martensite. Figure 5 shows that the highest amount of reverse-austenite is observed in the aged sample. Furthermore, Table 3 indicates that the sample aged at 490 °C exhibits the highest content of reverse-austenite.The mechanical properties of the sample are closely correlated with the reverse-austenite content, as depicted in Figure 8. Notably, the sample aged at 490 °C exhibits greater toughness with only a marginal reduction in strength. However, the relationship between austenite and the strength toughness of 18Ni300 is not a simple linear correlation because of factors such as precipitates and the martensite morphology. Overall, it is evident that reverse-austenite significantly enhances the toughness and marginally decreases the strength. With an increase in the reverse-austenite content from 0.1% to 6.9%, the elongation after fracture improves by 72.5%, whereas the tensile strength decreases by 2.3%.

The printed samples of 18Ni300 maraging steel manufactured by SLM display a distinct molten pool and a microstructure comprised of coarse martensite and a small proportion of residual austenite. Following the aging treatment, a ductile phase called reverse-austenite is generated. After the post-solution and aging treatments, the microstructure exhibits uniform and dense plate-like martensite with no notable presence of the austenite phase. A direct aging treatment at 490 °C is considered the optimal heat-treatment process for achieving an ideal balance between strength and toughness. At this temperature, the microstructure exhibits the highest reverse-austenite content (volume fraction: 7.7%). The ultimate tensile strength is 2012.8 MPa, and the elongation after fracture reaches a peak value of 6.9%. Therefore, a direct aging treatment at 490 °C is regarded as the most optimal heat-treatment process.

The fine reverse-austenite within the maraging steel manufactured via SLM serves as a toughening phase, enhancing the toughness without significantly compromising the strength. With an increase in the reverse-austenite volume fraction from 0.1% to 6.9%, the elongation after fracture experiences a 72.5% improvement, albeit at the expense of a 2.3% decrease in the ultimate tensile strength. Thus, the reverse-austenite is advantageous for achieving exceptional overall mechanical properties in maraging steel manufactured via SLM. The fine reverse-austenite plays a pivotal role in enhancing themaraging steel. However, in the maraging steel manufactured via SLM using 18Ni300, precipitation strengthening constitutes the primary mechanism with a limited effective range of precipitation temperatures. Further research is necessary to increase the reverse-austenite content, while maintaining adequate precipitation strengthening.

Selective laser melting (SLM) technology offers a wide range of design freedom, high density, and strong metallurgical bonding; thus, it is highly suitable for processing workpieces with complex shapes. A conformal cooling mold formed via SLM can improve the cooling efficiency and decrease the injection cycle time. However, only a few types of mold steel materials are suitable for 3D printing because of long processing time and high costs. First, conventional processes can be employed to manufacture conventional parts; subsequently, complex parts can be built using SLM. With this approach, the manufacturing efficiency can be improved and costs can be reduced. In this study, a new type of 3D printing die steel material (AM40) is deposited on a commercial H13 substrate using SLM. The effects of heat treatment (HT) on the microstructure and mechanical properties of AM40/H13 bimetallic structural materials are studied, and the deformation and cracking behaviors of the bimetallic molds are revealed.

In this study, AM40 steel powder and annealed H13 steel sheets are used. SLM is used to deposit AM40 onto the H13 substrate. Subsequently, quenching and tempering are performed to study the effects of the heat treatment. The particle size distribution is characterized using a laser particle size analyzer, whereas the microstructure and fracture morphology are characterized using optical microscope (OM) and scanning electron microscope (SEM). The grain morphology, orientation, and local misorientation of the bonding interface are characterized using electron backscatter diffraction (EBSD). Additionally, a Vickers microhardness tester is employed to measure the microhardnesses of the as-built and heat-treated samples. Tensile tests are performed using a fatigue testing machine.

No crack defect is observed at the interface of the as-built AM40/H13 bimetallic structure and the unique Marangoni molten pool at the interface (Fig. 7). Moreover, fine cellular and columnar martensite structures are observed in the AM40 region (Fig. 8). The microstructure of H13 is coarsened austenite (Fig. 5), and the distinct microstructural inhomogeneity is observed at the bonding interface. After quenching and tempering, the characteristics of the molten pool disappear, and uniform lath martensite microstructures form in the H13 region (Fig. 8). The inhomogeneity of the grain size and misorientation at the interface are eliminated (Fig. 10). Moreover, the diffusion width of element at the interface increases from 440 μm to 500 μm (Fig. 9). Additionally, the hardness of the as-built AM40/H13 at the bonding interface is 642 HV, which is higher than those of AM40 (529 HV) and H13 (202 HV). The average hardness of HT-AM40/H13 at the bonding interface decreases to 480 HV (Fig. 11), thus indicating that the hardness difference between AM40 and H13 is eliminated by the heat treatment. The tensile strength of HT-AM40/H13 increases significantly from 644 MPa to 1436 MPa (Fig. 12). Furthermore, some dimples, along with a cleavage pattern, are observed in the fracture (Fig. 14), thus indicating that the fracture mode is a combination of ductile and brittle. The increase in the tensile strength and ductility of the heat-treated AM40/H13 bimetallic alloy is analyzed based on the microstructure and fracture morphology of the bonding interface.

In this study, the as-built AM40/H13 bimetallic structure does not exhibit crack defects at the interface, and the microstructure is heterogeneous. Marangoni convection and cellular and columnar structures are observed in the weld pool at the interface. The alloying elements are evenly distributed at the interface, thus indicating good metallurgical bonding. After heat treatment, the grain size and dislocation density near the interface are similar, thus eliminating the inhomogeneity of the interface structure. The elements at the interface diffuse, and the diffusion width increases by 60 μm. The hardness at the as-built AM40/H13 bimetallic H13 side is the lowest (202 HV), followed by that at the AM40 side (529 HV); by contrast, the interface hardness is the highest (642 HV). Tensile deformation and cracking of the bimetal preferentially occur at the H13 side, with a strength of 644 MPa and fracture elongation of 29%, thus indicating ductile fracture. After heat treatment, the hardness of H13 increases to 483 HV, which is equivalent to that of AM40 (479 HV) after heat treatment, and the inhomogeneity of the hardness is eliminated. In addition, the tensile strength of HT-AM40/H13 increases significantly from 644 MPa to 1436 MPa, which is between those of AM40 and H13. The fracture is preferentially located at the AM40 side, far from the interface. Further, some dimples and cleavage patterns are observed, thus indicating that the fracture mode is a combination of ductile and brittle.

Laser melting deposition is used to prepare ATI 718Plus samples to study the effects of three heat treatment regimes on their microstructure evolution, hardness, and room temperature tensile properties. These regimes include direct aging heat treatment, solutionizing and aging heat treatment at 982 ℃, and high-temperature solutionizing and aging heat treatment at 1020 ℃. The aim is to elucidate the phase transformation behavior and mechanical property changes of laser additive ATI 718Plus under different heat treatment regimes and provide guidance for the selection of heat treatment processes used in the laser additive manufacturing of ATI 718Plus.

The plasma rotating electrode process is used with ATI 718Plus powder with particle diameter of 45?105 μm to prepare wrought ATI 718Plus superalloy substrates. The experiments are performed on the laser additive manufacturing system shown in Fig. 2, which consists of a 4000 W continuous wave fiber laser, an inert atmosphere processing chamber, a coaxial nozzle, and a powder feeding device. A well-formed ATI 718Plus sample is prepared using a unidirectional reciprocating scanning method with the following parameters: a laser power of 1200 W, scanning speed of 0.8 m/min, protective gas flow rate of 10 L/min, carrier gas flow rate of 15 L/min, and powder feed rate of 13 g/min. The geometric dimensions of each sample are 50.0 mm×58.0 mm×2.5 mm. Three heat treatment regimes are employed, as shown in Fig. 2(b). The analyzed samples are mechanically ground with SiC paper and polished using diamond suspensions and a colloidal silica suspension to prepare metallographic samples. Then, the polished samples are etched with No.2 waterless Kailing's reagent for optical microscope and scanning electron microscope (SEM) investigations. Uniaxial tensile tests are carried out at room temperature using a universal testing machine with a constant displacement rate of 1 mm/min.

After laser deposition, a large number of Laves phase areas form in the interdendritic region (Fig. 3). This hard and brittle phase deteriorates the mechanical properties of the additive-manufactured ATI 718Plus samples. The as-deposited sample mainly exhibits an epitaxial growth columnar dendritic morphology, with a large number of brittle long-chain Laves phases precipitated between dendrites, which consumes a significant amount of Nb, Mo, and other strengthening elements, severely reducing the mechanical properties of the as-deposited sample. After the direct aging heat treatment, the long-chain Laves phase morphology remains unchanged, and the η and γ′ phases precipitate heavily between dendrites. The solution and aging heat treatment system can effectively reduce the size and content of the Laves phase. With an increase in the solution temperature, the size and content of the Laves and η phases gradually decrease, and the γ′ phase uniformly precipitates. The hardness significantly increases after heat treatment (Table 2), but the hardness differences between the three heat treatments are relatively small. The room temperature tensile properties are shown in Fig. 8. Compared to the as-deposited sample, after heat treatment the samples exhibit significant increases in both the yield strength and tensile strength, while the elongation at fracture decreases and then increases. The yield and tensile strengths increase by 67.7% and 51.9% after the direct aging heat treatment, respectively, while the elongation at fracture decreases by 13%. After the solution aging (SA) heat treatment at 982 ℃, although the strength improvement is not as significant as that after the direct aging treatment, the yield and tensile strengths still increase by 63.6% and 45.6%, respectively. At the same time, the elongation at fracture increases by 3% compared to that of the as-deposited state. The strength improvement is the smallest after the 1020 ℃ SA, with a yield strength increase of only 62.0% and tensile strength increase of 34.2%, but the plasticity is significantly improved, with an elongation at fracture increase of 25.8% compared to that of the as-deposited state.

