Inconel 718 alloy, renowned for its outstanding high-temperature strength, oxidation resistance, and corrosion resistance, has found extensive applications across aerospace, energy, and other critical sectors. The advent of the additive manufacturing technology, particularly the laser powder bed fusion technology, has opened up novel avenues for the fabrication of Inconel 718 alloy and its composites. Nevertheless, the intricate microstructure of components produced by this technology gives rise to selective dissolution during subsequent electrochemical machining. This phenomenon not only deteriorates the surface quality but also restricts the further expansion of its applications.

In this study, laser powder bed fused Inconel 718 alloy and TiB?/Inconel 718 composite are used to innovatively address the surface quality problem faced by traditional electrolytic grinding. The core purpose of this research is to deeply explore the intrinsic relationship between the microscopic properties of materials and the processing technology, so as to optimize the processing quality and expand the application range of materials. Specifically, this research uses X-ray photoelectron spectroscopy (XPS) analysis technology to determine the chemical composition and elemental distribution of the superpassivation films of two materials. Through a detailed analysis of the composition of the superpassivation film, the aim is first to reveal its formation mechanism and the influence on the surface properties of the material during processing, and then to study the parameters of the electrolytic grinding process including different feed rates and rotational speeds and the relationship between the microstructure of materials and the electrolytic grinding surface quality . Scanning electron microscope (SEM) is used to observe the microstructural changes of materials under different process parameters in real time, and the surface quality parameters after electrolytic grinding can be accurately determined in combination with surface roughness measuring instruments and other equipment. Through comprehensive and detailed experimental data collection and analysis, the influence mechanism of microstructure on surface quality under different process conditions is deeply explored.

Compared to Inconel 718 alloy, TiB?/Inconel 718 composite can form a denser superpassivation film during electrolytic machining. The root cause of this phenomenon is that the addition of TiB₂ induces the formation of more homogeneous microstructures inside the material. At the microscopic level, the uniformly distributed TiB₂ particles act as a “backbone” and provide numerous stable attachment sites for the growth of superpassivation films. These sites make the superpassivation film more uniform in the formation process, and due to its strong interaction with the TiB₂ particles, the superpassivation film can be firmly attached to the TiB? particles, effectively avoiding being washed away by the electrolyte, thus ensuring the integrity and stability of the superpassivation film and improving the corrosion resistance of the material. Both materials are composed of Ni(OH)?, Fe(OOH), Fe?O?, Cr(OH)?, Cr?O?, TiO?, MoO?, Nb?O?, and other compounds. It is worth noting that the added TiB? does not participate in electrochemical dissolution throughout the electrochemical processing process, and mainly plays a role in enhancing the microstructural stability and promoting the formation of superpassivation films in the material. When it comes to the study of the effect of process parameters on surface quality, both Inconel 718 alloy and TiB?/Inconel 718 composite are available at a feed rate of 1.33 mm/s and a rotational speed of 1000 r/min, the surface quality is at its best, the surface is extremely flat, and there are no pits. At the same time, the gouge phenomenon caused by stray corrosion has been significantly improved, and the gouge amount has been greatly reduced.

Compared with Inconel 718 alloy, the nano-TiB₂ particles in the microstructure of TiB₂/Inconel 718 composite are diffusely distributed in the dendrite, which forms a denser superpassivation film in electrolytic machining, and the optimal parameters in electrolytic grinding are 1.33 mm/s and 1000 r/min, in which the volumetric electrochemical equivalent of Inconel 718 alloy is 1.727 mm³/(A?min), and the volumetric electrochemical equivalent of TiB₂/ Inconel 718 composite is 1.796 mm³/(A?min).

Bone tissue engineering is a crucial approach for treating bone defects. NiTi alloys are regarded as highly promising materials for bone tissue engineering due to their exceptional superelasticity, shape memory effect, high specific strength, excellent corrosion resistance, and biocompatibility. However, their further development is constrained by traditional fabrication methods. Laser powder bed fusion (LPBF) technology can precisely control the porosity, phase transition temperature, and comprehensive properties of NiTi alloys, providing a new approach for preparing NiTi alloy bone scaffolds. Based on the application requirements of bone implants, this study systematically investigates the phase transformation temperature, mechanical properties, and superelastic behavior of LPBF-fabricated NiTi alloys, aiming to provide a theoretical foundation for the development of LPBF-fabricated NiTi alloy bone implants.

In this study, near-equiatomic NiTi alloy powder (Ni atomic fraction of 50.8%) is fabricated using an LPBF system. During the printing process, a process window is established by varying the laser power (P of 190?310 W) and scan speed (v of 900?1300 mm/s) to produce NiTi alloy samples with dimensions of 8 mm×8 mm×8 mm. The density of the samples is measured using the Archimedes drainage method, and the phase transformation temperatures are determined using differential scanning calorimetry (DSC), with the results visualized in a contour map. Additionally, the phase composition and microstructural evolution are analyzed using X-ray diffraction (XRD) and electron backscatter diffraction (EBSD). The mechanical properties and superelastic behavior are evaluated through uniaxial tensile tests and cyclic loading–unloading experiments.

An analysis of the contour maps of phase transformation temperature and relative density shows that under a constant laser power of 250 W, the phase transformation temperature of the samples is slightly higher than body temperature (37 ℃). Additionally, under these conditions, the relative density exceeds 99.5%, meeting the requirements for both phase transformation temperature and relative density (Fig. 2). When the laser power is kept constant at 250 W, with increased scanning speed, the phase transition temperature first increases and then decreases (Fig. 3). Three groups of representative samples (samples 1#, 2#, and 3#, with scanning speeds of 900, 1100, and 1300 mm/s, respectively) are selected for subsequent research. As the scanning speed increases, the grain size of the columnar crystals gradually decreases, accompanied by an increase in grain boundaries. Consequently, the average value of the Kernel average misorientation (KAM) progressively increases, indicating higher dislocation density and more pronounced changes in crystal orientation (Fig. 5). As the scanning speed increases, the cooling rate accelerates, temperature gradient decreases, and interface velocity increases, promoting the formation of finer grains. Sample 2# exhibits the highest tensile strength (625.6 MPa) and yield strength (336.8 MPa), demonstrating the best overall mechanical performance. However, excessively high scanning speeds lead to the formation of a small amount of unmelted powder and pores during the melting process. These defects act as stress concentration points during tensile testing, initiating and propagating cracks at lower stress levels and ultimately causing premature fracture, thereby reducing both the tensile strength and elongation at break of the material (Figs. 6 and 7). During the cyclic tensile loading–unloading tests, after the first cycle, sample 1# exhibits the highest recovery strain (10.09%), while sample 2# shows the lowest total strain and recovery strain compared to the other two samples. However, sample 2# demonstrates the highest deformation recovery rate of 94.02%. By the 10th cycle, samples 2# and 3# exhibited the same deformation recovery rate (99.51%). Considering the overall cyclic recovery performance, sample 2# exhibits the best deformation recovery behavior and optimal superelasticity (Figs. 8 and 9). Under the same stress level, sample 2# experiences less dislocation slip and martensitic plastic deformation, but a higher proportion of stress-induced martensitic transformation occurs. This results in a higher superelastic recovery rate for sample 2# compared to the other two samples.

This study investigates the influence of laser power and scanning speed on the relative density and phase transition temperature of an LPBF-fabricated NiTi alloy. A process window for LPBF of the NiTi alloy is established, and a method for optimizing the relationship between density and phase transition temperature is proposed. In addition, the impacts of different scanning speeds on the microstructure, mechanical properties, and superelasticity are analyzed. The findings indicate that at a laser power of 250 W and scanning speeds ranging from 900 mm/s to 1300 mm/s, the printed NiTi alloy exhibits a relative density higher than 99.5%, and its austenitic transformation finish temperature is approximately 39 ℃, close to human body temperature. As the scanning speed increases, the transition temperature and relative density of the alloy first increase and then decrease, while the grain size is gradually refined. The average grain size decreases from 34.4 μm to 20.5 μm. Both B2 austenite phase and B19' martensite phase are observed in the NiTi alloy formed by LPBF, but the peak of B19' martensite phase is weak. When the scanning speed is 1100 mm/s, the alloy exhibits optimal mechanical properties, with a tensile strength of 625.6 MPa and fracture strain of 14.67%. At a scanning speed of 900 mm/s, the sample exhibits the highest recovery strain (9.38%), while a scanning speed of 1100 mm/s results in a higher deformation recovery rate (99.51%), with the deformation recovery rate reaching 94.02% during the first cycle of tensile testing, demonstrating optimal superelastic performance.

Owing to the low melting point of Mg, it tends to evaporate and form metallic vapor during the laser powder bed fusion (LPBF) process, affecting the stability of the laser forming process, forming quality, and mechanical properties of as-built magnesium alloys. Ceramic reinforcements added to the magnesium alloys offer an effective approach to enhance the forming quality and mechanical properties of the magnesium alloys. However, there remains limited work on the LPBF-fabricated TiC/AZ91D magnesium matrix composites, and the effect of laser processing parameters on the forming quality and mechanical properties of these composites remains unclear. This study investigates the effect of processing parameters on the forming quality and mechanical properties of LPBF-fabricated 2%TiC/AZ91D (2% is mass fraction) magnesium matrix composites. The findings provide valuable insights for the development of high-densification, high-performance magnesium matrix composites via laser additive manufacturing.

TiC/AZ91D powder mixture was prepared by ball milling. The laser powder bed fusion equipment was used for the fabrication of TiC/AZ91D magnesium matrix composites using the powder mixture. The Archimedes method was employed to measure the relative density of the LPBF-fabricated samples. Surface roughness was measured using a laser spectral confocal microscope. Optical microscope (OM) was used to observe the distribution of defects and the morphology of melt pool of the LPBF-fabricated samples. Microstructural characterization was conducted using a scanning electron microscope (SEM) equipped with an energy dispersive spectrometer (EDS). The hardness was measured using a microhardness tester, and tensile testing was performed using a universal testing machine.

The surface roughness of the LPBF-fabricated TiC/AZ91D magnesium matrix composites decreases by ~41% as the laser energy density increases (Fig. 3). This improvement can be attributed to sufficient energy input at a higher laser energy density, enhancing powder melting and melt spreading and thus reducing the surface roughness. The relative density of the TiC/AZ91D magnesium matrix composites increases as the laser power increases and decreases as the scanning speed increases (Table 1). A maximum relative density of ~99.4% is achieved at a laser power of 140 W and scanning speeds of 200 mm/s and 300 mm/s. Under high energy input, relatively long existence time for melt pool and small cooling rate are conducive to the escape of residual gas and reduction of porosity. At a laser power of 140 W and scanning speed of 300 mm/s, only few spherical pores are observed, well correlating with the high relative density of the fabricated samples (Fig. 4). The depth-to-width ratio of melt pool increases with the laser power (Fig. 5), as high energy input at a high laser power can enlarge the melt pool. The microstructure of the LPBF-fabricated TiC/AZ91D magnesium matrix composites comprises white unmelted TiC particles, light gray network and cellular β-Mg₁₇Al₁₂ precipitates, and dark gray α-Mg matrix [Fig. 6(a)]. TiC particles with a high melting point, are partially melted and can serve as heterogeneous nucleation sites to refine grains of matrix and improve the mechanical properties of the material. In the LPBF process, the slow cooling rate allows more time for Al to diffuse into the matrix, resulting in large amounts of β-Mg₁₇Al₁₂ precipitated in the matrix. At a laser power of 140 W and scanning speeds of 200 mm/s and 300 mm/s, the hardness distribution of the LPBF-fabricated TiC/AZ91D magnesium matrix composites is relatively uniform (Fig. 7). The average microhardness decreases from (114.0±2.5) HV to (102.9±0.5) HV as the laser power increases, and increases from (101.4±0.3) HV to (104.1±0.5) HV as the scanning speed increases (Fig. 8). These changes in hardness are attributed to a reduction in residual stresses and grain refinement of the composites. The ultimate tensile strength of the LPBF-fabricated TiC/AZ91D magnesium matrix composites increases from ~296.5 MPa to ~335.5 MPa as the laser power increases, representing a 13% increase (Fig. 9). This enhancement is attributed to increased Orowan strengthening in the composites. The ultimate tensile strength of the samples increases with the scanning speed up to a peak of ~335.5 MPa at 300 mm/s. However, the elongation decreases with the increase in scanning speed, reaching a maximum of ~3.47% at 200 mm/s (Fig. 10). This is owing to the reduced energy input at higher scanning speeds, which increases the metallurgical defects that can act as crack initiation sites, resulting in a reduced elongation. Fracture analysis of the LPBF-fabricated samples reveals un-melted TiC particles and no cracks or apparent porosity observed at the surface (Fig. 11). The uniform distribution of β-Mg₁₇Al₁₂ phases effectively hinders the propagation of microcracks, enhancing the strength of the material and improving its overall mechanical properties.

In this work, the 2%TiC/AZ91D magnesium matrix composites are fabricated at different laser processing parameters. The effect of the laser power and scanning speed on the manufacturing quality and mechanical properties is investigated. The surface roughness of LPBF-fabricated TiC/AZ91D magnesium matrix composites reaches ~12.4 μm at optimal processing parameters, which is attributed to sufficient laser energy input, enhancing the melt flowability and spreading. At a laser power of 140 W and scanning speed of 300 mm/s, the energy input ensures complete powder melting, clear melt tracks, and stable melt pools. The LPBF-fabricated samples have few pores, achieving good forming quality with a relative density of ~99.4%. The microstructure of LPBF-fabricated TiC/AZ91D magnesium matrix composites is composed of unmelted TiC reinforcements, a network-like distribution of β-Mg₁₇Al₁₂ precipitates, and α-Mg matrix. The average microhardness of the LPBF-fabricated TiC/AZ91D magnesium matrix composites reaches (114.0±2.5) HV at optimal processing parameters. At a laser power of 140 W and scanning speed of 300 mm/s, the tensile strength and elongation of the magnesium matrix composites reach ~335.5 MPa and ~3.45%, respectively. This improvement is primarily owing to the Orowan strengthening caused by fine β-Mg₁₇Al₁₂ precipitates that can hinder the dislocation motion.

Laser powder bed fusion (LPBF) has broad application prospects in the manufacturing of high-performance and complex metal components. However, the traditional Gaussian laser beams have nonuniform energy distributions that lead to challenges such as poor molten-pool stability and high cooling rates during processing. These issues can result in defects such as spattering and porosity, which hinder the technology’s advancement. Beam shaping techniques can help redistribute laser energy and effectively address these distribution challenges. Research shows that flat-top laser beams possess unique advantages in controlling defect formation and microstructure development. However, the mechanisms through which these beam profiles affect stress development are not well understood. This study aims to investigate the effects of Gaussian and flat-top laser beams on temperature and stress fields during the LPBF process under the conditions of maintaining the same spot diameter and total energy input. The goal is to provide a theoretical foundation for optimizing process parameters and regulating stress distribution in LPBF.

This study employs the numerical simulation method to systematically examine the effects of Gaussian and flat-top laser beams on the LPBF forming process. A finite element model of multi-pass scanning was developed based on thermo-mechanical coupling theory, and the reliability of the model was verified through single-pass scanning experiments. In the simulation process, a comparative study of Gaussian and flat-top laser beams was conducted by controlling variables such as spot diameter and total energy input. Under the condition of scanning speed of 700 mm/s, we performed a comprehensive analysis of key thermal field parameters, including temperature peaks, temperature gradients, and cooling rates at both the center and edge regions of the laser spot for each laser type. The evolution of residual stress was also evaluated. In addition, the study analyzed the trends in temperature peaks and residual stress across various process parameter combinations, providing a theoretical basis for process optimization.

The distribution of laser energy significantly affects the spatial temperature distribution, thereby affecting temperature peaks, gradients, and cooling rates. At a scanning speed of 700 mm/s, the maximum temperature at the center of the flat-top laser beam is approximately 541.28 ℃ lower than that of the Gaussian laser beam, whereas the maximum temperature at the edge of the spot is approximately 127.95 ℃ higher. Variations in energy distribution also result in differences in temperature gradients and cooling rates. The Gaussian laser beam demonstrates higher temperature gradients and cooling rates at the center of the spot, whereas the flat-top laser beam exhibits higher values at the edge. Specifically, the maximum temperature gradients at the edges of the Gaussian and flat-top laser spot are 5.38×10⁷ ℃/m and 5.74×10⁷ ℃/m, respectively. By contrast, the maximum cooling rates at the center of the above spots are 8.66×10⁶ ℃/s and 5.82×10⁶ ℃/s (Figs. 7?8). Despite these variations, the trends in residual stress evolution show similar patterns (Fig. 10) and no significant differences in the final residual stress distribution are observed (Fig. 11). Further analysis indicates that the variations in temperature peaks are more pronounced for the Gaussian laser beam at the center of the spot, whereas the flat-top laser beam's temperature peaks are more sensitive to changes at the edge of the spot (Fig. 9). In addition, changes in laser power and scanning speed do not significantly affect the differences in residual stress between the two laser beam types (Fig. 12).

This study developed a finite element model for multi-pass scanning and performed a detailed analysis of temperature peaks, temperature gradients, cooling rates, and residual stress distributions at both the center and edge of the laser spot for Gaussian and flat-top laser beams under a scanning speed of 700 mm/s and identical total energy input. The study also explored how variations in laser power and scanning speed affect temperature peaks and residual stress. The findings reveal that the flat-top laser beam provides a more consistent temperature distribution within the spot. Specifically, the temperature peak at the center of the flat-top spot is approximately 541.28 ℃ lower than that of the Gaussian beam, whereas the temperature peak at the edge of the spot is approximately 127.95 ℃ higher. In addition, the Gaussian laser beam demonstrates greater temperature gradients and cooling rates in the center, whereas the flat-top beam shows higher values at the edge of the spot. The maximum temperature gradients are observed at the edge of the spots, measuring 5.38×10⁷ ℃/m and 5.74×10⁷ ℃/m for the Gaussian and flat-top beams, respectively. The maximum cooling rates occur at the center of the spots, recorded at 8.66×10⁶ ℃/s and 5.82×10⁶ ℃/s for the Gaussian and flat-top beams, respectively. The effects of variations in laser power and scanning speed on temperature peaks differ between the two laser beam types. At the center of the spot, the Gaussian laser beam is more significantly affected, whereas at the edge of the spot, the flat-top laser beam shows greater sensitivity to these changes. However, these variations do not significantly influence the differences in residual stress distributions between the two types of laser beams.

Titanium alloys exhibit excellent biocompatibility, corrosion resistance, and specific strength, making them ideal materials for medical implants and prostheses such as dental implants and joint replacements. However, titanium alloys application in severe environments is limited owing to their poor surface wear resistance. Nitriding is a common method to enhance the performance of titanium alloys. However, for complex structural parts, the traditional nitriding process is cumbersome and complex, and the thickness of the resulting nitrided layer makes it difficult to achieve the desired effect. To meet the demand for improving the wear resistance of TC4 alloy in the fields of aerospace, medical devices, and bio-implants, some scholars have proposed a new method of simultaneous nitriding in additive manufacturing. By combining the advantages of additive manufacturing and laser gas nitriding, high-performance titanium alloys with larger nitrided areas can be obtained while increasing the processing efficiency. At present, research on synchronous nitriding in the additive manufacturing process has achieved some results; however, there have been few studies on the effect of nitrogen volume fraction on the organization and properties of titanium alloys during the forming process of selective laser melting (SLM). As a result, simultaneous SLM nitriding is achieved through selective laser melting under atmospheres with different nitrogen volume fractions. This microstructure and properties of SLM-formed TC4 alloys with different nitrogen gas volume fractions are systematically analyzed ,and the role of nitrogen in the SLM process of TC4 alloy is explored, providing a research basis for further nitriding technology development.

The substrate is sanded and cleaned to remove surface impurities. Subsequently, high-purity nitrogen and argon are passed into the high-pressure gas supply proportioning system in different ratios for mixing, and the mixed gases are then passed into the atmosphere protection box to provide a processing environment with different nitrogen volume fractions for the subsequent selective laser melting. The morphology, organization, physical phase, and elements of SLM-formed specimens are analyzed using optical microscope (OM), scanning electron microscope (SEM), X-ray diffractometer (XRD) and energy dispersive spectrometer (EDS) to determine the effect of the nitrogen volume fraction on the microstructure of the formed specimens. Based on this, the nitrogen mechanism in the SLM process is discussed.

In this study, the effect of different nitrogen volume fractions on the microstructure of SLM-formed TC4 alloys is investigated. The experimental results show that when the nitrogen volume fraction is 25%, the microstructure does not show obvious changes, and the nitrogen element primarily exists in the form of a solid solution. With a further increase in nitrogen volume fraction, the titanium nitride organization appears and gradually increases, and its morphology is transformed into dendritic crystals (Fig. 4). Subsequently, the microhardness, wear resistance, and corrosion resistance of the formed specimens with different nitrogen volume fractions are analyzed, and it is discovered that the microhardness gradually increases with increasing nitrogen volume fraction (Fig. 8). However, its abrasion resistance does not increase with increasing nitrogen volume fraction, and its abrasion resistance under 100%N₂, 25%N₂, 50%N₂,75%N₂, and 0%N₂ is from the largest to the smallest (Fig. 9). The corrosion resistance is analyzed, and it is discovered that the corrosion current density of the specimens with nitrogen volume fractions of 25% and 50% is reduced by 29.4% and 56.1%, respectively, compared to that without nitrogen, indicating an improvement in the corrosion resistance. When the nitrogen volume fraction is increased further, the corrosion resistance begins to decrease (Fig. 11), owing to its higher hard phase content, which increases material brittleness and causes a large residual stress. In the process of electrochemical corrosion, microcracks are produced, and autocatalytic corrosion emerges. Finally, the mechanism of nitrogen action in the forming process is analyzed. It is shown that the pre-positioned TC4 alloy powder is nitrided under the action of the temperature field generated by SLM molding, and the pre-nitrided TC4 alloy powder melts into the specimen. In addition, nitrogen gas permeates from the gas/liquid interface to the molten pool surface, where it reacts with titanium to form nitrides. Driven by convection within the molten pool, these nitrides are subjected to mass transfer towards the interior of the molten pool (Fig. 14).

In this study, synchronous nitriding is achieved during SLM of a TC4 alloy in atmospheres with varying nitrogen volume fractions. Subsequently, an in-depth analysis of the microstructures, physical phases, and properties of the specimens fabricated with different nitrogen volume fractions is performed. In addition, the role of nitrogen gas during the SLM formation process is investigated. The results show that as the nitrogen volume fraction increases from 0 to 25%, the β-pillar crystals gradually decrease, and the nitrogen element mostly exists in the form of a solid solution. When the nitrogen volume fraction is increased from 50% to 100%, the titanium nitride organization appears and gradually increases, and the morphology changes to dendritic. In terms of properties, the TC4 alloy microhardness gradually increases with increasing nitrogen volume fraction. A moderate increase in nitrogen volume fraction improves the wear resistance of SLM-formed TC4 alloys, whereas an excessively high nitrogen volume fraction reduces their corrosion resistance. To ensure that the SLM-formed TC4 alloy has good wear and corrosion resistance, the nitrogen volume fraction should be controlled in the interval of 25%?50%. Nitrogen plays a role in the SLM forming process of the TC4 alloy in two ways: first, under the influence of the forming temperature field, it causes the pre-positioned TC4 alloy powder to undergo nitriding and the pre-nitrided powder melts and enters into the alloy; second, it permeates from the gas/liquid interface into the melt pool, achieving inward mass transfer with the help of the melt pool convection.

CoCrFeNi-based high-entropy alloys have received extensive research attention owing to their outstanding performance. By adding Al elements, the alloy properties can be further optimized to produce lighter high-entropy alloys, which show great potential for research and application in military, aerospace, nuclear industry protection, and many other fields. However, harsh acidic environments, such as acidic wastewater and acid rain, cause significant industrial corrosion, severely threatening the service performance of materials. Therefore, it is essential to enhance the corrosion resistance of the material to acidic conditions while ensuring its strength. In recent years, selective laser melting (SLM) technology has been widely used to prepare high-entropy alloy components towing to its desirable characteristics such as rapid forming, high precision, and good density. SLM is a layer-by-layer construction process that involves rapid cooling and heating cycles, which can easily lead to the accumulation of residual thermal stresses within the material. Studies have shown that heat treatment can improve the residual stress in alloys, thereby enhancing their mechanical and corrosion resistance properties. In this study, an AlCoCrFeNi high entropy alloy was prepared via SLM and post-processed under different heat treatment temperatures to reveal the influence of different heat treatment temperatures on the microstructure, mechanical properties, and corrosion resistance in H₂SO₄ solution with concentration of 0.5 mol/L.

Pre-alloyed AlCoCrFeNi high-entropy powder, with a near-atoms ration, was prepared via aerosol atomization. The stainless steel measuring 200 mm×200 mm×200 mm was selected as experimental substrate. The SLM forming AlCoCrFeNi high entropy alloy experiment used a laser system. Based on the preliminary process exploration, the determined optimal experimental parameters were as follows: laser power of 200 W, scanning speed of 1000 mm/s, layer thickness of 50 μm, spot diameter of 85 μm, scan rotation angle between adjacent layers of 67°, scan spacing of 80 μm, overlap rate of 50%, and the chamber was filled with protective gas (argon) during processing. A resistance furnace was used for the heat treatment process of the samples; a diffractometer was employed to analyze the microstructure composition of the alloy; a scanning electron microscope (SEM) was utilized for microscopic characterization of the samples; an energy dispersive spectrometer (EDS) was used to characterize and analyze the elemental distribution and composition in the alloy structure; the fully automatic Vickers hardness tester was used to measure the surface hardness of the samples before and after heat treatment; the diffractometer was used to analyze the residual stress values of the alloy samples; at room temperature, an electrochemical system workstation was used to perform electrochemical corrosion tests on the samples in a H₂SO₄ solution with concentration of 0.5 mol/L as the corrosive medium, and the SEM was used to observe the corrosion morphology.

The alloy exhibits good phase stability at 650 ℃. When the heat treatment temperature is increased to 850 ℃, the phase transitions from a single body-centered-cubic (BCC) phase to a BCC+ face-centered-cubic (FCC) biphasic solid solution structure. As the heat treatment temperature rises, the intensity of the FCC phase diffraction peak gradually increases, and following high-temperature heat treatment, the diffraction peak shifts toward smaller angles (Fig.2) Following high temperature heat treatment, the microstructure of the alloy is composed of a BCC phase and precipitated phase FCC (Fig.3). The elements exhibit the enrichment phenomenon (Fig.4). As the heat treatment temperature increases, the hardness and residual stress of alloy show a more obvious downward trend (Fig.6). The Nyquist curves of the original alloy, 650 ℃, and 850 ℃ heat-treated alloys basically show complete capacitive arcs, while those of 1050 ℃ and 1250 ℃ heat-treated alloys exhibit incomplete capacitive arcs. As the heat treatment temperature increases, the impedance value and phase angle of the alloys gradually increase, enhancing their corrosion resistance (Fig.7). White corrosion products and pitting are evident on the surface of the original alloy and 650 ℃ heat-treated alloy. As the heat treatment temperature increases, the corrosion products on the surface of the alloy gradually decrease, and only a large number of uniform pitting exists (Fig.9).