The strength and hardness values of the ATI 718Plus additive samples significantly increase after heat treatment. The best match between strength and plasticity is obtained after high-temperature solution and aging heat treatment at 1020 ℃. Compared with those of the as-deposited state, the tensile strength and elongation at the fracture of the sample increase by 34.2% and 25.8%, respectively, after the 1020 ℃ solution and aging heat treatment.

Laser powder bed fusion (LPBF) is an additive manufacturing (AM) process that has the advantages of forming complex-shaped parts and cutting costs. It is widely used in the aerospace, medical equipment, weapons manufacturing, and other industries. However, in the LPBF process, the material powder is repeatedly heated and melted under the effect of laser energy and then cooled and solidified, which facilitates the formation of a large thermal gradient and thermal stress in the parts, leading to warping deformation. This type of deformation significantly affects the dimensional accuracy and mechanical properties of parts. By combining sensor signal acquisition with data analysis, deformation defects can be detected during AM to reduce production costs and improve the quality of formed parts. The radiant light signal of the molten pool is sensitive to the thickness of the powder layer during the LPBF process, which may reflect the warping deformation that has already occurred. It is also correlated with the temperature of the molten pool, reflecting the peak temperature at that location, and is related to the temperature field of the sample. Therefore, it has the potential to monitor the thermal stress during warping deformation. To study the relationship between thermal stress-induced warping deformation and the radiant light signal of the molten pool, a method for monitoring warping deformation in the LPBF process by acquiring the radiant light signal of the molten pool is explored in this study. In this study, an overhanging sample is formed during the experiment, and the radiation signal of the molten pool is collected and analyzed. The results show that the radiant light signal can not only monitor warping deformation but also reflect formation process of warping deformation to a certain extent.

To collect and compare the radiation light signal of the molten pool during the forming process of the warped and normal samples, T-shaped overhanging structure samples are formed (Fig.2), and five samples with three different support structures and sizes are designed for the experiment (Table 1). In this process, three sensors collect the radiation intensity signals from the molten pool, and an upper computer records the coordinate data of the laser spots (Fig.1). After data alignment, each light intensity value corresponds to the coordinates of the laser spot during scanning. To further explain the variation trend of the light intensity signal along the long side (Y-direction) of the sample, the scanning section of the sample is divided into regions, and the average light intensity of each region is calculated. Three measurement points are selected on the sample, and the heights of the measurement points relative to the substrate plane are measured using a coordinate apparatus.

No evident warping deformation is observed in the forming process of samples S80-1 and S80-2, whereas the warping deformations of samples S25-1, S25-2, and S20 are larger (Fig.6). This result indicates that samples with smaller support areas are prone to warping deformation; however, no noticeable linear correlation is observed. The normal samples S80-1 and S80-2 produce a larger average light intensity at both ends, with a minimum value of 0.93 V, while warped samples S25-1, S25-2 and S20-1 produce lower light intensity at the same area (Fig.7). This phenomenon indicates that sample warping can be distinguished from the light signal of the molten pool. The light intensity distribution of the first overhanging layer is different between the warped and normal samples. The light intensity of the warped sample in the region where the corresponding lower layer is solid is significantly higher than that in other regions, forming a “wave peak” in the curve (Fig.8). The above phenomena indicate a correlation between the radiant intensity distribution and peak temperature at the corresponding position and reveal that the evolution trend of the light intensity between the layers of the samples with the same geometric structure. The light intensity of the normal sample fluctuates more between layers, whereas that of the warped sample fluctuates less (Fig.9).

In this study, three types of overhanging samples with different structures are formed, and the radiation light signal of the molten pool is collected. Combined with sample deformation measurements and statistical methods, the data are analyzed, and the following conclusions are obtained:

1) In the layer after warping deformation, the light intensity of the warped specimen decreases significantly in the warped region, while the distribution of the light intensity of the normal specimen is uniform without a notable gradient.

2) For the warped specimen, when the overhanging layer has just been formed, and the deformation has not yet occurred, the light intensity "crest" corresponding to the central solid region of the specimen is quite different from the light intensity in other regions of the layer.

3) The interlayer evolution trends of the light-intensity values of the warped and normal samples are different. With an increase in the number of formed layers, the influence of the overhanging structure on the light intensity signal gradually decreases, and the light intensity tends to stabilize after the fifth layer.

4) A sample with a smaller support area is more likely to produce warping deformation, but no notable linear correlation exists between these two factors.

The parameters for selective laser melting (SLM) directly affect the morphology and microstructure of the melt-forming process, which in turn affect the mechanical properties of the formed structure. Metal powder rapidly heats up and melts under high-speed laser irradiation, forming a metal molten pool. The complex heat and material exchange processes inside and outside the molten pool are difficult to detect in real-time using monitoring instruments. To address the defects generated during the SLM forming process of an AlSi10Mg alloy, this study employed experimental and numerical simulation methods to investigate the effects of forming parameters such as laser power and scanning speed on the morphology of single- and double-channel of the AlSi10Mg alloy.

FS271M laser selective melting equipment was used for single- and double-channel SLM forming of the AlSi10Mg powder. The aluminum substrate was preheated to 130 ℃, the forming cavity was filled with high-purity argon gas as a protective gas, and the oxygen volume fraction was controlled to be less than 0.15%. Table 2 lists the forming parameters. The melt length was set to 20 mm. To facilitate subsequent observation, 1 mm spacing was set for single-melt scanning, and the forming process was repeated five times. After forming and cooling, the morphology of the melt was observed and analyzed using an AM7031MT digital microscope. In addition, Flow-3D v11.1 software was used to simulate the single-channel laser selective melting forming process. A numerical simulation was conducted to investigate the physical effects and phenomena such as thermal radiation, heat conduction, solid-liquid phase transition, molten pool evaporation, gravity, surface tension, and the Marangoni effect derived from the SLM process.

Under different scanning speeds using a laser power of 300 W, the overall continuity of the formed melt is good, no obvious spheroidization is observed, and the degree of overlap is high. As the laser-scanning speed decreases, the width of the melt gradually increases, and a clear ripple morphology is generated at a scanning speed of 700 mm/s. When a 100 W power laser is used for melt forming, the discontinuity and spheroidization of the melt are more severe. The width of the laser heat-affected zone decreases with an increase in the laser scanning speed. The lower the scanning speed, the more obvious is the degree of oxidation and blackening of the powder molten pool. The oxidation effect of the AlSi10Mg powder during processing is a major reason for the low density of the formed structural components. In practical experiments and production, the first-layer premelting method can be adopted to consume as much residual oxygen in the cavity as possible, reducing negative oxidation effects during the molding process. Under the action of a low scanning speed and high energy density laser, the spattering and airflow of the molten pool become more intense, making it easier to produce small-particle spheroidization defects on the forming plane. The keyhole depth generated by the metal molten pool under steam recoil pressure can reach 100 μm. As the laser moves, the molten pool rapidly cools and solidifies due to the high thermal conductivity of the aluminum alloy materials. If the keyhole is not completely filled by the molten pool fluid, pore defects form. Therefore, avoiding keyhole generation while ensuring the continuity of the melt path is necessary. The discontinuity of the melt path is mainly caused by insufficient melting of the powder layer. Reducing the thickness of the powder layer can improve the discontinuity caused by insufficient energy. However, the selection of SLM forming parameters should consider the product-forming efficiency while ensuring the quality of structure forming. Reducing the thickness of the powder layer prolongs the structure-forming time and affects the forming efficiency, and increasing the preheating temperature reduces the energy required for melting. To investigate the effects of the preheating temperature on the morphology of the formed channel, a laser power of 100 W and scanning speed of 800 mm/s were selected as scanning process parameters, and the preset environmental temperature T₀ was gradually increased for calculation. At T₀=500 K, the discontinuity phenomenon in the forming area is eliminated.

This study investigated the single-layer melt forming of AlSi10Mg powder material through experimental and numerical simulation methods. It was found that the surface tension and melt recoil pressure play crucial roles in the evolution and motion of the molten pool. Even when high-purity argon gas is used as the protective gas for the experiment, because of the oxidizability of the AlSi10Mg material, residual oxygen still affects the quality of the melt forming. Therefore, the oxygen content in the forming cavity should be minimized as much as possible prior to forming. Because the AlSi10Mg alloy powder has a weak laser absorption ability, the energy absorption rate was set to 12% in this study. For a given powder bed with a thickness of 50 μm, a mobile laser beam with a linear energy density of 200 J/m is required to completely melt the powder layer. Under low-power 100 W laser scanning, because of the low energy density of the laser, the melt channel is prone to discontinuity and large-scale spheroidization. Increasing the input energy density by reducing the scanning speed does not effectively solve the problem of uneven melt channels. Obtaining a smoother filling in the keyhole formed under low-speed scanning is difficult, which reduces the quality of the melt channel formation. By increasing the preheating temperature, the laser line energy density required for melting can be reduced, and the morphology of the melt formed at low power can be improved.