This study employs selective laser melting technology to prepare AlCoCrFeNi high entropy alloys and investigates the effects of heat treatment temperature on the microstructure, mechanical properties, and electrochemical corrosion behavior in a H₂SO₄ solution with concentration of 0.5 mol/L. Microcracks are present within the melt pool and at the boundaries of SLM-formed AlCoCrFeNi high entropy alloys, which consist of a single BCC phase. The alloy exhibits good phase stability following heat treatment at 650 ℃. As the heat treatment temperature increases, the SLM-formed AlCoCrFeNi high entropy alloy transitions from a single BCC phase to a biphase solid solution structure of BCC+FCC, with the surface microhardness decreasing to 336 HV and residual stresses being relieved. The alloy treated at 1250 ℃ has optimal corrosion resistance.

Laser powder bed fusion (LPBF) technology for fabricating ceramic particle-reinforced nickel-based composites is one of the effective methods to enhance the mechanical properties of Inconel 718 alloy. In the reported studies, the ceramic mass fraction is predominantly below 2%. Although some studies have reported TiC mass fraction up to 5% in Inconel 718 alloy, the impact of TiC particles on mechanical properties has still not been thoroughly investigated. Consequently, the contribution of TiC particles to the mechanical properties in high TiC content TiC/Inconel 718 composites remains ambiguous, and the influence of TiC content on the microstructure is not yet well-defined.

In this study, TiC/Inconel 718 composites with TiC mass fraction of 1.5% and 3.0% were prepared by the LPBF technology. The effects of TiC content on the microstructures and mechanical properties of the composites were systematically analyzed. In addition, the effect of TiC particles on the tensile process of LPBF-TiC/Inconel 718 composites was analyzed.

The thermal conductivity of TiC is significantly higher than that of Inconel 718 alloy, which accelerates the cooling rate of the composite melt pool. This reduces the primary dendrite spacing in the LPBF-TiC/Inconel 718 composite, leading to a more uniform microstructure. The layered structure becomes less pronounced, and the microstructure is significantly refined. Furthermore, as the mass fraction of TiC increases, the curvature radius of the fusion line in the LPBF-TiC/Inconel 718 composite gradually increases, and the melt pool becomes flatter. This further promotes the epitaxial growth of dendrites within the melt pool, enhancing the texture strength of the LPBF-TiC/Inconel 718 composite. The addition of TiC particles significantly improves the mechanical properties of the LPBF-TiC/Inconel 718 composite and the yield strength increases when the mass fraction of TiC particles increases. The effect of TiC on the yield strength of the LPBF-TiC/Inconel 718 composite can be analyzed from three aspects: coefficient of thermal expansion (CTE) mismatch strengthening, load strengthening, and fine grain strengthening. The results show that the contribution of these three strengthening mechanisms to the yield strength is positively correlated with the TiC content. CTE mismatch strengthening and fine grain strengthening are the primary contributors to the increase in yield strength of the LPBF-TiC/Inconel 718 composite, while the contribution of load strengthening is relatively minor. Additionally, during the tensile process of the LPBF-TiC/Inconel 718 composite, TiC particles may debond and spall, forming pits that act as crack initiation sites and propagation paths. This increases the numbers of potential failure initiation points, thereby reducing the elongation of the composite.

As the TiC content in the composite increases, a notable transformation occurs in the microstructure of the TiC/Inconel 718 composite. Specifically, the microstructure becomes more uniform, and the distinct layer-band structure becomes less pronounced and eventually indistinguishable. This uniformity is attributed to the consistent distribution of TiC particles throughout the matrix, which facilitates a more homogeneous microstructure. The addition of TiC particles does not fundamentally alter the as-deposited microstructure of the composite. Instead, it enhances the characteristic of columnar crystal epitaxial growth, which is a key feature of the as-deposited state. This enhancement is observed through the refinement of the microstructure, where the primary dendrite spacing is significantly reduced from 502.2 nm to 355.3 nm. This reduction in dendrite spacing is a direct result of the increased nucleation sites provided by the TiC particles, leading to a finer and more uniform microstructure. Concurrently, the increased curvature radius of the fusion line in the TiC/Inconel 718 composite plays a crucial role in the growth of dendrites within the melt pool. This directional growth is induced by the geometric changes in the melt pool, which in turn enhances the texture strength of the material. The texture strength increases from 4.27 to 12.76, indicating a significant improvement in the material anisotropic properties and overall mechanical performance. Compared to the as-deposited Inconel 718 composite, which serves as a baseline for comparison, the TiC/Inconel 718 composite exhibits a marked enhancement in mechanical properties. This improvement is evident in the increased yield strength, ultimate tensile strength, and hardness of the composite. However, it is important to note that the elongation of each composite is reduced to varying degrees. This reduction in elongation is attributed to the presence of TiC particles, which can act as stress concentrators and potential crack initiation sites. Despite this decrease in elongation, the overall mechanical performance of the TiC/Inconel 718 composite is significantly superior to that of the as-deposited Inconel 718 composite, making it a promising material for applications requiring high strength and durability.

Inconel 738 alloy is a precipitation-reinforced, cast, nickel-based superalloy reinforced with γ′-Ni₃(Al, Ti) phase. The alloy owns a high volume fraction of the γ′-reinforced phase and has excellent creep resistance, strength, corrosion resistance, and good microstructure stability at high temperatures, rendering it a key material for aeroengine turbine blades. Using conventional methods of casting, forging, and machining, the complex structures of aeroengine turbine blades—which have multiple holes and thin walls—face high processing challenges and require long processing cycles. Recently, selective laser melting (SLM) technology has been especially useful for high precision and rapid manufacturing of complex thin-walled structures—including medium and small overhangs, complex internal cavities, and profiles. Thus, it serves as a new method for manufacturing aeroengine turbine blades. However, for many conventionally cast Inconel 738 alloy components, the extremely high-temperature gradients (≥10⁵ K/cm), rapid cooling rates (≥0.01 m/s), and cyclic thermal effects inherent to SLM technology can form cracks during SLM processing and subsequent heat treatments. The laser energy density plays a key role in determining the behavior of the molten pool formed from the nickel-based high-temperature alloy powder layer during the SLM process, which in turn influences the microstructure and mechanical properties of the formed aeroengine parts. Thus, the goal of this study is to investigate the effects of scanning speed and laser power on the densification of this alloy within the optimal range of laser energy densities, address the internal defects in additively manufactured Inconel 738 alloy, and optimize the parameters of the forming process.

Herein, we employed SLM technology as the preparation method and investigated the effects of laser power and scanning speed on densification of the SLMed Inconel 738 alloy. We used a one-factor experimental design and fabricated 50 Inconel 738 blocks. We analyzed the blocks using optical microscopy (OM) and calculated their densification using ImageJ software. The results reveal that—within a selected range of the printing parameters—an optimal interval of energy density exists that affects the densification of the fabricated alloy. In particular, higher scanning speeds lead to lower densification within this optimal interval, and the influence of the laser power on the densification aligns with that of the energy density. We then selected five blocks for mechanical property analyses, which were used to establish the relationship between the energy density, laser power, scanning speed, and the mechanical properties of the alloy. We determined the optimal processing parameters—i.e., those that resulted in the highest densification and the best mechanical properties—from these relationships.

Our analysis of the effects of different energy densities demonstrates that the densification of the Inconel 738 alloy initially increases and then decreases as the laser energy density increases during the SLM process. We find that the optimal range of laser energy densities is between 60 and 70 J/mm3. However, within a similar range of energy densities, the densification of the printed alloy samples displays fluctuations with substantial variations (Fig. 8). Moreover, as the scanning speed increases, the densification of the Inconel 738 blocks continuously decreases (Fig. 9). The relationship between laser power and densification follows a trend similar to that of the energy density, where the densification first increases and then decreases. In particular, peak densification occurs at an energy density (E) of 66.67 J/mm3 (Fig. 9). The samples obtained with a scanning speed of 1100 mm/s and laser powers of 260 W and 270 W exhibit similar energy densities and densifications. A comparison of the mechanical property curves for these two samples shows that their yield strengths, tensile strengths, and elongations are nearly identical. This indicates that—within the optimal range of energy densities—the laser power does not have a remarkable impact on densification. The samples obtained using a laser power of 270 W and scanning speeds of 1100 mm/s and 1150 mm/s exhibit similar energy densities, although the densification obtained with the scanning speed of 1100 mm/s (99.92%) is slightly higher than that obtained with 1150 mm/s (99.91%). The yield strength and elongation of the sample fabricated at a scanning speed of 1100 mm/s are considerably higher than those of the sample fabricated at 1150 mm/s. This demonstrates that, within the optimal range of energy densities, increased scanning speed adversely affects the mechanical properties, which is consistent with its effect on densification. The highest densification was observed in the samples fabricated with a scanning speed of 950 mm/s and a laser power of either 270 W (99.95%) or 260 W (99.97%); these were the two samples selected for mechanical property testing. We find that the yield strength of the sample fabricated with a laser power of 260 W and scanning speed of 950 mm/s is 563.199 MPa and that its tensile strength is 1446.412 MPa, both slightly higher than those of the sample fabricated with a laser power of 270 W and a scanning speed of 950 mm/s. The elongation of the sample fabricated with a laser power of 260 W and scanning speed of 950 mm/s reaches 9.2%, corresponding to the highest energy of plastic deformation. The sample found to have the optimal mechanical properties is the one fabricated with a laser power of 260 W, scanning speed of 950 mm/s, and pitch of 0.10 mm.

Herein, we used scanning electron microscopy, OM, and tensile testing to investigate the influence of the laser energy density on the densification, metallurgical defects, and mechanical properties of the SLM-formed high-temperature alloy Inconel 738. Based on the optimal laser energy density, we further analyzed the effects of the laser power, scanning speed, and pitch on the densification of the samples. Accordingly, we draw the following conclusions from these experimental results:

(1) The laser energy density plays a critical role in the densification of the SLM-formed Inconel 738 alloy samples. The densification initially increases sharply and then decreases gradually as the laser energy density increases. The maximum densification of Inconel 738 is approximately 99.97% at E=68.42 J/mm³.

(2) Although the optimal laser energy density for the SLM-formed Inconel 738 alloy is determined, other different parameters will still affect the densification under the same laser energy density. This occurs because the energy density is influenced by a combination of factors, including the laser power (P), scanning speed (v), and pitch (s).

(3) After determining the optimal energy density, the effect of the laser power on the densification is still not clear. The densification of the sample decreases as the scanning speed increases. Given a fixed laser energy density, the densification of the sample can be modified by adjusting the scanning speed.

(4) For the SLM-formed Inconel 738 alloy, the printing parameters that yield the optimal mechanical properties are as follows: E=68.42 J/mm³, P=260 W, v=950 mm/s, and s=0.10 mm.

Tantalum, which is used in bone implants in the medical field, must be processed with optimal process parameters according to individual differences. Selective laser melting (SLM) technology helps meet this requirement, and the forming process of SLM determines the shaping quality and performance of tantalum. To date, few studies have been conducted on the forming process and performance of tantalum by SLM, and it is particularly important to study the process of SLM-formed tantalum. In this study, to obtain tantalum samples with excellent forming performances, the SLM forming process parameters of tantalum are explored using the control variable method, and tantalum samples with a smooth surface and dense inner are obtained. The microstructures and mechanical properties of the tantalum samples are analyzed. This provides a reference for further optimization of the SLM forming process of tantalum.

In this study, three identical tantalum samples were fabricated under the same conditions by varying the laser power and scanning speed for SLM forming with pure tantalum powder as the material. The surface morphologies and internal defects of the samples were characterized by optical microscope (OM). The densities of the samples were measured using the Archimedes drainage method, and the microhardness of the bottom surface of the samples was characterized by a Vickers hardness tester. The optimal process window and parameters were obtained through comparative analysis. As the next step, cross-sectional erosion and electrolysis of the tantalum samples formed by the optimal process were carried out, and the microstructures were characterized by OM. The crystal structure and grain orientation were analyzed by scanning electron microscope (SEM) with an electron backscatter diffraction (EBSD) probe. Further, the tensile samples were prepared with the optimal process parameters, and the tensile test was performed using an electromechanical universal testing machine. The fracture morphology was characterized using SEM.

When the laser power is 350 W and the scanning speed is 750 mm/s or 850 mm/s, the cooling rate is faster. At this time, the energy density is low and insufficient to completely melt the powder. The temperature gradient along the scanning direction is high; therefore, the surface of the sample presents a discontinuous scanning trace, and there are many pores inside. The microhardness values are 247.15 HV_{and 224 HV,respectively. When the scanning speed is changed in the range of 350?550 mm/s, the time that the laser beam remains on the powder increases, the energy density of the input molten pool increases, and the powder melting degree improves. Although the density of the sample is higher at this time, the powder melting speed is faster; therefore, argon gas remains inside the sample and forms small holes. Small droplet splashes and pores are observed on the surface, and the local hardening effect on the surface results in higher microhardness values. When the scanning speed is 650 mm/s, the melting and solidification of the powder reach equilibrium. No obvious pores or splashes are observed inside the sample, and the surface of the melt channel is flat and pinnate. The microhardness of the sample is 260.92 HV (Figs. 3, 5, 7, and 8).}

When the scanning speed is fixed at 650 mm/s and the laser power varies within the range of 150?300 W, the laser energy density is relatively low, the amount of powder melted per unit volume is lower, surface spheroidization and agglomeration are obvious, and unfused defects appear inside. The density is less than 99.4%, and the microhardness increases from 148.05 HV to 253.62 HV. When the laser power is increased to 350 W, the laser energy density increases, which effectively suppresses the spheroidization and agglomeration. The surface of the sample forms a regular melting channel, and the internal bonding is tight, without obvious defects. When the laser power is 400 W, the laser energy is high, molten pool is unstable, and surface is continuously flat with splashes and pores. Although the sample is dense, there are pores inside the sample, and there are cracks and unfused defects at the edges (Figs. 4, 6, 7, and 8).

The optimal process window and parameters are determined by analyzing the surface morphology, internal defects, and density. The lengthwise-section microstructure of the tantalum samples prepared with the optimal process parameters has relatively large axis-to-diameter ratios and grows across layers along the forming direction. The crystal orientation is not obvious, which is related to the interlayer rotation angle of 67° (Figs. 2, 9, and 10). The mechanical properties are better than those of tantalum samples formed by traditional methods, with a yield strength of 668 MPa. The fractures of tantalum tensile samples prepared using the optimal process parameters have macroscopic plastic deformation characteristics, and there are dimples with different sizes, indicating that the fracture mechanism is ductile fracture (Figs. 11 and 12 and Table 3).

In the process of forming tantalum by SLM, a higher scanning speed or lower laser power produces a larger temperature gradient, which causes obvious spheroidization on the surface of the sample, accompanied by pores or unfused defects inside, and the density of the sample is low. A lower scanning speed or higher laser power produces excess laser energy, resulting in excessive melting of the powder. At this time, the cooling rate of the molten pool is low, resulting in cracks at the edge, a hardening effect on the surface with a small amount of spatter, and pores inside.

Large-scale integral hydraulic multi-way valves in engineering machinery feature complex internal channels and numerous overhanging structures with interconnected oil passages, posing significant demands on sand core strength during casting. Traditional manufacturing techniques achieve a casting success rate of only about 60%, presenting substantial challenges such as high mold costs and lengthy processing cycles, hindering rapid product launch. Three-dimensional (3D) printing technology offers several potential solutions. Recent studies by domestic and international researchers have successfully applied selective laser sintering (SLS) and binder jetting (3DP, PCM) technologies to relatively simple structural components such as casings, boxes, and cylinder bodies. However, successful instances of rapid manufacturing of hydraulic valves are scarce. Owing to inherent weaknesses in the strength and density of 3D-printed sand cores, the rapid casting of large, complex internal cavity-structured integral multi-way valves using 3D printing remains problematic, with a casting success rate below 60%. Furthermore, research and applications in this field are limited. Therefore, there is an urgent need to develop rapid manufacturing technologies for integral hydraulic multi-way valves based on sand mold 3D printing, aiming to advance sand mold additive technology in engineering machinery, equipment manufacturing, aerospace, and other fields.

This study conducted a comparative analysis of the performance of two typical 3D printing processes for rapid sand mold manufacturing to identify the most suitable 3D printing technology for integral hydraulic multi-way valves. Focusing on a representative integral hydraulic multi-way valve, the research examined its structural characteristics and casting challenges through failure case analysis. Using finite element simulations, the filling processes of single-layer and composite casting systems were investigated by comparing metal flow velocity during mold filling. Based on these analyses, a composite casting system was developed specifically for 3D printing applications. Additionally, exhaust systems were designed for both internal sand cores and external molds, incorporating reinforced core and high-temperature-resistant coating. Finally, rapid casting experiments and mass production verification were conducted for integral hydraulic multi-way valves.

A rapid casting solution for 3D-printed integral hydraulic multi-way valves was developed, incorporating a “composite casting system, conformal exhaust, strength enhancement, and temperature-resistant coating” approach. Through rapid casting and testing verification, the results demonstrate that the integral multi-way valve exhibits good overall forming quality, with well-formed internal oil passages showing a straightness of ≤0.28 mm/100 mm. Internal inspection of sectioned valve bodies reveals no defects such as pores or flash with dimensions ≥0.3 mm. The average hardness deviation within the valve body is ≤5%, confirming the feasibility of rapid casting for large-scale integral hydraulic multi-way valves using 3D printing technology.

A comparative analysis of different 3D printing technologies, including selective laser sintering (SLS) and binder jetting (3DP), was conducted to evaluate sand mold properties such as strength, printing accuracy, and gas evolution rate. The results indicate that 3DP sand mold printing technology is more suitable for the rapid casting of large-scale integral multi-way valves. By comparing metal flow velocity in single-layer and composite casting systems, a composite casting system for integral hydraulic multi-way valves was developed based on 3D printing technology. This system leverages the advantages of single-layer casting while preventing continuous impact on local sand cores, thereby improving casting success rates. A performance enhancement method for 3D-printed sand cores was proposed, incorporating “conformal exhaust, strength enhancement, and temperature-resistant coating,” which increases overall bending strength and effectively prevents internal defects in valve bodies. The developed rapid casting process successfully produces integral hydraulic multi-way valves, with CT scans showing no casting defects such as shrinkage cavities or porosity exceeding the specified size limits. The valve body exhibits uniform and stable internal hardness, meeting industrial application requirements. The technical achievements of this research demonstrate the feasibility of mass production and provide a foundation for the rapid manufacturing of large components with complex internal structures.

TC11 titanium alloy, an α+β high temperature titanium alloy with excellent comprehensive mechanical properties, is widely used as critical structural components in the aerospace industry. Compared to traditional “integral forging + machining” techniques, wire laser additive manufacturing offers advantages such as high deposition efficiency, low cost, and enhanced material utilization when producing TC11 components. However, process pores tend to occur during the deposition process, which can degrade the overall mechanical properties of the components. Existing research in the welding field has demonstrated that oscillating laser technology, achieved through rapid laser oscillation, creates an “eddy” flow in the molten pool, unifying and ordering the flow state within it, enhancing keyhole stability, and extending bubble escape time, thereby effectively inhibiting the porosity formation. To suppress the porosity defects in wire laser additive manufacturing, laser oscillating is introduced into the process, resulting in a new technique known as laser oscillating wire additive manufacturing. To investigate the process characteristics of laser oscillating wire additive manufacturing of TC11 titanium alloy, the influence of process parameters on single track deposition was studied, and the microstructure and microhardness analysis of single track multi-layer deposition samples were conducted.

Deposition experiments were conducted on TC11 wire material using a laser oscillating wire additive manufacturing system. Cross sections and longitudinal sections (each with a length of 20 mm) for single track deposition were cut from the substrate using wire electrical discharge machining. After mounting, grinding, and polishing, the samples were etched using Kroll reagent. The samples were then observed and characterized for their microstructures using an optical microscope and a field emission scanning electron microscope. Subsequently, the deposition morphological dimensions and porosity were measured using Image J image processing software. Microhardness tests were performed on a digital display microhardness tester.

Result and Discussions The linear oscillating laser exhibits excellent suppression of porosity. When the amplitude of the linear oscillating laser is 1.0 mm, the suppression of porosity is highly significant. However, as the amplitude continues to increase, it greatly degrades the deposition surface quality without significantly improving the porosity suppression effect. The effectiveness of frequency in suppressing porosity varies with amplitude. At amplitudes of 1.0, 1.5, and 2.0 mm, a frequency of 50 Hz achieves excellent porosity suppression effect with a low porosity rate of 1.72 %. However, at an amplitude of 0.5 mm, a frequency of 150 Hz is required to obtain a significant improvement in porosity suppression (Fig. 3). The linear oscillating laser exhibits an energy distribution characteristic with high energy on both sides and low energy in the middle (Fig. 6), which is consistent with the bimodal structure feature of the molten pool (Fig. 4). As the equivalent line energy density increases, the molten pool gradually changes from a bimodal structure to a typical arc-shaped shape. Further research has been conducted into the microstructure and microhardness of the deposited layers of TC11 titanium alloy produced through laser oscillating wire additive manufacturing. The results indicate that the microhardness distribution is consistent with the microstructure evolution pattern. The combined effects of refined grains at the top and the abundance of fine acicular martensite α′ result in a gradual increase in microhardness along the deposition direction (Fig. 9).

The linear oscillating laser exhibits excellent suppression of porosity. When the amplitude exceeds 1.0 mm or the frequency is too high, it can adversely affect the surface quality of single track deposition. Considering both porosity and deposition surface quality, the laser oscillating parameters are optimized to be an amplitude of 1.0 mm and frequency of 50 Hz. The laser power and scanning speed are the determinants of the equivalent linear energy density of the oscillating laser, which affects the metal melting state and the geometric characteristics of the molten pool. With the optimized laser oscillation parameters, when the laser power increases from 800 W to 1400 W, the deposition layer width and melt depth increase by 1.645 mm and 1.214 mm, respectively. When the scanning speed increases from 0.4 m/min to 1.0 m/min, the deposition layer width and melt depth decrease by 1.422 mm and 0.512 mm, respectively. The wire feeding speed directly determines the amount of material fed into molten pool and is positively correlated with the height of the deposition layer. During the deposition process, the equivalent energy distribution of the oscillating laser exhibits a characteristic of high energy on both sides and low energy in the middle, which is consistent with the bimodal structure feature of the molten pool. As the equivalent line energy density increases, the molten pool gradually changes from a bimodal structure to a typical arc-shaped shape. In single track multi-layer TC11 deposition samples, the grain size gradually decreases along the deposition direction, in the range of 79.7?223.4 μm, with a transition from coarse β-columnar grains to fine equiaxed grains. The martensite α′ is more elongated and numerous at the top of the multi-layer deposition layer. The typical microstructures of multi-layer deposited TC11 samples are characterized by basket-weave and colony structures. The refinement of grains and the abundance of fine acicular martensite α′ at the top are the main reasons for the gradual increase in microhardness along the deposition direction, reaching a maximum hardness of 499.5 HV.

The joining components of titanium alloy and steel have been widely used in aerospace, automobile manufacturing and other fields because of their high strength and light weight. However, the formation of Ti-Fe intermetallic compounds seriously affects the mechanical properties of joints welded with titanium alloy and steel dissimilar metals, which limits their further application. The selection of appropriate welding methods and the use of the interlayer or filler materials are of great significance for obtaining efficient and reliable titanium alloy/steel joints. Dual-beam laser welding can accurately control the composition of intermetallic compounds in dissimilar metal welding, and is more stable than single laser welding. The use of single interlayer or filler material can inhibit the formation of Ti-Fe intermetallic compounds, but other intermetallic compounds might be formed. Therefore, in this study, the dual-beam laser welding method is used and V is added as an interlayer on the basis of using CuSi3Mn1 filler wire to optimize the interface structure of the joint and improve the mechanical properties of the joint.

TC4 titanium alloy and QP980 high-strength steel are used as the base materials, and CuSi3Mn1 alloy is used as filler wire. A novel welding method with coaxial pulsed-continuous dual-beam laser heat source assisted by V interlayer is adopted. The microstructure and mechanical properties of the joints with different V interlayer thicknesses are compared. The macroscopic morphology of the joint is observed by optical microscope, and the mechanical properties are tested by universal tensile testing machine. The element distribution and phase composition of the interface microstructure of TC4/QP980 joints with and without V interlayer are tested and analyzed by electron probe X-ray micro-analyzer.

The V interlayer-assisted dual-beam laser welding process can obtain TC4/QP980 welded joints with smooth weld surface and high strength. When the thickness of the intermediate layer is small, the high energy heat input of the dual-beam laser destroys the V interlayer, resulting in a large influx of CuSi3Mn1 filler wire and forming a thick intermetallic compound reaction layer between the V interlayer and TC4. In contrast, when the thickness is large, there are pores and incomplete welding (Fig. 2). The joints with a 0.10 mm-thick V interlayer have the better mechanical properties, and the maximum tensile shear load reaches 457 N/mm (Fig. 3). The steel-side interfaces are all composed of Fe solid solution (ss) + Cu solid solution (Figs. 4 and 5). However, the microstructures of the TC4-side interfaces are different. TC4-side interface without V interlayer is mainly composed of TiCu and Ti₅Si₃ intermetallic compounds with large hardness and brittleness, and the thickness of the intermetallic compounds layer is about 200 μm (Fig. 6). The microstructure of the TC4-side interface using V interlayer is optimized to two layers, and the total thickness is reduced to about 110 μm (Fig. 7). The upper layer is composed of α-Cu/(Fe, V, Cu, Si) ss/V, and the thicknesses of the ss layers are less than 2 μm. The microstructure of the lower layer is V/(Ti, V) ss+Ti₂Cu/Ti₂Cu+TiCu/(Ti, Cu) ss+Ti₂Cu/α-Ti. The formation of Ti₂Cu, which has lower brittleness than TiCu, and the solution bonding both improve the bonding strength of the TC4-side interface.

In this study, a novel welding method with coaxial pulsed-continuous dual-beam laser heat source assisted by V interlayer is proposed. The main problem of TC4 titanium alloy and QP980 high-strength steel welding is that it is easy to form Ti-Fe intermetallic compounds between them, which leads to stress concentration and finally the joints fracture. The thickness of the V interlayer is adjusted to match the dual-beam laser welding process. The maximum tensile properties can be obtained when the thickness of the V layer is 0.10 mm. The application of a V interlayer can improve the interface microstructure and reduce the thickness of the bonding layer to about 100 μm. Overall, the proposed coaxial pulsed-continuous dual-beam laser welding method assisted by V interlayer can effectively suppress the formation and thickness of intermetallic compounds in TC4 and QP980 high-strength steel welded joints, optimize the interface microstructure, and thus improve the mechanical properties of welded joints. This study provides a new theoretical support for solving the problem of dissimilar metal welding of titanium alloy and steel, and a reference for the application of dual-beam laser welding technology in high-performance dissimilar metal structures.