In recent years, laser additive manufacturing based on direct laser deposition has attracted widespread attention because of its flexibility and efficiency. This technology has a wide range of applications and high additive manufacturing efficiency. It is widely used in the aerospace, rail transit, and ship component maintenance equipment fields. However, high-strength aluminum alloys such as those in the Al-Zn-Mg-Cu series have a high content of alloying elements. During the solidification process, the semi-solid range of the alloy may exceed 100 ℃, which can easily leave gaps between aluminum grains. The α-Al layer of "liquid film" generates cracks under the action of thermal stress, making laser deposition repair difficult. AlSi10Mg alloy, as a cast aluminum alloy, has the characteristics of a short semi-solid range and high strength, and is suitable for additive manufacturing and the laser repair of high-strength aluminum alloy components. However, during the laser deposition process, process fluctuations often cause defects such as pores to appear in the components, leading to cracks and ultimately component failure during use. Therefore, exploring a method to eliminate pores in components produced using AlSi10Mg laser deposition is of great significance for improving the mechanical properties and service life of AlSi10Mg components.

Atomized AlSi10Mg alloy powder with a particle size ranging from 53 μm to 150 μm is adopted. During the laser deposition process, the laser power is 2700 W, deposition speed is 600 mm/min, powder feed rate is 5.8 g/min, overlap amount is 2.5 mm, argon flow rate is 5 L/min, and protective argon amount is 20 L/min, resulting in a single-layer thickness of 0.5 mm. After depositing eight layers to achieve a thickness of 4 mm, stir friction treatment is performed on the deposited AlSi10Mg alloy. The height of the mixing needle of the mixing head is 4 mm, with a four-prism shape and diameter of 6 mm at the end of the prism. During the stirring friction treatment process, the rotational speed is 800 r/min, stirring speed is 100 mm/min, and variation in the stirring friction treatment passes is 5 mm. Subsequently, the laser deposition of eight-layer AlSi10Mmg alloy is continued on the surface of the AlSi10Mg alloy after the stir friction machining, and then stir friction machining is used.

Wire cutting is used to cut the AlSi10Mg alloy into five samples, and stir friction-assisted laser deposition is conducted, followed by room-temperature rolling treatment. Rolling deformation values of 20%, 46%, and 68% are achieved on three of the samples. The five tensile specimens of the AlSi10Mg alloy are treated as mentioned above, and their strength and elongation values are measured using a tensile testing machine. After vibration polishing, the five metallographic samples are observed using a scanning electron microscope and backscattered electron diffractometer, and their microhardness values are measured. A thin film sample of the AlSi10Mg alloy is prepared and its microstructure is observed using a transmission electron microscope after electrolytic double spraying. The fracture of the tensile specimen is observed using the scanning electron microscope.

The hardness values of the AlSi10Mg alloy in the five different states are listed in Table 1. It can be observed that the hardness of the deposited AlSi10Mg alloy is approximately 109 HV. Because at high temperatures, the solid solubility of the Si element in the α-Al matrix is relatively high, and when the temperature rapidly drops, it is difficult for the Si element to recover from α-Al matrix, and a large number of Si atoms on α-Al matrix play a role in solid solution strengthening. The Al matrix plays a role in solid solution strengthening. After stir friction processing, the solid solution strengthening effect is significantly weakened, and the hardness of the AlSi10Mg alloy decreases to 75 HV. Based on the hardness values of the rolled AlSi10Mg specimens listed in Table 2, it can be observed that the rolling process improves the effects of dislocation strengthening and fine grain strengthening in the AlSi10Mg alloy, ultimately increasing the hardness of the laser-deposited AlSi10Mg alloy after stir friction processing to 116 HV. As shown in Fig. 9, after stir friction processing, the strength of the AlSi10Mg alloy is close to 200 MPa, and the elongation distribution is 33%?40%. It can be seen that stir friction processing can simultaneously improve the strength and plasticity of the laser-deposited AlSi10Mg alloy. Figure 10 shows that there are a large number of dimples in the tensile fracture surface of the AlSi10Mg alloy in the laser deposition state and stir friction processing state, indicating that the fracture mode of both AlSi10Mg alloy specimens is the plastic fracture mode. The research on hardness shows that the strength and elongation of the laser deposited AlSi10Mg alloy cannot reach high levels. However, after stir friction processing, the larger shoulder pressure and stirring effect eliminate the porosity defects in the alloy, reduce the stress concentration, and thus significantly increase its elongation.

After friction stir processing, the columnar α-Al and eutectic phases in the laser deposited AlSi10Mg become equiaxed α-Al grains and Si particles, while the Mg₂Si precipitate phase is significantly refined. After rolling, when the deformation of the AlSi10Mg alloy increases to 68%, dislocation strengthening further refines the grains.

After rolling, the dislocation strengthening effect in the alloy is significantly enhanced. In addition, increasing the rolling amount also brings about a fine grain strengthening effect. Therefore, the hardness of the AlSi10Mg alloy can be increased to 116 HV at most, exceeding the microhardness of the laser deposited AlSi10Mg alloy.

Although the solid solution strengthening effect in the laser deposited AlSi10Mg alloy is remarkable, the solidification defects in the alloy lead to the formation of early cracks during the tensile process, which results in an alloy strength of less than 200 MPa and an elongation of less than 20%. After friction stir processing, the strength and toughness of the AlSi10Mg alloy are simultaneously improved, with a strength close to 200 MPa and an elongation of 33%?40%. After rolling, the dislocation strengthening effect of the AlSi10Mg alloy gradually increases, and its strength continues to rise, reaching a maximum of approximately 400 MPa. The localized hardening area in the alloy leads to a decrease in its plastic deformation ability, and the elongation gradually decreases to 25%.

Hybrid material structures have various applications in the automotive industry owing to their light weight. Stainless steel, which exhibits good corrosion resistance and remarkable mechanical properties, is widely used in automotive applications. Glass fiber reinforced plastics (GFRPs) that exhibit high specific strength and cost performance have replaced existing materials in applications requiring lightweight materials. Single-side resistance spot welding of stainless steel and GFRP can help combine the advantages of the two materials. However, owing to the difference in the thermal physical properties and chemical structures of these two materials, the combined strength cannot meet industrial requirements. Improving mechanical interlocking and chemical bonding is an effective approach for enhancing the joint performance. The laser joining process can be used to fabricate micro-textures and change the surface chemical state. Thus, micro-textures on the surface of stainless steel are prepared using a nanosecond laser, and the strengthening mechanism of the interface under the influence of the micro-textures is studied.

Initially, 304 stainless steel and GFRP are selected as base materials. The 304 stainless-steel sheets are subjected to laser texturing. The cruciform mesh micro-texture is selected as the basic morphology of the stainless-steel surface. The grid line uses contained multiple equally spaced scan lines, and a laser processing system supporting software is used to preset different micro-texture widths. The number of laser scanning times is set as 10, and the micro-texture width is set as 0.1?0.5 mm. An optical digital microscope and a field-emission scanning electron microscope are used to detect the laser texture, surface morphology, and fracture surface of the joint. A constant-temperature heating platform and a high-temperature wetting angle measurement system are used to measure the GFRP contact angle on the stainless-steel surface to characterize its wettability. A universal material testing machine is used to conduct tensile-shear tests on the 304 stainless steel/GFRP single-side resistance spot welding joints.

The introduction of micro-textures on the surface of stainless steel significantly improves the wettability of the surface. The surface of stainless steel changes from an untreated non-wetting state to a wet state after laser treatment. As the width of the micro-texture increases, the wettability initially increases and then decreases (Fig. 5). When the micro-texture width is 0.2 mm, the wettability reaches the optimum value. The interior of the micro-textures is completely filled with molten GFRP. When the micro-texture width is too large, the molten GFRP cannot completely fill the interior of the micro-textures (Fig. 7). C and Fe diffuse at the interface, and an element diffusion layer is formed (Figs. 8 and 9). When the micro-texture width is 0.2 mm, the tensile-shear force reaches the maximum value of 3548 N, which is 385% higher than that of the untreated stainless steel/GFRP single-side resistance spot welding joint. The tensile-shear force first increases and then decreases as the micro-texture width increases. Compared with the case of the joint without micro-textures, after laser treatment of the stainless-steel surface, a large amount of the resin-glass fiber mixture is observed in the center area of the fracture of the joint (Fig. 11). The fracture mode changes from an interfacial fracture to a mixed form of interfacial and cohesive fractures. Corresponding to the wettability and joint tensile-shear force, the bonded-area ratio first increases and then decreases, indicating an improvement in mechanical properties.

Laser texturing is used to improve the performance of stainless steel/GFRP single-sided resistance spot welding joints. After the nanosecond laser treatment, the wettability of the molten GFRP on the stainless-steel surface is significantly improved, and the state changes from non-wetting to wetting. The introduction of the micro-textures improves the mechanical properties of the stainless steel/GFRP resistance spot welding joint. When the micro-texture width is 0.2 mm, the tensile-shear force of the stainless steel/GFRP single-side resistance spot welding joint reaches the maximum value of 3548 N. Compared to the case wherein the micro-textures are not introduced, the tensile-shear force of the textured joint is 731 N. The introduction of the micro-textures increases the contact area between the stainless steel and GFRP, thereby significantly enhancing mechanical interlocking. When the micro-texture width is suitable, the GFRP completely fills the inside of the micro-textures. When the micro-texture width is too small or too large, the GFRP does not completely fill the inside of the micro-textures owing to the influence of wettability. In addition to mechanical interlocking, Fe and C chemically diffuse at the interface to form a compound layer, which further improves joint strength.