In recent years, additive manufacturing of metal composites has been studied extensively, and nanoscale ceramic particles have been introduced into the metal matrix as a reinforcing phase, which can improve the microstructure of the alloy and the comprehensive properties of a material. The additive-manufactured composite specimen, which has great potential for industrial production and application, is found to have a large residual stress. Consequently, further heat treatment is required. However, most of the studies directly use the heat treatment system of forgings, and the microstructural characteristics of additive manufacturing are quite different from those of the forgings. Therefore, studying the corresponding heat treatment system based on the characteristics of additive manufacturing is necessary. This study considers TiB₂/In718 composites as the research object to explore the effects of different heat treatment methods on the microstructure and precipitated phase characteristics of the composites to provide a certain reference for subsequent research.

The TiB₂/In718 composites were prepared by laser additive manufacturing equipment and the resulting TiB₂/In718 composites were subjected to high-temperature homogenization heat treatment. The morphologies of the precipitated phases were observed using a field emission scanning electron microscope (SEM), which was combined with the energy dispersive spectrometer (EDS) to analyze the distribution of elements in the alloys.

The TiB₂/In718 composites were subjected to high-temperature heat treatment. The microstructures of TiB₂/In718 composites exhibit dendrites following heat treatment at 1175 ℃, and no recrystallization grains are observed, indicating that the recrystallization temperature of TiB₂/In718 is above 1175 ℃. When the high-temperature homogenization temperature is 1200 ℃, the dendrite substructure disappears, grain morphology is equiaxed, and large-angle grain boundaries exhibit jagged arch traces. Moreover, some recrystallized grains can be observed near the original grains (Fig. 1), which is characteristic of the early stage of recrystallization. It provides a favorable location for the generation of recrystallized grains, which nucleate through the grain boundary arching mechanism. When the high-temperature homogenization temperature is 1225 ℃, the TiB₂/In718 composite material undergoes obvious recrystallization, a large number of fine equiaxed grains are formed, grain refinement occurs, and the grain morphology is fully equiaxed, indicating that the TiB₂/In718 composite is completely recrystallized. The In718 alloy can be completely recrystallized following heat treatment at 1100 ℃ for 1 h, indicating that the recrystallization temperature of In718 increases following the addition of TiB₂ nanoparticles. This is because the TiB₂ strengthening phase particles hinder the slippage of dislocations and the migration of grain boundaries, which is inconducive to the nucleation and growth of recrystallization and hinders the recrystallization process, requiring higher crystallization temperatures. When the high-temperature homogenization temperature is 1250 ℃, the recrystallized grains grow further. Additionally, during the high-temperature homogenization process, the TiB₂ nanoparticles decompose, and owing to the low solid solubility of B atoms in the nickel matrix, the decomposed B atoms segregate at grain boundaries to form borides. Following direct aging (DA) heat treatment, a large number of Laves phases are present between the dendrites of TiB₂/In718 (Fig.5), and this formation of the Laves phase is related to the microscopic segregation of Nb elements in the dendrite part during the solidification of the alloy. Meanwhile, TiB₂ dissolves in the matrix during solidification, resulting in the precipitation of Nb and other elements from the supersaturated matrix and diffusion to the dendrites. This leads to the segregation of Nb and other elements between the dendrites and the formation of the Laves phase. Following solution treatment and aging (STA) heat treatment, the Laves phase partially dissolves, and its morphology changes from chain to block (Fig. 6).

When the high-temperature homogenization treatment is performed at 1175 ℃, the laser-solid-formed TiB₂/In718 composites do not undergo an obvious recrystallization process. When the temperature exceeds 1200 ℃, obvious recrystallization occurs, recrystallization grains are formed, and with the increase of solution treatment temperature, many columnar crystals are transformed into equiaxed crystals, and the recrystallization phenomenon becomes prominent. When the temperature increases to 1225 ℃, the recrystallized grains of TiB₂/In718 composites grow. Additionally, when the samples are treated with a solid solution above 1200 ℃, the TiB₂ nanoparticles react with the matrix to form boride. With the increase in the number of TiB₂ nanoparticles and the number of cell crystals, the distribution morphology of the Laves phase also changes from chain to network distribution, and the shape of the Laves phase changes from a long chain to a block. Following STA heat treatment of TiB₂/In718 composites, the distribution and structure of alloying elements become more uniform, and the layer structure is basically eliminated. In addition, during the heat treatment process, the Laves phase partially dissolves, and the Laves phase morphology changes from a long chain to a block.

In the process of laser deep penetration welding, the laser beam interacts with the material to produce photo plasma, metal vapor, plume, and welding defects such as spatter and porosity. Among them, the root cause of small porosity is the instability of small holes, the collapse of the small hole, and the gas that enters the deep penetration welding hole is too late to escape during the solidification of the weld pool, which eventually affects the stability of the welding process. In order to study the influence of gas protection modes on the laser deep welding process of 15 mm thick 316H stainless steel and reduce the influence of photo plasma, metal vapor, plume, etc., on the welding process, this paper carries out a study on the influence law of different types of gas protection modes on the laser deep penetration welding process. This paper provides a theoretical basis for the application of the laser deep penetration welding technology in 316H stainless steel plates.

First, the gas protection device is specially designed, and 5 gas channels are designed, namely molten pool protection gas (MP), transverse blowing protection gas 1 (TB1), transverse blowing protection gas 2 (TB2), side blowing protection gas (SB), and tail dragging protection gas (ST) (Fig. 2). Second, seven different types of gas protection modes are established. Finally, the key process parameters such as laser power are kept unchanged, and only the gas mode is changed. The welding test is carried out successively from the gas protection mode A to G. Meanwhile, the images of plume and metal vapor in the welding process under different gas protection modes are observed with the help of visual high-speed photography (Fig. 6). Furthermore, the cross-sections of 7 beads are compared and analyzed, and the cross-sectional morphologies of beads are indirectly used to characterize the suppression effect of different gas protection modes on the plume glow during welding.

The addition of SB has little effect on weld penetration depth and width. The cross-sectional weld morphology of bead 7 is regular, the weld penetration depth reaches 10.1 mm, and the weld width reaches 2.5 mm (Fig. 4). However, from the observation of high-speed photography, it is found that the addition of SB has a blowing effect on the narrow and long plume similar to the laser beam focusing shape, and can reduce the spatter during the welding process (Fig. 6). Among them, the bead 6 has the worst shape (compared with the other 6 wbeads), especially irregular shape appears at the upper end of the weld, and it is tilted toward one side near the upper end of the weld. Combined with the gas mode F corresponding to bead 6, the interference of TB1 (Ar) and TB2 (Air) in this gas mode leads to drastic changes in the transient plume flow state, resulting in turbulence in the air flow near the weld pool. This results in poor deformation of the surface weld (Fig. 4). Under the G type gas protection mode, key welding process parameters such as laser power, welding speed, and defocusing amount are optimized again, achieving double-side formation from single-pass welding of 316H stainless steel (Fig. 8). The microstructures of the base metal, weld metal, and heat-affected zone are all austenitic metal. The lowest impact energy of a single sample is 198 J, which is far greater than the minimum 90 J requirement. There is a shear lip on the impact fracture of the weld, and obvious plastic deformation occurs.

The addition of SB has almost no effect on the cross-sectional morphology of the bead, and the weld penetration depth and width are basically the same. However, from the observation of high-speed photography, it is found that the addition of SB has a blowing effect on the narrow plume similar to the laser beam focusing shape, and can reduce the splash during the welding process. Finally, it is determined that the gas protection mode of type G is the optimal gas configuration in the process of laser deep penetration welding. Under the G type gas protection mode, the key welding process parameters such as laser power, welding speed, and defocusing amount are further optimized to achieve double-side formation from single-pass welding of 316H stainless steel. The weld forming is excellent, and the relevant indicators meet the technical requirements. The investigation here provides a theoretical basis and a practical experience for the application and development of laser deep penetration welding technology in 316H stainless steel medium plates.

Laser deep penetration welding (LDPW) of high-reflectivity alloys presents significant challenges due to the complex interaction between laser and material. Traditional methods for monitoring weld penetration depth are often inefficient or lack real-time accuracy, leading to inconsistent welding quality. Optical coherence tomography (OCT) has emerged as a promising technique for non-destructive, real-time monitoring of weld depth. However, the effectiveness of OCT-based penetration depth monitoring heavily depends on algorithmic optimization. This study aims to develop an improved penetration depth monitoring algorithm tailored for OCT, enhancing its accuracy and robustness when applied to high-reflectivity alloy welding. By optimizing signal processing and feature extraction, this research seeks to advance real-time penetration depth measurement, contributing to more precise and efficient laser welding processes.

This study proposes an improved OCT weld penetration depth extraction algorithm, which optimizes the existing OCT measurement method through four key steps of “noise reduction?ghosting elimination?data fitting?depth correction.” First, background noise from the charge coupled device (CCD) camera is removed using brightness distribution analysis, where a threshold is set to retain only valid data points. Second, the local outlier factor (LOF) algorithm is employed to detect and eliminate optical ghosting effects by analyzing density variations, and the primary data region is identified using a depth-based segmentation approach. Third, a moving average filter is applied to the keyhole depth data to smooth measurement fluctuations and obtain a continuous depth curve. Finally, based on metallographic measurements, a multivariate regression model is established to analyze the influence of welding parameters (laser power, welding speed, defocus amount, etc.) on the correlation between keyhole depth and actual weld penetration depth. This model is used to correct OCT measurement errors. The optimized algorithm significantly improves the accuracy and reliability of OCT-based penetration depth monitoring, providing a more robust method for real-time welding depth measurement.

The proposed optimization algorithm effectively improves the accuracy of weld penetration depth measurement based on OCT. By integrating signal denoising, ghost image elimination, curve fitting, and penetration depth correction, the algorithm significantly enhances the reliability of depth estimation under various welding conditions. The results show that the optimized measurement data closely align with actual metallographic depths, with relatively small error, confirming its feasibility for online monitoring. Moreover, the algorithm performs well in both linear and oscillating welding processes, demonstrating adaptability to different materials and welding parameters. These verification results further validate the robustness and high precision of the algorithm, making it a viable solution for accurately measuring weld penetration depth in industrial applications.

This study addresses the challenge of low penetration depth measurement accuracy in laser deep penetration welding of highly reflective aluminum and copper alloys. An optimized OCT-based algorithm, incorporating signal denoising, ghosting elimination, data fitting, and depth correction, significantly enhances measurement precision by mitigating noise and ghosting interference. Validation results demonstrate that the optimized algorithm improves accuracy across different materials and welding parameters, maintaining error within 6%, making it a reliable approach for real-time weld quality monitoring and closed-loop control.

The presence of welding defects in the weld seam significantly impacts the strength and service life of structural components. Traditional detection methods often suffer from limitations such as high computational costs, slow inference speed, and large model sizes, which restrict their practical applications. This paper aims to address these challenges by proposing a lightweight real-time weld defect classification model ,TDRE-YOLO-cls, that achieves high accuracy while maintaining a small model size and fast inference speed.

Carbon steel is widely used in construction, bridges, shipbuilding, and automobile manufacturing due to its excellent mechanical properties, ease of processing, and relatively low cost. Selecting appropriate welding methods is crucial for ensuring the safety and reliability of carbon steel structures. However, in practical operations, defects such as dents and holes cannot be completely avoided. These defects can not only weaken the load-bearing capacity of weld joints but also lead to failures or accidents. Therefore, timely and accurate detection of weld defects is essential for ensuring the safe and reliable operation of welded structures.

With the development of artificial intelligence, non-destructive testing has shown significant improvements in precision and efficiency. However, traditional methods like magnetic particle inspection are limited to ferromagnetic materials and require cumbersome preparation, radiographic testing poses potential health risks and is costly, and ultrasonic testing is slow and complex in data processing. In contrast, laser-based non-destructive testing, despite its inability to detect internal defects, is widely used for surface defect detection due to its high precision, high sampling rate, and compact hardware.

We modified the YOLOv8n-cls architecture by introducing the Re-Parameterized Reshaping Convolutional Representation (RCR) module into shallow layers and the Spatial Pyramid Pooling (SPP) and Downsampling Position-Specific Attention (DPSA) modules into deep layers. The RCR module leverages RepConv blocks to efficiently extract multi-scale features. Meanwhile, the DPSA module employs special downsampling and compression mechanisms to reduce model parameters. Additionally, we proposed a Compressed Squeeze and Excitation (CSE) attention mechanism tailored for DPSA to enhance the extraction of critical information.

Specifically, we replaced the shallower C2f modules in the YOLOv8n-cls architecture with the RCR modules containing RepConv to obtain multi-scale features while controlling inference time . We introduced a DPSA module with special downsampling and compression mechanisms to effectively reduce the model parameter size. Finally, to further enhance the model ability to extract key information, we developed a dedicated attention mechanism for the DPSA module.

Experimental results showed that TDRE-YOLO-cls outperforms YOLOv8n-cls in several key metrics of Top-1 accuracy increased by 2.4%, weighted precision increased by 2.3%, and weighted recall increased by 2.4%. Notably, our model achieved these improvements while reducing the model parameter count by 52.1% and maintaining an inference time of 0.9 ms per frame (Table 2). To comprehensively evaluate the performance and generalization ability of TDRE-YOLO-cls, we conducted extensive experiments on an expanded dataset consisting of burrs, dents, holes, and no obvious defects, totaling 3792 samples. Training involved 200 epochs, with validation and test sets used for performance assessment.

Further comparison with various existing models, including YOLOv8n-cls, YOLOv11n-cls, YOLOv8s-cls, YOLOv11s-cls, MobileNetV3, and ShuffleNetV2, demonstrated the superiority of TDRE-YOLO-cls in terms of accuracy, inference time , and model size. On the test set, TDRE-YOLO-cls showed a Top-1 accuracy improvement of 2.4%, a weighted precision increase of 2.3%, and a weighted recall enhancement of 2.4%, while maintaining an inference time of 0.9 ms per frame and reducing the model size by 52.1% (Table 2). Ablation studies confirmed the effectiveness of each component in our proposed model, demonstrating its robustness and generalizability (Table 3).

Additionally, we performed a detailed analysis on the performance of TDRE-YOLO-cls across different types of weld defects, including burrs, dents, holes, and no defects. Our results indicated that TDRE-YOLO-cls achieves balanced performance across all categories, although there is still room for improvement, particularly in detecting dents. Future work will focus on optimizing the model to better handle specific types of defects, thereby enhancing overall accuracy and reliability.

The TDRE-YOLO-cls model effectively balances real-time performance, accuracy, and model size, making it suitable for industrial applications where hardware resources are limited. The proposed modifications, including the introduction of RCR, SPP, and DPSA modules, have been shown to significantly improve the model performance without compromising inference time . Future research will continue to refine the model, particularly focusing on improving its performance for challenging defect types such as dents, to achieve even higher overall accuracy. Moreover, we plan to explore the integration of advanced techniques such as semi-supervised learning and transfer learning to further enhance the model capabilities.

Because the fluid flow behavior in a molten pool is closely associated with the mechanical properties of the joint, the analyses and regulation of welding temperatures and flow fields are of great practical importance. The main objectives of this study are to reveal the interaction of the oscillating laser beam with the molten pool and to elucidate the underlying relationship between the welding temperature and flow fields as well as the microstructure of the joint.

The welding experiments were conducted using an oscillating laser wire-filling welding system with a butt joint configuration. A three-dimensional transient multiphysical numerical model was constructed to simulate the oscillating laser welding process of 5083 aluminum alloy. The microstructures of the joints were observed using an AOSVI M320P-HK830 polarizing microscope (PM) and a Zeiss Ultra 55 LE scanning electron microscope (SEM). The chemical compositions of the different phases of the joints were analyzed using an XFlash^?4010 energy dispersive spectroscope (EDS) detector. X-ray diffraction (XRD) analyses of the joints were performed using a D8 Advance X-ray diffractometer for phase identification. Nondestructive X-ray testing was performed to detect weld pores.

The results show that molten pool stirring produced by an oscillating laser beam facilitates heat transfer and thus reduces the temperature gradient of the molten pool. In addition, the oscillating laser beam stirs the molten pool to drive the fluid flow in the same direction and form vortices, thereby increasing the stability of the molten pool. When the oscillaton frequency is 100 Hz, a vortex with a flow rate of up to 0.7 m/s is formed, which significantly improves the stability of the molten pool. Furthermore, degassing the molten pool, refining the grain size of the weld, and promoting the precipitation of the β phase in the weld can be achieved by virtue of the oscillating laser beam stirring the molten pool. As the oscillation frequency increases from 0 to 100 Hz, the area fraction of β phase in the weld increases from 3.37% to 12.3%, whereas the weld porosity decreases from 5.02% to 0.2%.

Oscillating laser beam stirs the molten pool to facilitate heat transfer and reduce the temperature gradient of the molten pool. The stirring of the molten pool caused by the oscillating laser beam drives the fluid to flow in the same direction and forms vortices, which thereby increase the stability of the molten pool and result in a sound weld. Oscillating laser welding refines the grain size of the weld and hence increases the amount of grain boundary by stirring the molten pool, thereby promoting the precipitation of β phase in the weld. The molten pool created by oscillating laser welding is wide and shallow compared with that created via conventional laser welding; thus, the gas bubbles in the former have sufficient time to leave the molten pool, which reduces the gas porosity. Meanwhile, the stirring effect of the oscillating laser beam on the molten pool helps gas bubbles escape from the molten pool surface, which contributes to the reduction in weld porosity.

Laser shock micro-hydraulic bulging (LSMHB) is a novel high-strain-rate forming technology that combines the advantages of laser shock forming and hydraulic shock forming. It is characterized by high strain rate, excellent surface quality, and stable forming, effectively improving the forming performance of metallic materials, which has attracted increasing attention. Previous studies have mainly focused on the use of cylindrical liquid chambers, however the influence of liquid chamber geometry on forming performance has not been extensively explored. In conventional LSMHB, excessive thinning and localized stress concentrations are often observed at the fillet regions, leading to surface defects and even failure. Therefore, this study aims to improve the forming process by using a conical frustum liquid chamber, replacing the conventional cylindrical chamber. This research systematically investigates the effects of laser energy on material forming depth, thickness distribution, surface quality, and microhardness, using TC4 titanium alloy foil as the experimental material. The goal is to improve forming uniformity and address the excessive thinning at the fillet regions.

To evaluate the performance of the conical frustum liquid chamber in the LSMHB process, a series of experiments were conducted at varying laser energy levels (30%?90%). The TC4 titanium alloy foil used in this study had a thickness of 50 μm. The liquid chamber design was optimized to improve pressure distribution based on numerical simulations and the response surface methodology. During the experiments, the laser system generated a high-intensity nanosecond pulse, and the laser energy was precisely adjusted to ensure repeatability and consistency. Forming depth and thickness distribution were measured using the digital microscope, and surface roughness and 3D morphology were detected using the microscope. Microhardness analysis was performed using cross-sectional samples obtained through cold mounting and measured with the micro Vickers hardness tester to assess the impact of laser energy on material strengthening. Finally, the experimental results were analyzed to assess the advantages of the conical frustum liquid chamber in improving forming performance compared to the cylindrical liquid chamber.

The experimental results show that as laser energy increases, the forming depth increases uniformly. At 90% laser energy, the maximum forming depth reaches 287.30 μm, surpassing the forming depth of the cylindrical liquid chamber under the same conditions. This improvement is attributed to the optimized conical frustum liquid chamber, which allows the laser shock wave to be transmitted more evenly to the workpiece surface. Additionally, the thinning ratio at the fillet region remains stable across all laser energy levels, avoiding the excessive thinning and microhardness spikes typically observed in conventional LSMHB processes. Surface roughness analysis indicates that as laser energy increases, the surface roughness changes moderately. Microhardness tests show a consistent increase in material hardness with increasing laser energy and a minimum hardness increase of 17.28% at 90% laser energy. This enhancement is attributed to the high-strain-rate deformation mechanism, which promotes dislocation accumulation and strain hardening, thus strengthening the material. The study on the forming performance of TC4 shows that the optimized liquid chamber design effectively improves the forming capability of LSMHB process.

This study demonstrates that the introduction of the optimized liquid chamber in the LSMHB process results in a smoother thickness transition at the fillet region, improving the microstructural hardness of TC4 titanium alloy foil and significantly enhancing the forming performance. The conical frustum liquid chamber allows for a more uniform transmission of the laser shock wave, effectively solving the excessive thinning and microhardness spikes observed in the fillet region of traditional cylindrical liquid chambers. The results provide valuable theoretical and experimental support for optimizing the LSMHB process.

Metal oxides have garnered significant attention in modern materials science because of their exceptional piezoelectricity, semiconductor properties, and optical performance, which enable diverse applications. Compared with conventional two-dimensional (2D) architectures, three-dimensional (3D) structural designs can enhance performance by increasing the energy density of energy-storage devices, improve the mechanical properties of materials, and achieve functionalities that are not inherent to base materials. However, conventional lithography-based fabrication is limited to simple 2D patterning, which restricts micro/nanodevice integration. Additive-manufacturing techniques such as selective laser sintering and direct ink writing have been employed for 3D metal-oxide fabrication; however, their resolution remains constrained (10?100 μm) by nozzle dimensions and material rheology. Hence, two-photon polymerization (TPP) combined with thermal sintering is proposed, which involves the utilization of metal-ion-doped photoresists to create submicron 3D oxides after the removal of organics. However, chemically active ions (e.g., Fe3?) inhibit polymerization by quenching free radicals. Meanwhile, alternative scaffold-absorption-based methods immerse hydrogel templates in metal salt solutions to adsorb ions without interfering with TPP. Nonetheless, ion absorption reduces the hydrogel swelling capacity, thus resulting in low metal loading and porous sintered structures. Consequently, a universally applicable, high-precision fabrication strategy is urgently required that balances material adaptability with dense, defect-free 3D metal-oxide nanoarchitectures. The development of such methods can facilitate advanced applications in photonics, flexible electronics, and energy storage by fully exploiting the unique properties of 3D metal oxides.

This study proposes a versatile and efficient method for fabricating 3D metal-oxide nanoarchitectures. Leveraging the coordination mechanism between metal ions and carboxyl groups within a polymer scaffold, the approach integrates multiple metal ions into TPP-printed 3D polymer templates. Subsequent thermal sintering removes the organic framework, thus yielding high-quality 3D metal-oxide architectures with sub-500 nm feature sizes (minimum of 410 nm). The key fabrication steps are as follows: 1) TPP-based femtosecond laser printing of 3D polymer scaffolds; 2) immersion in 0.1 mol/L metal salt solutions for 120 min to facilitate metal-carboxyl coordination and ion infiltration; and 3) high-temperature sintering to decompose the polymer and form dense metal-oxide structures. This study systematically investigates factors affecting metal-scaffold coordination, such as solution concentration and immersion duration, to optimize ion uptake and structural fidelity. Compared with conventional hydrogel-based adsorption methods, this coordination-driven strategy enhances the metal-loading capacity and minimizes the porosity of the final sintered structures. The proposed method demonstrates broad applicability across various metal oxides while circumventing polymerization inhibition caused by reactive ions (e.g., Fe3?) in direct TPP approaches.

To achieve high shape fidelity in sintered metal-oxide nanoarchitectures, a novel TPP photoresist with enhanced metal adsorption efficiency is developed in this study, which significantly improves the ion-loading capacity of the scaffold during immersion [Fig. 1(b)]. The polymer scaffold, which is fabricated via femtosecond laser TPP, enables efficient metal ion incorporation via a dual mechanism [Fig. 1(d)]. First, the dissociation of carboxylic acid groups (—COOH) in the polymer network generates negatively charged —COO^- groups, thus inducing electrostatic repulsion-driven swelling to facilitate metal-ion diffusion. Second, stable coordination bonds form between oxygen lone pairs in carboxylate groups and metal ions, thus ensuring robust chemical adsorption. This synergistic approach overcomes the limitations of conventional hydrogel-based methods and minimizes the porosity of the final structures. Benefiting from TPP submicron resolution and controlled shrinkage during sintering, the method affords a NiO “buckyball” architecture with features measuring 410 nm at the minimum [Fig. 1(e)]. Complex architectures, including CoO woodpile arrays and Ni_0.5Co_0.5O curved lattices, demonstrate the capability of this technique for fabricating intricate 3D geometries and multi-metal-oxide systems. The key process parameters are systematically optimized as follows: 1) The acrylic acid (AAc) content in the photoresist is optimized to balance the carboxyl density and structural integrity, thus maximizing the ion-loading rate (Fig. 2); 2) the laser energy and exposure time are calibrated to ensure scaffold fidelity, which is critical for preserving structural details after sintering (Fig. 4); and 3) immersion conditions are optimized to enhance coordination efficiency without compromising scaffold stability. Post-sintering characterization (Figs. 6 and 7) confirms the formation of phase-pure metal oxides, whereas mechanical tests reveal the exceptional mechanical properties of the CoO microstructures (Fig. 8).

This study presents a highly efficient and versatile method for fabricating 3D metal-oxide nanoarchitectures by integrating femtosecond-laser TPP with a metal-ion coordination mechanism. Unlike conventional nanofabrication techniques, our approach introduces metal ions into the polymer scaffold after printing, thereby effectively avoiding the interference of metal ions during polymerization. By optimizing the photoresist formulation and coordination reaction conditions, we achieve high-efficiency metal-ion loading. Subsequent thermal sintering removes the organic framework, thus successfully yielding 3D metal-oxide structures with features measuring 410 nm at the minimum. By decoupling the printing and metallization stages, the proposed method enables the fabrication of intricate architectures with diverse metal-oxide compositions while maintaining submicron precision. In particular, its application spans from single-component oxides (e.g., NiO and CoO) to multi-metal-oxide systems (e.g., Ni_0.5Co_0.5O), thus highlighting its broad material adaptability. The proposed technique provides a novel paradigm for creating 3D functional devices to be used in cutting-edge applications such as energy-storage devices, photodetectors, and semiconductor sensors.