Ductile iron has been extensively used in various automotive components such as crankshafts and differential housing owing to its relatively low density and capacity for significant tensile strength. 20MnCr5 is a robust and tough alloy steel commonly employed in the production of gears and shafts. Establishing effective welding between the shaft body and the gear material is a significant research challenge. However, the notable disparity in the thermal properties between ductile iron and alloy steel hinders the performance of the welding joint. The high carbon content of ductile iron promotes carbon segregation at the welding interface and exacerbates the formation of microcracks, thereby considerably increasing the complexity of the welding process. Owing to its high energy density, laser welding offers the advantage of generating welds with more precise heat-affected zones. In this study, a novel continuous-pulse coaxial dual-beam laser is employed as a welding heat source to enhance the surface quality of the weld seam. The high-quality welding of ductile iron and alloy steel is achieved by decreasing the laser input power and diminishing pore formation. We hope that our novel welding strategy and findings will be helpful in understanding the bonding mechanism of ductile iron and alloy steel and provide more application space for their connectors.

In this study, QT500-7 and 20MnCr5 are employed as the base materials, with ERNiCr-3 as the filling wire. A novel continuous-pulse dual-beam laser is used as the heat source. First, the pulsed laser power is varied with a constant continuous laser power to determine the optimal combination of heat sources. The laser action position is then adjusted to further enhance the weld strength. Microstructures are observed using a metallographic microscope, and mechanical performance testing and analysis are conducted using a tensile testing machine. The microhardness of the weld is measured using a microhardness tester. Additionally, the fracture behaviors of different specimens are analyzed using a field-emission scanning electron microscope.

The use of a continuous-pulse coaxial dual-beam laser as a welding heat source (Fig. 2) produces high-quality welding joints. When the pulsed laser power is varied, the weld formation varies considerably (Fig. 4). The weld seam is found to have no defects, such as cracks or pores. When the laser action position shifts toward the steel side, the heat input on the ductile iron side gradually decreases. This reduction in the heat input suppresses the diffusion of carbon, leading to a significant decrease in the hardness values of the heat-affected and bond zones on the QT500 side (Fig. 13). The cross-sectional morphology of the weld reveals significant changes in the melting amount of the QT500-7 side base material, with the centerline shifting toward the ductile iron side when the laser action position is changed (Fig. 6). The segregation line of carbon caused by the high carbon content of the nodular cast iron is solved by changing the laser position to reduce the heat input on the side of the nodular cast iron (Fig. 7). The best mechanical properties of the joint are obtained under a pulsed laser power of 440 W and offset of 0.2 mm. In summary, a continuous-pulse coaxial dual-beam laser can yield high-quality welding joints. Better dual-beam laser welding parameters can be achieved by adjusting the laser power and action position. Furthermore, carbon segregation issues can be effectively resolved by reducing the heat input on the side of the nodular cast iron by changing the laser action position, and pulsed laser stirring proves useful.

In this study, a coaxial dual-beam laser welding technology is proposed to address the challenges of welding ductile iron QT500 and alloy steel 20MnCr5. The main problems are the precipitation of martensite and ledeburite in the heat-affected and bond zones on the QT500 side, which results in carbon segregation. The pulsed laser power and position are adjusted in this study. When the laser action position is shifted toward the steel side, the decreased heat input suppresses the diffusion of carbon, leading to a significant decrease in the hardness of the heat-affected and bond zones on the QT500 side. The best mechanical properties are achieved under a pulsed laser power of 440 W and laser offset of 0.2 mm. The continuous-pulse coaxial dual-beam laser welding technology not only improves the carbon segregation phenomenon on the ductile iron side but also reduces the formation of welding cracks. Overall, the proposed novel coaxial dual-beam laser welding technology is effective in improving welding quality, specifically for ductile iron and alloy steel dissimilar metals. The joint exhibits high-quality and high-performance characteristics by reducing carbon segregation and minimizing hardness values. This study advances the field of welding and provides a potential solution for the welding of dissimilar metals with different material properties.

Laser and electrochemical hybrid machining is a composite processing method that combines laser and electrochemical processing. It can be used to process hard conductive materials. It can accelerate the electrochemical dissolution rate, avoiding recasting layers, thus improving the surface quality. This study proposes a tubular electrode-coupled laser and electrochemical hybrid machining technology that uses a newly designed tubular electrode. This realizes coaxial transmission of laser and electrochemical energy inside the tubular electrode and controllable coupling at the processing gap, which is suitable for high-quality small hole processing with a high aspect ratio. A coupling mechanism dominated by laser and an electrochemical processing is proposed based on the controllable adjustment of the laser and electrochemical energy at the processing gap. The effects of the temperature rise in the laser irradiation zone on the electrolyte conductivity, current density, liquid-phase mass transfer, and electrochemical dissolution rate, as well as the effects of bubbles and impurities generated during electrolysis on the laser energy. Material removal models for laser and electrochemical hybrid machining are established, and preliminary simulation analysis and experimental research on laser and electrochemical hybrid machining are conducted.

This study introduced a tool for laser and electrochemical hybrid machining with a tubular electrode that confined the electrolyte and laser beam coaxially or asynchronously. In addition, it utilized a coaxial optical fiber inside the tubular electrode to enable total internal reflection of the laser, thereby achieving independent control of laser and electrochemical energy within the tubular electrode. Based on this process, a coupling mechanism for the laser and electrochemical energy was explored, as well as the mechanisms where the laser and electrolysis dominate in the hybrid machining process. By investigating the temporal and spatial distributions of local temperature and stress induced by coupled energy, we study the influence of laser on mass transport and electrode potential in the micro-region of electrochemical machining. A theoretical model for the kinetic behavior of materials removal under the action of hybrid energy was established, and a preliminary simulation analysis of laser and electrochemical hybrid machining was conducted. The results of this study laid a theoretical foundation for the fabrication of complex structures with high quality and aspect ratio.

First, the influence of laser power density on the machining capability of workpiece materials is explored (Fig.2). When the laser power density is low, the laser affects the thermal and electrochemical parameters of the workpiece material and the changes in the electrolyte's electrical conductivity, electrolytic current density, ion diffusion rate, bubble rate, and electrode potential within the machining gap through thermal effects. When the laser power density reaches the electrolyte breakdown threshold, the laser impacts the laser and electrochemical hybrid machining process through both thermal and mechanical effects. Second, based on the controllable adjustment of the laser and electrochemical energy within the tubular electrode, the state changes in the coupling region caused by these energy are classified into three mechanisms: laser-assisted electrochemical machining, laser and electrochemical hybrid machining, and electrolysis-assisted laser machining (Fig.4). Furthermore, through theoretical analysis and preliminary simulation studies, the electric field and current density distributions in the laser and electrochemical hybrid energy field, the flow field distribution, the temperature distribution, and the resulting machining surface are investigated. This facilitates in the evaluation of material removal at different locations on the workpiece during the laser and electrochemical hybrid machining processes. Finally, three-dimensional morphologies of blind holes produced by the only electrochemical machining and laser and electrochemical hybrid machining are compared. The advantages of the hybrid laser and electrochemical processing are confirmed (Fig.9). It successfully manufactures through-holes with a diameter of 1.26 mm and a high aspect ratio of 16∶1 and through-holes with a diameter of 1.25 mm and high aspect ratios of 42∶1 (Figs.10 and 11).

Laser and electrochemical hybrid machining typically suffer from defects such as stray corrosion caused by electrochemical machining and resolidification defects caused by laser machining. To avoid the occurrence of defects and improve the surface quality, this study introduces a tool for laser and electrochemical hybrid machining with a tubular electrode. This enables the coaxial transmission of laser and electrochemical energy within the tubular electrode and the controlled coupling at the machining gap, thereby effectively preventing defects such as stray corrosion and resolidification of layers. This approach is suitable for fabricating complex structures with high quality and aspect ratios. Based on the controllable adjustment of the laser and electrochemical energy, this study proposes mechanisms in which laser and electrolysis dominate, and both cooperate in hybrid machining. The thermal effects of the laser on the laser and electrochemical hybrid machining and the influence of the pulse width of electrolysis on the process are analyzed. This study establishes a theoretical model for the kinetic behavior of material removal under the action of hybrid energy. Preliminary investigations are also conducted on the time and spatial distribution of the hybrid energy field and its impact on the machining surface using simulation models. Through experiments, the advantages of laser and electrochemical hybrid machining are verified. Small holes with a diameter of 1.25 mm and aspect ratio of up to 42∶1 without resolidified layers are successfully produced. This study lays a theoretical foundation for the fabrication of complex structures with high quality and aspect ratio.

Femtosecond laser direct writing (FLDW) has been widely used in material processing to improve material performance due to its high flexibility, true three-dimensional capability, and wide applicability to various materials. Photonic integrated circuits (PICs) constructed by FLDW are advantageous in terms of high stability and strong resistance to interference, making them suitable for applications in optical interconnects, biosensing, quantum communication, and quantum simulation. With the continuous expansion and enrichment of these applications, miniaturization of photonic devices has become an inevitable trend. However, the integration density of PICs is significantly limited by the loss caused by large curvature waveguides (including 90°bending, 180° bending, and S-shaped bending) due to the low refractive index contrast of waveguides produced by single-shot FLDW. Although various methods have been reported to optimize the bending loss of large curvature waveguides, none of them can simultaneously meet the requirements of high integration density and wide applicability range. In this work, we employ a method of multiple laser modifications to enhance the refractive index contrast between the core and cladding of waveguides, optimize the cross-sectional refractive index distribution of the core, and achieve a bending loss as low as 0.64 dB/cm for S-shaped bent waveguides with a radius of 20 mm. Since the modification lines are written inside the waveguide and completely consistent with the bending shape and formation of the waveguide, this method possesses the characteristics of high integration density and wide applicability range, providing an important basis for the miniaturization of PICs.

This paper analyzes the causes of bending loss in waveguides and proposes a method of multiple laser modifications to enhance the refractive index contrast between the core and cladding of waveguides, optimizing the cross-sectional refractive index distribution of the core. Then, the mode field distribution within the bent waveguide and the bending loss of the bent waveguide before and after modification are simulated using professional optical waveguide simulation software, COMSOL and Rsoft, respectively. Finally, S-shaped bent waveguides and modification lines are written in alkaline-earth borosilicate glass using a 1030 nm femtosecond laser. By adjusting the scanning order, center spacing, writing power, angle, density, writing mode, and number of layers of both the waveguide and the modification lines, the mode conversion loss between the straight waveguide and the bent waveguide is effectively reduced, as well as the radiation loss of the bent waveguide. In addition, the central wavelength of the testing laser is set to 808 nm. After adjusting the laser to vertical polarization using a polarization controller, the laser is coupled into the waveguide through a polarization-maintaining fiber. The output light is received by a power meter after removing the scattered light using an iris filter.

In the simulation part, the bending loss of the modified bent waveguide is significantly reduced compared with the unmodified bent waveguide, as demonstrated by the comparison of bending losses before and after modification using Rsoft simulations (Fig.3). The waveguide parameters remain unchanged during the simulations. In the experimental section, cross-sectional microscope images of the bent waveguide before and after modification are compared (Fig.4), and it is observed that the dimensions of the two waveguides are similar, indicating that the added modification lines do not occupy any additional space outside the waveguide. In addition, we provide experimental and simulated mode field distributions before and after adding modification lines, and observe that after adding modification lines, the mode field of the bent waveguide is to some extent closer to the center of the waveguide. Subsequently, different writing orders for the modification lines and the waveguide are designed (Fig.5), and the minimum bending loss is achieved with the optimal writing order. Furthermore, considering the flexibility in writing the modification lines, experimental investigations are conducted on the center spacing between the modification lines and the waveguide, as well as the power of the modification lines (Fig.6), the density and angle of the modification lines (Fig.7), and the number of layers and writing mode of the modification lines (Fig.8). These parameters could alter the refractive index distribution of the bent waveguide cross section, thereby influencing the magnitude of bending loss. Therefore, by selecting appropriate parameter combinations, the bending loss can be minimized.

In this study, we employ a method of inscribing modification lines inside bent waveguides using femtosecond laser to reduce the bending loss. The power of 380 mW, the scanning speed of 40 mm/s, and the depth of 190 μm are selected as writing parameters of the waveguide. Experimental results demonstrate that at a position of 20 mm, utilizing the optimal writing order and the side-writing approach, along with the innermost modification lines positioned at a center spacing of 0.3 μm from the waveguide, a power of 300 mW, an encapsulation angle of 10°, a density of 10°, and a layer number of 2, the bending loss of the S-shaped bent waveguide could be reduced to 0.64 dB/cm. These experimental findings are consistent with the Rsoft simulation results. This method offers a more convenient and flexible option for integration in photonic chips, contributing to further improvements and advancements in their development and applications.

With the rapid development of the aerospace, ship, power, and energy fields, single-crystal Ni-based superalloys have been widely used in aeroengine and gas turbine components because of their excellent comprehensive performance. This has resulted in an increase in the quality requirements for related microhole structures, which has translated to higher processing technology requirements. Waterjet-guided laser drilling technology, when compared with other traditional microhole processing techniques, such as electrochemical machining, electrical discharge machining, and “dry laser” processing, has the advantages of a large working distance, no thermal damage, neat cutting, and no obvious taper. However, the high specific strength and low thermal conductivity of single-crystal Ni-based superalloys make them prone to defects such as poor microhole surface morphologies and large tapers during processing. Hence, it is crucial to investigate the effects of the processing parameters on the microhole surface morphologies and taper for high-quality machining of superalloy microholes.

This study investigates the mechanism and experimental research of waterjet-guided laser drilling of the single-crystal Ni-based superalloy, DD91. First, the effects of the laser single-pulse energy, scanning speed, feed time, and scanning time on the surface morphologies and tapers of microholes are studied by setting up single-factor experiments. Then, based on the single-step spiral scanning mode [Fig.2(a)], a multistep spiral scanning mode drilling method [Fig.2(b)] is proposed to improve the defects of poor microhole surface morphologies and large tapers. In the multistep spiral scanning mode, the coupled energy beam repeatedly scans the innermost circle (circle 1) N times, cut across the material to form a prefabricated hole at the center of the microhole, and then scans the second circle (circle 2) to the outermost circle (circle N) N times with a single-step spiral scanning mode to complete the processing of the filling circle and hence widen the aperture and improve the microhole geometry. Finally, the quality of microhole machining via the single/multistep spiral drilling methods is compared under the appropriate processing parameters. The microhole surface morphologies are observed using optical microscope, the entrance and exit apertures are measured via ultra-depth-of-field microscope, and the corresponding taper is calculated.

During waterjet-guided laser trepanning on metals, material removal is dominated by laser ablation through mechanisms such as photothermal mechanisms, including material melting, evaporation, and sublimation. The water jet, with its high heat capacity, can provide good heat management as well as clean molten material and debris from the ablation zone (Fig.3). As the laser single-pulse energy increases, the material removal rate also increases, which enlarges the exit diameters and causes the taper to increase (Fig.5). A pulse energy that is too low will lead to serious microhole surface morphology damage (Fig.4). With an increase in scanning speed, the ablation time per unit area decreases, which leads to a worsening of the circularity of the hole (Fig.7), a decrease in the exit diameter, and an increase in the taper (Fig.6). As shown in Fig. 9, the entrance diameters of the microholes are all steady at approximately 1025 μm, regardless of how many feeds are applied. The exit diameters increase with an increasing number of feeds and reach a saturation value (approximately 1000 μm) after the feed time is over 6 (Fig.9). Multiple feeds can improve the circularity of the microhole (Fig.8). When the scanning time is 1, the microhole taper is smallest, but the dimensional accuracy is low. With an increase in the scanning times, the quality of the microhole deteriorates, the entrance aperture decreases linearly, the exit aperture first decreases and then becomes saturated, and the taper of the microhole first increases and then decreases (Figs.10 and 11). Based on the above results, the appropriate processing parameters are selected to compare the quality of microhole machining via the single/multistep spiral drilling methods. The surface morphologies and taper of the microhole processed using the multistep spiral drilling method are obviously improved (Fig.12 and Table 2). This is because a prefabricated hole at the center of the microhole can discharge debris and water from the bottom of the hole, reduce the interference with laser transmission, and improve the surface morphologies and taper of the microhole.

The variations in the laser single-pulse energy, scanning speed, feed time, and scanning time on the surface morphologies and taper of microholes using the single spiral drilling method are investigated. A multistep spiral scanning mode drilling method is proposed to improve the defects of poor microhole surface morphologies and large tapers caused by the single-step spiral scanning mode. The quality of microhole machining using the single/multistep spiral drilling methods is compared under appropriate processing parameters. The experiments indicate that increasing the single-pulse energy and reducing the scanning speed can improve the surface morphology of microholes and reduce the microhole taper. With an increase in the feed times, the surface morphology of the microhole gradually improves, and the microhole taper initially decreases and then saturates. As the number of scanning rounds increases, the surface morphology of microhole gradually deteriorates, and the microhole taper first increases and then decreases. The taper of microholes processed using the multistep spiral method is only 0.29°, which is a 70% reduction compared to that using the single-step spiral method, and the dimensional deviation and roundness are controlled within 20 μm.

Surface-enhanced Raman scattering (SERS) plays an important role in trace detection and other fields of research. The use of periodic micro-nano structures has been a popular method for realizing high-performance SERS substrates. Traditional preparation technologies can prepare micro-nano structures with high precision and excellent quality. However, most of them have the limitations of strict environmental requirements, low efficiency, and high material dependence; therefore, it is necessary to develop alternative micro-nano periodic structure preparation techniques. Laser-induced forward transfer (LIFT) has been used to transfer almost all types of materials; however, it is limited in preparing periodic micro-nano structures. Laser interference lithography (LIL) has the advantage of high efficiency in preparing periodic micro–nano structures with large areas. Therefore, combining LIFT with LIL technology, a novel technique of laser interference-induced forward transfer (LIIFT) is proposed to transfer the periodic Ag micro-dots under three-beam laser interference in this study, which can overcome the shortcomings of traditional technologies and realize large-area metal micro-nano array structure manufacturing in a rapid, low-cost manner. Finally, the SERS properties of the Ag micro-dots transferred via LIIFT are tested and analyzed.

The donor film to be transferred consists of a quartz substrate, a polyimide(PI)sacrificial layer film mixed with carbon nanoparticles (CNPs@PI), and an Ag film. The preparation process of the CNPs@PI thin film is described in Ref.[16], and the Ag film is evaporated onto the CNPs@PI film (Fig.1). The principle of transferring a micro- dot via LIIFT is shown in Fig.2. The azimuth angles (0°, 120°, and 240°) and transverse-magnetic (TM) mode of the polarization direction are chosen. Moreover, the effect of the interference period on the transferred microstructure is discussed. Finally, the Raman spectra of RhB on the transferred Ag micro-dots with different periods (9, 11, and 15 μm) are tested to analyze the SERS property of the transferred Ag micro-dots.