Microactuators, defined as nanoscale to microscale structures capable of actuation, hold significant potential for executing tasks such as serving as micro-sensors, micro-valves, and components in microrobotic systems. Among actuation mechanisms of microactuators, magnetic actuation stands out due to its wireless operation, precise directional control, and stable response. However, current magnetic microactuators predominantly rely on soft magnetic materials (e.g., Fe₃O₄ for doping or Ni for coating), which exhibit limited magnetic strength, hindering effective motion in low-Reynolds-number fluids. Hard magnetic materials, such as sintered NdFeB powders (5 μm average diameter, high saturation magnetization, and coercivity), offer a promising solution but face fabrication challenges in achieving microscale and three-dimensional (3D) geometries. Existing methods, for example, femtosecond laser direct writing struggles to integrate opaque and large NdFeB particles into high-resolution 3D structures, as particle sizes exceed the resolution limits of two-photon polymerization (TPP). To address this, the study here introduces a femtosecond laser mold-assisted fabrication strategy to overcome these limitations. By combining TPP-based template fabrication with NdFeB-polymer slurry infusion, we demonstrate the successful fabrication of hard magnetic microactuators with complex 3D architectures. The resulting microactuators exhibit superior motion performance compared to a soft magnetic counterpart, particularly in high-viscosity environments, which is attributed to the high remanence and coercivity of NdFeB. This approach bridges the gap between high-resolution 3D microfabrication and hard magnetic material integration, enabling the development of robust, functionally advanced microactuators or microrobots for biomedical and microfluidic applications.

The fabrication process begins with femtosecond laser patterning of templates in positive photoresist. A Ti∶sapphire femtosecond laser oscillator generates pulsed laser beams (central wavelength of 800 nm, pulse duration of 75 fs). The laser beams are expanded and directed through galvanometric scanners before being focused by a 60× oil-immersion. 3D scanning is achieved through coordinated control of x-y galvanometer mirrors and a z-axis precision stage. Approximately 0.1 mL positive photoresist is dispensed onto a coverslip and spin-coated at 1000 r/min for 10 s. The substrate undergoes pre-baking at 110 ℃ for 1 h to eliminate bubbles and enhance structural stability. Subsequent femtosecond laser exposure creates the microactuator architecture in the photoresist layer [Fig. 4(a)]. The developed template is obtained after 30 min immersion in developer [Fig. 4(b)], followed by thorough cleaning in deionized (DI) water. A slurry containing NdFeB particles and epoxy precursor is prepared at a mass ratio of 1∶1 and deposited onto the template surface [Fig. 4(c)]. Vacuum degassing at ambient temperature effectively removes entrapped air bubbles. Excess slurry is removed using lint-free wipes prior to thermal curing at 80 ℃ for 2 h [Fig. 4(d)]. Magnetization is performed using a pair of permanent magnets with a 1.2 T magnetic field intensity [Fig. 4(e)]. Final release of microactuators is achieved through template dissolution in ethanol, with subsequent transfer to DI water for further tests [Fig. 4(f)].

The developed femtosecond laser template-assisted method enables robust fabrication of hard magnetic microactuators with complex 3D geometries and tunable NdFeB content [Figs. 5(a)?(c)]. The microactuators retain structural integrity, with NdFeB particles uniformly encapsulated within epoxy resin [Fig. 5(b)]. Magnetic torque generation is directly related to the NdFeB mass fraction, enabling programmable actuation performance. Under rotating magnetic fields [Fig. 6(a)], the hard magnetic microactuators exhibit distinct rotational modes [Fig. 6(b)], with step-out frequencies increasing proportionally with applied field strength [Fig. 6(c)]. Crucially, the high remanence of NdFeB confers exceptional performance in viscous solutions. In 60 cP glycerol solution, the hard magnetic actuator outperforms a soft magnetic counterpart that cannot rotate under identical conditions [Fig. 6(f)]. The results highlight that hard magnetic materials overcome the torque limitations of soft magnetic materials. This advance expands the operational scope of magnetic microactuators for biomedical applications requiring robust fluidic manipulation.

This study presents a method for fabricating 3D microscale structures containing hard magnetic NdFeB materials through femtosecond laser two-photon polymerization combined with a template infusion approach, successfully realizing the preparation of hard magnetic microactuators. A systematic investigation of this technique demonstrates its capability to process diverse geometries, employ multiple materials, and achieve controlled magnetic content within specific ranges. The fabricated hard magnetic microactuators exhibit excellent structural integrity and leverage the advantages of hard magnetic materials, demonstrating superior locomotion capabilities in liquid environments. Comparative actuation experiments in high-viscosity fluids reveal that the fabricated hard magnetic microactuators have significant performance advantages over conventional soft magnetic microactuators. This methodology holds potential for broader applications in manufacturing hard magnetic microstructures, actuators, and microrobots, potentially enabling more demanding biomedical applications in the future.

Waveguide Bragg gratings are integrated photonic devices with high mechanical stability. In recent years, the femtosecond laser direct writing technology has provided a flexible method for the fabrication of complex optical components in the field of photonics, which has the advantages of being non-contact, non-destructive, maskless and “cold processing”, and has been utilized to fabricate integrated waveguide Bragg grating structures in a variety of bulk materials. This optical sensor can be applied to a wide range of environmental measurements such as temperature, strain, and refractive index. Sapphire crystals are widely used to fabricate sensors that can be used in extreme environments because of their excellent optical properties, ultra-high melting point, and electrical insulation. However, there are fewer reports on the fabrication of waveguide Bragg grating structures on sapphire crystals. In this paper, single mode transmission in sapphire crystals is achieved by optimizing the sizes of multilayer depressed cladding waveguides. And a uniform Bragg grating structure is integrated in the single mode waveguide. Our goal is to realize the fabrication of single mode waveguide Bragg grating integrated structures in sapphire bulk material with temperature sensing using the femtosecond laser direct writing technology.

Firstly, the repetition rate of the femtosecond laser is set to 200 kHz, and the effect of pulse energy on sapphire scribing is investigated for the same scanning speed (2 mm/s). After optimization of the processing parameters, a 10 mm long five-layer single mode depressed cladding waveguide is fabricated along the c-axis of the optical axis using a pulse energy of 25 nJ at a depth of 110 μm below the surface of the undoped sapphire bulk material. The cladding of the waveguide is composed of femtosecond laser-induced negative refractive index modified regions. And the waveguide properties are observed using an end-coupled system. Using the same repetition rate, the scanning speed and pulse energy are varied to integrate a second-order uniform Bragg grating structure with a period of 888.25 nm in the waveguide core. The morphology of the machined waveguide Bragg grating structure is observed using a microscope. After that, waveguide Bragg grating spectral characterization is carried out to observe the spectral variation of gratings with different lengths. Finally, the sensing performance of waveguide Bragg gratings at different temperatures is analyzed.

After fabricating a five-layer depressed cladding waveguide by the femtosecond laser direct writing technique, the waveguide is characterized using an end-face coupling system. The mode profile of the waveguide can be observed by a beam profiler and the mode profiles of waveguides with different sizes are compared (Fig. 5). By reconstructing the refractive index distribution from the obtained optical measurements, the normalized cutoff frequency of the waveguide is calculated to be lower than the normalized cutoff frequency required for single mode, thus confirming the single mode characteristics of the waveguide. The single mode waveguide transmission loss of 0.78 dB/cm is obtained based on the cutback method. Later on, a second-order uniform Bragg grating structure is integrated in a sapphire single-mode waveguide, and two Bragg resonance reflection peaks with different central wavelengths are found by characterizing its reflection spectrum. It is analyzed as a birefringent effect caused by transverse electric (TE) and transverse magnetic (TM) mode excitations, and this polarization phenomenon may be caused by the asymmetry of the refractive index change. The two reflection peaks correspond to 10.76 dB and 13.1 dB side lobe suppression ratios and 3 dB bandwidths of 0.327 nm and 0.347 nm, respectively (Fig. 9). Finally, the temperature sensing performance is tested by increasing the sensor temperature from room temperature to 650 °C. After three repeated experiments, the results show that the central wavelengths of the reflection peaks of the two modes shift toward longer wavelengths when the temperature increases, and the average temperature sensitivities of the TE and TM modes reach 24.8 pm/°C and 24.5 pm/°C, respectively (Fig. 11).

In this paper, single mode waveguide Bragg grating integrated structures are prepared in undoped sapphire bulk crystals using the femtosecond laser direct writing technique. This integrated structure consists of two structures, a single mode waveguide and a waveguide Bragg grating. The single mode properties of the waveguide are demonstrated based on mode profile observation and reconstruction of the refractive index distribution. And the transmission loss of the single mode waveguide is 0.78 dB/cm at 1550 nm. After that, a point-by-point method is used to integrate a second-order uniform Bragg grating in a single mode waveguide, which has a narrow 3 dB bandwidth (<0.35 nm) and a high reflectivity (>90% for TE mode and >95% for TM mode). When the temperature around the sensor is in the range of 100 ℃ to 600 ℃, the average temperature sensitivity reaches 24.8 pm/°C for TE mode and 24.5 pm/°C for TM mode. The fabricated sensor has promising applications in the field of temperature sensing.

The conventional separation of complex microstructure optical components leveraging mechanical contact has the problems of low separation efficiency and serious chipping at the edge of the separation. Ultra-fast laser separation technology offers the advantages of high separation quality, fast speed, and contact-free separation. However, laser distortion is caused when microstructure optical components are cut using the laser . Analyzing the causes of distortion and developing a method to eliminate it is the focus of high-quality and high-efficiency separation of complex microstructure optical components using ultrafast lasers. It has implications for the extended use of ultrafast lasers to separate optical lenses with non-planar structures.

This study utilizes the optical software Zemax to model the optical system and processing environment and analyzes the light path to reveal the underlying reason for the distortion of the laser light incident on the microstructure optical components. A liquid-phase matching medium-assisted ultrafast laser separation manufacturing technology is employed. A compound system is formed by placing the optical component into the liquid medium with a refractive index matching its own. This method reduces the distortion of the operating laser at the microstructures of the optical component, thereby enabling the establish of the modified surface inside the component. This study adopts a heat bath to supply a constant thermal environment for the entire optical component. The four sides of the modified surface of the optical component are heated uniformly at the same time, the thermal stress induces the cracks to extend from the surface to the inside under the guidance of the microcracks, and the cracks are finally realized.

This study derives that the sudden change in the refractive index of the propagating medium during the propagation of the laser from the environment medium to the target medium is the underlying reason for the distortion of laser through the functional microstructures of optical components. It causes the laser to generate irregular dispersion and reflection, deviating from the original path. Liquid phase-assisted ultrafast laser separation of complex microstructure optical components is proposed (Fig. 1).

This study utilizes Sellmeier formula [formula(3)] and Cauchy dispersion formula [formula(4)] to derive the refractive index values of optical component materials and liquid-phase refractive index-matching media under the operating laser wavelength, respectively, and selects the liquid-phase matching liquid whose refractive index matches the refractive index of the optical materials under the operating laser wavelength, such that the solid-liquid compound with the same refractive indices can be obtained.

This study obtains the Bessel beam length [formula(1)] and displacement [formula(2)] control relationship in different liquid-phase matching media. This study simulates the optical field distribution of the laser as well as the energy distribution of the Bessel beam primary axis after the laser enters the optics under three different values of the matching deviation (Fig. 5), and the results show that the smaller the matching deviation, the better the quality of the separation. In the presence of large matching deviations, the separation surface will show defects consistent with the microstructural cycle of the lens surface, primarily natural cracks that result in a morphologically uncontrollable separation. The laser cannot process the continuous modified surfaces required for separation inside the microstructure optics when the matching deviation is greater (Fig. 8).

An ultrafast laser separation process window for H-BaK3 is experimentally obtained, and the irregular separation of a variety of microstructure optical components proves the effectiveness and wide applicability of the technique.

The laser separation of complex microstructure optical components is analyzed by optical simulation to determine the reason for the failure of laser separation, i.e., the operating laser produces distortion at the microstructures to obstruct the formation of the Bessel beam inside the components, which leads to the inability to process the modified surfaces that cover the cross section of the components. Liquid-phase refractive index matching media are used to suppress laser distortion. Following analysis of the distribution of the optical field and the quality of the separated facets under different matching deviations, it is concluded that the use of a refractive index matching liquid that matches the optical material at the operating laser wavelength can effectively suppress distortion and obtain an excellent modified surface. Adopting the heat bath method to conduct contactless and uniform heating on the modified surface guarantees excellent separation quality. Compared with the traditional mechanical separation of optical components, liquid-phase-assisted laser separation of complex microstructure optical components can realize rapid and high-quality separation of optical components in bulk. The separation speed for H-BaK3 can reach 100 mm/s, the cut surface roughness is as low as 2.4 μm, and the chipping size is less than 20 μm, which provides a reference window for the laser modification separation process. This technology expands the application of ultrafast laser separation and can fulfill the demand for the customized separation of optical components.

In this study, we investigate the laser-induced structural coloring of titanium nitride (TiN)-coated stainless steel substrates via direct laser writing. As a noncontact surface modification method, laser processing offers several advantages over conventional coating techniques, such as ease of operation, high processing efficiency, and eco-friendliness. However, research on the laser processing of thin-film-deposited metal surfaces is at its early stage and the color formation speed remains slow. Unlike the conventional coating technologies physically depositing film layer, the magnetron sputtering coating technology can form uniform nanoscale films on stainless steel surfaces. We use a nanosecond fiber laser to prepare coloring patterns directly onto a TiN-coated stainless steel surface, creating a distinctive structural color. Through appropriate experiments, the effects of three factors—laser scanning speed, scanning interval, and output power—on pattern colors are assessed independently. The reflectance spectra and energy spectrum of structural color pattern and temperature analysis provide insights into the interaction between the laser and stainless steel surface. The mechanism of the laser-induced structural colors on the coated stainless steel surface is identified. The structural colors are not angle-dependent.