The transferred Ag micro-dot structure is observed on the receiving substrate, as shown in Fig. 3(a1)?(a3). While the interference periods of the three beams are 9 μm, the outline of the micro-dot structure can be observed, but there are many overlapping areas between adjacent micro-dots, as shown in Fig.3(a1). While the interference period increases to 11 μm, a micro-dot with clear edge can be observed, as shown in Fig.3(a2). Upon increasing the period to 15 μm, the clarity of the micro-dot is further improved [Fig.3(a3)]. These results, combined with Fig. 3(b1)?(b3), indicate that an increasing number of nanoparticles are distributed in the middle regions of a single micro-dot, leading to an improved resolution of each micro-dot with an increase in the period. The average diameter of the nanoparticles in the consistent area in the middle regions of the micro-dots with different periods is 129?141 nm, as shown in Fig.3(c1)?(c3). These results, combined with the statistical results of the numbers of Ag nanoparticles in the consistent area (Fig.4), indicate that, with an increase in the interference period, the average size of the Ag nanoparticles changes slightly, whereas the number of nanoparticles increases significantly. This is because the maximum intensity and contrast of the three-beam interference light field increase with the laser interference period (Fig.5), which affects the transferable mass of the Ag film and the distribution of Ag nanoparticles on the receiving substrate.

The SERS characteristics of Ag micro-dot substrates with periods of 9, 11, and 15 μm are tested by selecting Rhodamine B(RhB) as the analyte with a concentration of 10^-3 mol?L^-1 and compared with the RhB Raman spectrum on the bare Si substrate. The results show that, while using an Si substrate covered with the transferred Ag nanoparticle micro-dots, the Raman intensity of RhB is significantly enhanced compared with that of the bare Si substrate. Simultaneously, the Raman intensity of RhB on the Ag nanoparticle micro-dot substrate increases with the micro-dot period (Fig.7). This can be attributed to the local surface plasmon resonance effect. With the increase in interference period, the density of the transferred Ag nanoparticles increases evidently, which can create more and stronger “hot spots,” making it easier to excite strong Raman signals.

Periodic Ag nanoparticle micro-dots are realized efficiently with three-beam laser interference LIIFT. The distribution of Ag nanoparticles in each micro-dot is controlled by adjusting the laser interference period. With the increase in the period of the laser interference from 9 μm to 15 μm, the density of Ag nanoparticles in the micro-dot increases. Finally, the Ag nanoparticle micro-dot substrates prepared by the three-beam LIIFT technology show significant SERS characteristics. Moreover, the Raman intensity of RhB on the transferred Ag micro-dots increases with the period of the micro-dots. This study demonstrates the potential application of LIIFT technology in the field of SERS chips.

High manganese steel (HMnS) has good deformation hardening properties. Under impact load, its surface rapidly hardens, thereby improving surface hardness but maintaining good toughness inside. However, under low stress wear conditions, it often exhibits a low hardening behavior accompanied by early surface wear. The pre-hardening treatment of the HMnS surface can improve its mechanical properties under low impact and low stress conditions. Therefore, scholars have proposed various surface pre-hardening treatment methods for HMnS, such as mechanical shot peening, explosive hardening, laser shock, and ultrasonic shock. Laser remelting is the process of using a laser beam to melt the surface of a material and improve its microstructure and mechanical properties through the rapid solidification of the molten pool. Unlike the equiaxed grains of cast HMnS, laser remelted HMnS often forms typical columnar and equiaxed dendritic structures due to the high temperature gradient and high cooling rate during solidification. Therefore, exploring the hardening behavior and wear resistance of laser remelted HMnS under ultrasonic rolling is of great significance.

This study used continuous cast Mn13 steel plate for laser remelting, and its cross-sectional microstructure was equiaxed grains. The laser power used was 700 W, the laser wavelength was about 960 nm, the scanning speed was 5 mm/s, the spot diameter was 1.2 mm, and the overlap rate was 50%. To prevent oxidation during laser remelting, high-purity argon with volume fraction of 99.99% gas was selected as the protective gas. An ultrasonic rolling strengthening device was used to treat the surface of HMnS after laser remelting with an amplitude of 4 μm. The vibration frequency was 40 kHz, and the static pressures were 100 N and 200 N, respectively. The samples were sequentially ground, polished, and corroded using silicon carbide sandpaper, metallographic grinder, and aqua regia solution. Measurement and analysis of laser remelted HMnS before and after ultrasonic rolling were carried out using field emission electron probe microanalyser, electron backscatter diffractometer, field emission scanning electron microscope, roughness profilometer, Vickers hardness tester, pin disc rotary friction and wear tester, and three-dimensional profilometer.

During the laser remelting, due to the high cooling rate and temperature gradient, the solidification structure consists of columnar and equiaxed dendrites, without obvious defects such as cracks and pores and without precipitation of cementite. After ultrasonic rolling with static pressures of 100 N and 200 N, the surface hardness increases by 120.48% and 173.82%, respectively. It can be seen that the microstructure of laser remelted HMnS also has deformation hardening, especially with outstanding surface hardness properties. The wear test shows that without ultrasonic rolling, the depth and width of the wear marks are the highest. In contrast, the depth and width of the wear marks are the lowest when the static pressure of ultrasonic rolling is 100 N. The volume wear rate without ultrasonic rolling is 6.945×10^-5 mm³/(N·m), and those under ultrasonic rolling with static pressure of 100 N and 200 N are 4.93×10^-5 mm³/(N·m) and 5.95×10^-5 mm³/(N·m), respectively. The ultrasonic rolling hardening mechanism of laser remelted HMnS is as follows. During the ultrasonic rolling, the surface of laser remelted HMnS undergoes severe plastic deformation, which is essentially dislocation slip and deformation twinning. Normally, high-frequency ultrasonic rolling can obtain nanograins on the surface of the material. Unlike the equiaxed grain structure of cast HMnS, laser remelted HMnS has a high interdendritic Mn content, while the intra-dendrite Mn content is lower. So the stacking fault energy within the dendrites is lower than that between the dendrites, making it easier to form twins within the dendrites. Twins can still expand between adjacent dendrites, forming twins that can penetrate multiple dendrites. The results indicate that the small angle grain boundaries and Mn segregation do not inhibit the formation and expansion of twinning. Due to the small angle grain boundaries and Mn segregation that can hinder the movement of dislocations, laser remelted HMnS exhibits good deformation hardening behavior.

This study uses laser remelting technology to obtain non-uniform solidification structure on the surface of HMnS, and investigates the hardening behavior and wear resistance of non-uniform solidification structure of HMnS under ultrasonic rolling. The solidification structure of laser remelted HMnS is composed of thinner equiaxed dendrites and columnar dendrites growing perpendicular to the bonding surface. There are many small angle grain boundaries formed in the solidification structure, and there is Mn segregation at the small angle grain boundaries. The non-uniform structure of laser remelted HMnS forms a dense twinning and thin severe plastic deformation layer under ultrasonic rolling, indicating its twinning hardening behavior. The thickness of the severe plastic deformation layer is 3?4 μm. The wear test shows that when the static pressure of ultrasonic rolling is 100 N, the twinning hardening and severe plastic deformation of the surface significantly increase the surface hardness of HMnS, and the volume wear rate is reduced by 29.01% compared to that of the surface without ultrasonic rolling. The wear mechanism is light adhesive wear and abrasive wear.

AISI 1045 steel (45 steel) has good plasticity, ductility, and excellent mechanical properties and is widely used in automotive manufacturing. However, the surface of the material can be damaged by friction, leading to a deterioration in the performance of components in contact with moving parts. Further, damage caused by corrosion and abrasion accelerates the expansion of cracks and risks fracture. Partial remanufacturing is an effective way to reservice damaged parts, maximizing the residual value of the material. Laser additive manufacturing is considered the most promising remanufacturing technology for rebuilding the geometric features of damaged parts and restoring their mechanical properties; however, it faces problems concerning material properties. In this study, an innovative combination of laser directed energy deposition (LDED) and laser shock processing (LSP) processes is proposed for the remanufacture of damaged 45 steel, utilizing the respective advantages of each process.

Experiments were conducted on 45 steel, whose chemical composition is displayed in Table 1. H13 tool-steel powder was used as the laser deposition powder for the experiments, and its chemical composition is displayed in Table 2. The laser composite remanufacturing process was realized using LDED and LSP equipment, and the specimens were fabricated according to LDED and LSP experimental parameters: a 2 mm thick layer was deposited on the substrate using LDED, milled to a smooth surface, and treated with LSP. Finally, the specimens were cleaned using ultrasonic vibration. A dry slip abrasion test was carried out on an HT-1000 spherical disc high-temperature tribometer based on the ASTM standard G99-95.

The number of small pores around the contact area significantly reduced with increased laser power. Comparisons of samples before and after LSP show that the strained areas exhibit inhomogeneous surfaces (Fig.2). With increasing laser power, the microhardness gradually increases; LSP significantly improves the microhardness of the LDED repair layer (Fig.4). The LDED-1200 W specimen has broad martensitic laths with a small number of fine needles; in comparison, in the LDED-1800 W specimen, the lath size decreases, the grain boundaries increase significantly, and internal refinement occurs with some dislocation (Fig.7). LSP induces significant refinement of surface grains, forming tiny nanoparticles with non-sequential orientation; the impact extends downward along the depth, and a large number of discrete dislocation structures, including dislocation tangles and cells, were found near the impact surface (Fig.8). The LDED specimen has a large worn area, with deep grooves and ridges parallel to the sliding direction, and almost the entire worn area is severely abraded (Fig.11); conversely, the worn surface of the LDED+LSP specimen is smooth, with only a small number of scratches and grooves in the middle of the wear, no large worn area, and no obvious adhesion phenomenon on the surface (Fig.12).

A combination of LDED and LSP post-treatment was used to repair damaged 45 steel. The main conclusions are as follows:

(1) With increased LDED laser power, the powder is fully dissolved under high heat, the forming quality of the repair layer is improved, the quantity of internal holes is reduced, the porosity is reduced, the martensite lath-like structure is refined, and the cementite in the structure is dissolved.

(2) Plastic deformation of the material occurs under the influence of LSP and the surface grains undergo refinement, forming nanograins with an approximate size of 30?50 nm. Subsequently, the deformation influence extends along the depth, generating many dislocations and forming high-density dislocation structures.

(3) The main wears on the LDED-restored layer are plowing and adhesive wears, with a small contribution from abrasive wear; whereas the main wear mechanism of the LDED+LSP-restored layer is adhesive wear accompanied by abrasive wear.

(4) LSP induces nanograin and dislocation reinforcement to refine the material structure, which effectively eliminated the internal pores, compacted the structure, and realized surface hardening, thus improving wear resistance. Meanwhile, the post-treatment process is accompanied by the dissolution of primary cementite, which further improves wear resistance.

Laser cladding technology offers an effective repair solution for turbine blades, aircraft landing gears, propellers, and large gears. Typically, these workpieces are fixed to the worktable, after which surface repairs are carried out using the laser cladding head. However, gravitational effects, complex material flow, vaporization, and thermo-physical processes within the melting pool pose challenges in achieving desired shapes and properties, affecting the morphology of the cladding layer and the overall surface quality. Laser remelting, which is a surface post-treatment method, can improve the performance and morphology of a surface without changing the original equipment, thereby saving post-treatment time. This study aims to elucidate the modifications in the laser cladding layer morphology at various tilting angles. Through finite element simulations, we simulate the impacts of substrate tilting angles on the resulting morphology of the laser cladding layer. Furthermore, our investigation aims to identify the optimal location for in-situ remelting by analyzing the distribution of the molten pool velocity and temperature fields. This study involves conducting multi-track experiments to compare the macro-morphology of the cladding layer before and after remelting. We aspire for our findings to serve as a valuable reference guiding the understanding of how in-situ remelting influences morphology modification and microstructure evolution in laser cladding layer on inclined substrates.

This study employed 45 steel substrates and 316L powder materials. A finite element simulation software was used to numerically simulate the laser cladding and in-situ remelting process on inclined substrates. The laser cladding simulation was carried out at different tilting angles, and a tilting angle of 30° was selected for the in-situ remelting simulation. The molten pool velocity and temperature distribution across the three remelting positions were simulated and compared. Subsequently, practical in-situ remelting experiments were performed on tilted laser cladding layers. The changes in the surface morphology of the molten cladding layer before and after remelting were examined using an optical microscope. Mechanical properties were assessed using Vickers hardness testing, friction and wear testing, and three-dimensional profilometry. These analyses were aimed at comprehending the microstructural transformations induced by in-situ remelting.

The morphology of the laser cladding layer on the inclined substrate is significantly affected by gravity. After in-situ remelting, the molten pool tends to flow along the direction of gravity. When the substrate tilt angle is less than 90°, the leading angle is inversely proportional to the substrate tilt angle, whereas the trailing angle is directly proportional. However, when the substrate tilt angle exceeds 90°, the morphology of the molten pool becomes highly unstable. The regular downward-sliding morphology gradually transforms into unpredictable irregular patterns, resulting in splattering and droplet formation (Fig. 4). Numerical simulations of the molten pool velocity field and temperature field during the in-situ laser remelting of the tilted fused cladding show that the optimal remelting position is in the middle of the overlapping region of the cladding tracks (Fig. 6). The primary factor behind this phenomenon is that during remelting, not all the cladding layers on both sides of the laser are fully melted. As the cladding layers on either side of the laser melt, they flow toward the center of the cladding structure. However, the flow of the molten pool encounters resistance from gravity and the existing cladding layers. Consequently, the molten pool bulges upward. Additionally, because of the impact of the laser and powder, a portion of the velocity field forms leftward vortices, whereas another portion of the velocity field increases the height of the molten pool. After remelting, the central region between the two cladding layers exhibits spherical bulges. This effect occurs primarily because a section of the cladding layer is remelted. As the grooves fill with the flowing molten pool, any excess material moves toward the surface of the cladding layer owing to gravitational forces. Given the distance from the center of the heat source, there is an increase in the temperature gradient, accompanied by vaporization of the excess solute. This vaporization lifts the excess material, forming a mass that contracts and solidifies into uniformly sized liquid beads within the molten pool (Fig. 8). The structural changes at the bottom of the cladding layer are minimal compared with the upper portion, where the laser remelting significantly impacts the structure. Consequently, the heat-affected zone from laser remelting primarily concentrates in the upper-middle region of the cladding layer. Observations reveal that the post-remelting dendritic structures at the top of the cladding layer almost vanish. A substantial proportion of the dendritic structures fracture, transforming into equiaxed grains. Furthermore, the grains at the top undergo noticeable refinement (Fig. 9). After laser remelting, an intensified temperature gradient occurs at the molten pool interface, accelerating the crystallization rate. Elevated cooling rates impede grain growth while promoting heightened nucleation, thereby refining the grain structure. Consequently, in-situ laser remelting densifies the cladding layer structure via these mechanisms.

Gravity has a significant impact on the morphology of the cladding layer, and as a result, the cladding layer tends to flow along the direction of gravity after melting. When the tilt angle of the substrate is greater than 90°, the molten pool morphology is extremely unstable. It transitions from the original downward-dripping regular morphology to an unpredictable irregular morphology, often resulting in splashing or even dripping, which substantially damages the morphology of the cladding layer. In-situ remelting in the overlap zone of the cladding channel brings about the most substantial improvement in the surface flatness of the cladding layer. The surface unevenness is reduced significantly, from 0.165 to 0.056, because of remelting. Our results show that the in-situ laser remelting primarily affects the microstructure in the middle and upper parts of the cladding layer. The dendritic crystals at the top of the layer nearly vanish, and a large number of dendritic crystals transform into isometric crystals. After remelting, the height of the cladding layers reduces by 16.1% compared to their pre-melting state, and the roughness significantly decreases by 69.5%. The surface hardness of the cladding layers increases by 70 HV after remelting, accompanied by a 76.3% reduction in wear rate.

After an aircraft is in service for a certain period, various factors, such as external forces, light, and humidity, affect the paint layer on its skin surface. This leads to aging, cracking, peeling, and other phenomena. Hence, removing the original coating on the metal material surface becomes a necessary step before repainting it. Laser cleaning offers advantages, including non-contact, environmental friendliness, precision, and no secondary pollution. It can replace traditional mechanical polishing and chemical paint removal methods, enhancing the cleanliness of aircraft surface paint layers. Due to the strict safety requirements of aircraft, it is crucial to understand whether the use of lasers for removing surface paint layers impacts the fatigue properties of the aircraft. In this study, a nanosecond pulse laser is used to clean aviation aluminum alloys coated with a fluid-resistant epoxy primer coating. Subsequently, the effects of the laser cleaning process on the microstructure and fatigue properties of the aluminum alloys are examined.

A pulsed laser was used to remove the surface coating of the aviation aluminum alloys. The effects of the laser cleaning process on the microstructure, mechanical properties, and fatigue properties of the aviation aluminum alloys were examined via appearance inspection, optical microscope (OM), scanning electron microscope (SEM), in situ temperature detection, mechanical property detection, and high cycle fatigue test.

The results show that when the laser power and pulse frequency are 80 W and 100 kHz, the epoxy primer coating on the surface of the aluminum alloy can be removed. However, some residual paint layer remains on the surface of the sample. When the laser power and pulse frequency are 500 W and 500 kHz, the maximum surface temperature does not exceed 115 ℃, and no obvious heat affected zone is observed. Laser cleaning increases the surface roughness, with partial ablation and melting occurring within the depth range of 10 μm. After laser cleaning, the hardness of the material increases. As the laser power, frequency, and energy density increase, the rate of hardness increase decelerates. The tensile property results indicate that the tensile strength, yield strength, and elongation of the sample after laser cleaning are slightly lower than those of the blank sample. Through high-cycle fatigue testing, when compared to those of the blank sample, the fatigue properties of the painted sample after laser cleaning decrease by 11.76%. This mainly stems from the increased surface roughness caused by laser cleaning. However, after anodizing and painting treatments, laser cleaning does not further exacerbate the fatigue damage caused by anodizing.

Analyses are conducted on the appearance, microstructure, roughness, in-situ temperature, mechanical properties, and fatigue performance of the samples after laser cleaning. With a laser power and pulse frequency of 80 W and 100 kHz, a residual paint layer remains on the sample surface. However, at elevated levels of 500 W and 500 kHz, oxidation might appear on the substrate surface. The process of laser cleaning tends to increase surface roughness, causing partial ablation and melting within the depth range of 10 μm. The surface temperature during this procedure increases in tandem with the increase in laser power and pulse frequency. But even at peaks of 500 W and 500 kHz, the maximum surface temperature stays below 115 ℃. After cleaning, the material hardness increases. However, as the laser power, frequency, and energy density increase, the increase in hardness decelerates. There is a minor reduction in the sample tensile strength, yield strength, and elongation. When compared to the untreated samples, those cleaned by laser but not anodized or painted show a reduction in fatigue properties by 9.34%. In comparison, samples that undergo anodizing and painting processes after cleaning experience a reduction in fatigue properties by 13.84%. Specifically, after painting, laser cleaning results in a decrease in fatigue properties by 11.76%. Notably, laser cleaning does not further increase the fatigue damage due to anodizing.

Civil aircrafts operate under severe conditions, leading to issues such as peeling and cracking in the aircraft paint layer. This makes localized or comprehensive paint removal and surface maintenance pivotal during C-check or D-check procedures. Current paint removal methods, such as manual grinding and chemical stripping, are widely used. However, they have significant shortcomings. Manual grinding often lacks precision and can damage the aircraft surface, while chemical stripping is complicated and environmentally detrimental. These challenges hinder sustainable and efficient advancements in the civil aviation industry. Laser paint removal has emerged as a promising solution, offering high precision, reduced pollution, and automation possibilities. It is rapidly gaining global attention. However, a knowledge gap exists in understanding the exact mechanism of paint removal during laser ablation, particularly regarding the impacts of single-pulse thermal stress and plasma-induced removal. This study bridges this gap by first determining the vaporization point and strength limit of the paint layer. Then, single-pulse laser ablation simulations are compared with experimental results to better understand thermal stress paint removal during laser cleaning. This research sheds light on paint removal mechanisms and evaluates the impacts of varying scanning speeds on laser paint removal efficiency. Hence, the aim is to offer valuable insights and references for enhancing the use and development of laser paint removal techniques on civil aircraft skin.

In this study, the vaporization point and strength limit of the paint layer are first determined via thermogravimetry and stress-strain experiments. Subsequently, a finite element analysis of the single-pulse laser ablation-thermal stress paint removal process is conducted using the COMSOL software. Experiments on laser paint removal, both single-pulse and multi-pulse, at varying scanning speeds (ν=1000, 900, 800, and 700 mm/s), are performed on the composite paint system of the LY12 aluminum alloy substrate using nanosecond pulsed fiber lasers. Then, a white light interferometer (WLI) is employed to inspect and analyze the crater profile created by the single-pulse laser. An optical microscope (OM) is used to characterize the resulting surface and cross-sectional morphology from the multi-pulse line scanning laser paint removal. Additionally, a scanning electron microscope (SEM) and an energy-dispersive X-ray spectroscope (EDS) are utilized to analyze the microstructure and compositional changes on the cleaned surface after the paint is removed.

After the completion of single-pulse laser irradiation (t=200 ns), the paint ablation process does not cease. The accumulated heat causes the surface temperature of the residual paint layer to further increase until t=750 ns. When the temperature falls below its vaporization point, the ablation paint removal process concludes (Figs. 6 and 7). At this moment, the action depth of thermal stress generated by the coupling of the temperature field on the surface of the residual paint layer gradually increases from the bottom to the edge of the crater, while the corresponding values progressively decrease (Fig. 9). The maximum value (Fig. 9, σ_max=2.7×10⁷ Pa) approaches the strength limit of the paint layer [Fig. 11(a), σ=2.68×10⁷ Pa], leading to a U-shaped stress-damage zone on the surface of the residual paint layer (Fig. 12). This results in physical damage, such as delamination and fragmentation of the residual paint layer (Fig. 13). During the laser ablation process, both plasma impact and shielding effects coexist. The plasma impact causes the actual width of the ablated crater to be larger than the simulated results, while the shielding effect results in the experimental ablation depth of the crater being smaller than the simulated results (Fig. 14). In the process of single-pulse laser paint removal, the ablation and plasma paint removal effects gradually intensify before the laser irradiation ends and then gradually weaken after the irradiation completes. However, the thermal stress paint removal effect remains unchanged (Fig. 15). In the multi-pulse surface scanning laser paint removal experiments, the actual spot overlap rate is higher than the theoretical value (Fig. 17), resulting in the removal depth of the paint layer being greater than that of the single-pulse results. As the scanning speed gradually reduces, the removal effect of the paint layer gradually enhances due to ablation, the plasma effect gradually weakens, and the depth of paint removal and the deposition amount of β-type copper phthalocyanine along with the functional oxidized particles gradually increase [Figs. 16(a), (c), (e), and (g), and Table 10].

In this study, the single-pulse nanosecond laser thermo-mechanical coupling paint removal transient process is simulated using the COMSOL software. The results show that the behavior of ablative removal of paint layers does not finish at the end of one pulse cycle, and the heat accumulated on the cleaning surface extends the ablation process by 550 ns. The simulated thermal stress value is slightly larger than the experimentally determined tensile limit of the paint layer. This causes the residual paint layer on the cleaning surface to produce a delamination and cracking zone ranging from 0.6 μm to 2.8 μm. The laser ablation paint removal process triggers plasma impact and shielding. This makes the experimental crater paint removal width larger and the depth smaller than the simulation results. In the single pulse laser paint removal process, the ablation and plasma paint removal effects gradually intensify before the laser irradiation ends and then gradually weaken after the irradiation ends, but the thermal stress paint removal effect remains unchanged. As the scanning speed gradually decreases, the ablative paint removal effect strengthens, the plasma effect weakens, the thickness of the paint layer removal increases, and the amount of functionally oxidized particles deposited on the paint layer also increases. When ν=1000 mm/s, the topcoat is partially removed and the primer is slightly damaged. The topcoat is removed cleanly, and the primer is partially removed when ν=900 mm/s. However, for ν=800 mm/s and 700 mm/s, the area and depth of the residual primer continue to decrease with the reduction in the scanning speed, and the oxidized film is exposed.

Diamond is a wide-bandgap semiconductor material with several excellent physical and chemical properties. It has an ultrawide bandgap of 5.5 eV, which is significantly higher than those of GaN, SiC, and other wide-bandgap semiconductor materials. In addition, it has a low dielectric constant, low friction coefficient, high carrier mobility, high electron drift speed, and high thermal conductivity. These unique properties make diamonds have an important application value in optics and microelectronics. Because of the high hardness of diamonds, the traditional mechanical polishing method, which yields low polishing speeds and has high costs, cannot achieve an ideal effect. Ion-beam etching is a highly efficient noncontact surface-polishing method for super-hard and brittle substrates. However, it is unsuitable for industrial production because of its high cost. Laser processing is a noncontact processing technique that can handle curved surfaces. It has high processing efficiency and can achieve high-quality processing of various hard materials. Therefore, laser polishing can be used to polish the diamond film using the high energy of the laser to ablate the edges of the diamond particles. It can reduce the surface roughness and flatten the film, but typically induces a surface microstructure or nanostructure on the film surface and introduces a graphite layer. Although several polishing methods for diamond films have been developed, they have limitations, and it is difficult to satisfy the increasing application requirements. To solve these problems, we propose a composite polishing method that uses laser polishing combined with ion-beam etching. By further optimizing the polishing process parameters, a diamond surface without a modified layer is obtained, and the roughness is reduced. The results of this study provide technical support for diamond micromachining and related microdevice preparation.

The research object of this study is a diamond film prepared via chemical vapor deposition (CVD). The CVD diamond film was first prepolished using a femtosecond laser. The incidence angle of the laser was varied, and the diamond surface was initially polished by controlling the femtosecond laser output power and exposure time. The three-dimensional (3D) surface morphology and roughness of the diamond films were characterized and analyzed using 3D laser microscopy. Next, the power parameters of the nanosecond laser were controlled, and fine polishing was performed. The effect of nanosecond laser machining on the surface roughness of the films was assessed. Subsequently, the effect of the ion-beam etching time on the roughness of the CVD diamond was analyzed. The morphology of the polished diamond films was observed using cold-field emission scanning electron microscopy. The Raman scattering spectra of the samples were measured using Raman spectrometry to analyze the changes in the graphite layers during different polishing processes.

After femtosecond+nanosecond machining and ion-beam etching, the roughness of diamond surface decreases significantly, from 4 μm without etching to 0.47 μm after etching. In addition, the graphite layer formed by the thermal effect during laser processing can be effectively removed, and the diamond surface can be polished without modification and with high smoothness.

1. Using the femtosecond+nanosecond polishing method to polish the surface of diamond film can effectively reduce the surface roughness and produce a smooth surface.

2. Laser-polished diamond is typically converted into graphite because of the thermal effect that accumulates on the diamond film surface, which ablates the film surface and forms a graphite layer on the surface. By bombarding the laser-polished surface structure with an ion-beam, graphitization can be effectively eliminated, and an unmodified layer can be formed.

3. The use of field mirrors to polish diamond films can result in efficient large-area processing. With shortened scanning and ion-beam etching time, rapid preparation can be achieved, creating conditions for the industrial application of diamond films.

In this study, the effect of femtosecond+nanosecond+ion-beam polishing on the roughness of CVD diamond films was investigated. Ideal surface roughness can be achieved by selecting suitable laser processing and ion-beam polishing parameters. The composite polishing technology of laser polishing and ion-beam etching can effectively polish CVD diamond films. By controlling the femtosecond laser output power and exposure time and varying the laser incident angle, rough polishing of the diamond surface can reduce the roughness and formation of the graphite layer. Nanosecond processing can be fine-processed after applying femtosecond rough processing, however, owing to the thermal effect, a graphite layer is formed during processing. Finally, the graphite layer is effectively removed via ion-beam etching. High-quality polishing is achieved without modifying the layers. Compared with the roughness of approximately 4 μm before polishing, the surface roughness of the composite polished diamond film decreases significantly, with the minimum value reaching 0.47 μm. The proposed method polishes diamond surfaces and provides support for the micromachining and fabrication of micro-optical components on diamond surfaces.