In this study, we investigated the effects of the scanning speed, scanning interval, and output power on the sample surface while setting the other parameters as constant. We calculated and analyzed the variation trends of the pattern color with respect to these three parameters. First, by varying the scanning speed in the range of 100?300 mm/s, we analyzed its effect on the surface structural color. Second, by setting scanning intervals of 0.02, 0.05, 0.07, and 0.10 mm, we observed the micro-nanostructures of the surface and discussed the effect of scanning intervals on the surface structural color. Next, for a fixed focal length of 156 mm, scanning speed of 100 mm/s, repetition frequency of 20 kHz, scanning interval of 0.05 mm, and an output power range of 46%?66% (in 4% steps), the trend of the structural color changes was investigated and the effects of output power were analyzed. Based on the experimental data, we determined the parameter range of the surface structural colors. By examining the microstructure and energy spectrum of the sample surface, we investigated the mechanism of color formation. Finally, we verified the color stability.

First, we discuss the effect of the laser parameters on the sample surface color. At low scanning speeds, a transition from yellow to violet, then to blue, and finally to silver-gray is observed, and the color transition is relatively fast. With the increase in the scanning speed, the color becomes lighter and the color transition speed decreases owing to a decrease in the total laser output energy (Fig. 2). Second, at a small scanning interval, the iridescence of surface structures at a scanning interval of 0.02 mm is observed under a microscope (Fig. 3). Under the conditions of the other parameters being set as constants while varying the output power, only the color brightness changes with increasing output power (Fig. 4), and yellow, violet, and blue are observed. This is the reason why the nonoverlapping laser beam part reduces during pattern marking. Based on the above experiments, different color-mixing techniques are used to obtain coloring patterns on the sample surface (Figs. 5 and 6). Table 1 summarizes the coloring parameters obtained from the above experiments, and the reflectance spectra of the five colors are measured (Fig. 7). Finally, based on the experiments, the microstructure of the sample surface is observed via scanning electron microscopy (Fig. 8). According to the elemental content analysis (Table 2), the presence of Cr, Mn, Fe, and Ni in the energy spectroscopy test results confirms the laser radiation penetrating the TiN coating and interacting with the stainless steel surface, which results in the melting of the stainless steel surface and its penetration into the TiN coating (Fig. 9). Thus, we conclude that the color formation on a metal surface is primarily dependent on the color of surface oxides or nitrides.

Basic structural colors are obtained via direct writing on the surface of a TiN-coated stainless steel plate using a nanosecond fiber laser. The results show that the metal surface color changes regularly with the scanning speed, scanning interval, and laser output power. The main colors are yellow, pink, blue, purple, and green. Various color-mixing techniques and unique color patterning can be used to produce consistent and clear images. However, the wide range of wavelengths for each color can result in color mixing, which affects the purity and saturation. In the green region, fine patterns obtained by laser scanning are observed, particularly for the laser-induced periodic surface structures (LIPSSs) with a period of approximately 1 μm. Subsequent energy spectrum analysis and temperature field calculations imply the presence of Ti, N, O, Cr, Mn, Fe, and Ni on the metal surfaces, which confirms that the laser radiation interacts with the stainless steel surface through the TiN coating, thereby melting and penetrating it. The color of the TiN coating or stainless steel surface under laser irradiation is primarily dependent on the color of surface oxides or nitrides and the LIPSS. The LIPSS occupies a smaller area than the oxide color area and is not the primary coloring mechanism that is affected by the angle or light source.

GH4169 is a nickel-based superalloy widely used in manufacturing aero-engine blades. However, it is affected by ultra-high cycle fatigue due to the low impact energy and small spot diameter of microscale laser shock peening. To address these issues, microscale laser shock peening reinforcement is employed to improve the fatigue properties of the material while controlling the depth of propagation of the shock wave to achieve macro-deformation of the thin blades and synergistic enhancement of fatigue properties. Currently, the effect of microscale laser shock peening on the ultra-high cycle fatigue properties of GH4169 alloy has not been investigated in domestic and international studies. Consequently, this effect is investigated in this study using an ultrasonic fatigue testing machine.

In this study, GH4169 alloy was used, and the fatigue specimen size was obtained through theoretical design and simulation. The specimens were subjected to microscale laser shock peening using three process parameters: 62 mJ impact energy and 1 time impact (62 mJ&1 time ), 62 mJ impact energy and 3 times impacts (62 mJ&3 times ), and 82 mJ impact energy and 1 time impact (82 mJ&1 time ). Subsequently, axially symmetric ultrasonic ultra-high cycle fatigue tests were conducted at room temperature, and fatigue fracture morphologies were analyzed using scanning electron microscope (SEM) with energy-dispersive spectroscope (EDS). Confocal laser scanning microscopy (CLSM) was used to analyze the surface morphologies strengthened by microscale laser shock peening, while X-ray diffraction (XRD) analysis was used to determine the distribution of residual stress in the surface layer. Additionally, the microstructure of the surface layer after microscale laser shock peening was observed using electron backscatter diffraction (EBSD). The analytical results were synthesized to reveal the strengthening mechanism of microscale laser shock peening in improving the ultra-high cycle fatigue properties of GH4169 alloy.

The S-N curves (Fig. 4) indicate that microscale laser shock peening improves the ultra-high cycle fatigue properties of GH4169. Fatigue fracture morphologies (Figs. 5?8) show that slip induces crack initiation, while microscale laser shock peening inhibits surface crack initiation. Additionally, surface morphology analysis (Figs. 9 and 10) reveals that microscale laser shock peening increases surface roughness. However, fatigue test results show that the increase in surface roughness due to laser shock peening is not the primary factor affecting the failure mechanism and properties of ultra-high cycle fatigue. Combining fatigue test results with surface residual stress distribution data (Fig. 11) indicate that the residual compressive stress introduced by microscale laser shock peening inhibits surface crack initiation. Furthermore, a combination of fatigue test results and microstructure EBSD analysis (Figs. 12 and 13, Table 4) reveals that refined surface grains and high-density dislocations contribute to fine grain and dislocation strengthening, respectively, thereby improving the ultra-high cycle fatigue properties of the material.

In this study, the influence of microscale laser shock peening on the ultra-high fatigue properties of GH4169 alloy is investigated using ultrasonic fatigue tests combined with the analysis of the fracture morphologies, surface morphologies, residual stress, and microstructure. This comprehensive approach enables a detailed understanding of the strengthening mechanism of microscale laser shock peening in improving the ultra-high cycle fatigue properties of GH4169 alloy. The main conclusions are as follows: microscale laser shock peening improves the ultra-high cycle fatigue properties of GH4169 alloy, allowing the material to withstand higher stress amplitudes at the same fatigue life. When the fatigue life reaches 1×10⁸ cycles, 62 mJ&1 time, 62 mJ&3 times, and 82 mJ&1 time specimens endure cyclic stress amplitudes of 450, 500, and 475 MPa, respectively, representing increases of 12.5%, 25%, and 18.75%. Microscale laser shock peening improves surface roughness, with R_a values of 0.926, 1.020, and 0.935 μm for 62 mJ&1 time, 62 mJ&3 times, and 82 mJ&1 time specimens, respectively. Increasing the impact energy and the numbers of impacts increases surface roughness but does not contribute to surface fatigue failure. Microscale laser shock peening introduces residual compressive stress within a depth range of 400?500 μm, with a gradient distribution. The 62 mJ&3 times specimen exhibits a maximum residual compressive stress of -761 MPa below the surface. Additionally, microscale laser shock peening refines grains of the surface layer and increases dislocation density, reducing the average diameters of the surface grains to 5.23, 4.09, and 4.68 μm, respectively. The average kernel average misoreintation (KAM) values of the surface layer reach 0.25°, 0.29°, and 0.27°, respectively, while the geometrically necessary dislocation (GND) densities increase to 0.78×10¹⁴, 0.9×10¹⁴, and 0.85×10¹⁴/m2, respectively. After microscale laser shock peening, the sources of fatigue cracks in GH4169 alloy shift from the surface to the interior because the residual compressive stress suppresses surface crack initiation and effectively balances the tensile stress. Additionally, surface layer grain refinement and the introduction of high-density dislocations contribute to fine grain and dislocation strengthening, respectively, further inhibiting surface crack initiation and improving the fatigue properties of the material.

To address the poor adhesion and high interface resistance of TiO₂ electrodes, a novel approach combining laser cladding and dealloying is proposed to fabricate porous Ti-based coatings with metallurgical bonding on pure titanium substrates. In this paper, we investigate the polarization characteristics of Cu-Ti alloys with different compositions, the effect of laser cladding on dealloying, and the photocatalytic performance of laser-clad, dealloyed, and porous Ti-based electrode materials.

Laser cladding technology was used to deposit Cu-Ti alloy powders with different compositions onto Ti substrates, forming a metallurgically bonded Cu-Ti alloy coating. Electrochemical dealloying was then performed to selectively remove the Cu component, resulting in a porous Ti-based electrode material. The coatings were characterized via scanning electron microscope (SEM), X-ray diffractometer (XRD), electrochemical workstation, and ultraviolet (UV)-visible spectrophotometry. The elemental composition, phase structure, polarization curves, UV-visible absorption spectra, transient photocurrent, and electrochemical impedance spectroscopy (EIS) results were analyzed.

Laser cladding produces Cu-Ti alloys with varying compositions, each generating different phases that directly influence the electrochemical dealloying process. Specifically, the Cu₂₈Ti₇₂ alloy mainly consists of CuTi₂, Ti, and CuTi phases; the Cu₃₃Ti₆₇ alloy contains CuTi₂, Ti, and CuTi phases; and the Cu₄₂Ti₅₈ alloy exhibits similar phases. As the Cu atomic fraction increases to 47%, the Cu₄₇Ti₅₃ alloy consists only of Ti and CuTi phases. A further increase in the Cu content leads to more complex phase compositions. For instance, the Cu₅₂Ti₄₈ alloy contains Ti, CuTi, Cu₄Ti₃, and Cu₂Ti phases, while Cu₅₈Ti₄₂ includes Ti, CuTi, Cu₄Ti₃, Cu₂Ti, and Cu₄Ti phases. The phase composition of the Cu₆₃Ti₃₇ alloy is simplified to Ti, Cu₂Ti, and Cu₄Ti phases, and the Cu₆₈Ti₃₂ alloy consists mainly of Cu₃Ti, Cu₄Ti₃, and Cu₄Ti phases. As the Cu atomic fraction reaches 73% or higher value, the alloy phases stabilize. The Cu₇₃Ti₂₇ alloy is composed of Cu, Cu₂Ti, and Cu₄Ti phases, whereas the Cu₇₇Ti₂₃, Cu₈₃Ti₁₇, and Cu₈₇Ti₁₃ alloys predominantly consist of Cu and Cu₄Ti phases. The polarization curves for the Cu-Ti alloys exhibit distinct regions, including active-dissolution, passivation, and over-passivation regions, with different dealloying characteristics. Dealloying occurs rapidly in the active-dissolution and over-passivation regions, whereas the process slows down significantly or even halts in the passivation region. When the Cu atomic fraction exceeds 87%, the alloys only exhibit active dissolution; when the Cu atomic fraction falls below 87%, the alloys exhibit passivation behavior. In the >28%?<73% Cu atomic fraction range, the active-dissolution region disappears, and dealloying only proceeds under over-passivation. When the Cu atomic fraction decreases below 28%, the over-passivation region disappears, and dealloying ceases, indicating a dealloying limit of 28% (Cu atomic fraction). The porous Ti-based electrode material exhibits significant absorption peaks at 200 nm and 300 nm, with increasing absorbance as the wavelength increases. This indicates a strong response to UV light and enhanced visible-light absorption. Transient photocurrent measurements reveal a rapid and clear response to light. Electrochemical impedance spectroscopy (EIS) data fitted using Z-view software reveal a relatively low impedance for the porous Ti-based electrode, with solution resistance and charge transfer resistance of 8.66 Ω and 12.13 Ω, respectively, providing favorable conditions for photocatalytic reactions.

In this study, the dealloying corrosion behavior of Cu-Ti alloy precursors with different compositions prepared via laser cladding is investigated, revealing the impact of laser cladding on the separation limits and critical potential of Cu-Ti alloys. The photocatalytic performance of the porous Ti-based electrodes prepared via a combined laser-cladding?dealloying process is also analyzed. Based on the results, the following conclusions can be drawn. The Cu-Ti alloys prepared via laser cladding exhibit significant phase compositional differences, which result in distinct characteristics in their polarization curves, which in turn can generally be divided into three regions: the active-dissolution, passivation, and over-passivation regions. The electrochemical dealloying process proceeds rapidly in both the active-dissolution and over-passivation regions; however, it is extremely slow or even halts in the passivation region. For a Cu atomic fraction of ≥87%, the Cu-Ti alloy exhibits only an active-dissolution region; when the Cu atomic fraction is <87%, the polarization curves exhibit passivation behavior. For a Cu atomic fraction of >28%?<73%, the active-dissolution region disappears, and sustained dealloying requires the over-passivation process. For a Cu atomic fraction of ≤28%, passivation prevents continued dealloying, indicating that the separation limit of Cu-Ti alloys prepared via laser cladding is 28% (Cu atomic fraction). Based on the characteristic potentials of the polarization curves for Cu-Ti alloys with different compositions, dealloying composition?potential diagrams for Cu-Ti alloys can be successfully constructed. These diagrams can guide the selection of dealloying processes for Cu-Ti systems and provide a reference for the dealloying corrosion behavior study of other alloy systems containing passivating elements. The porous Ti-based electrode materials prepared via the combined laser-cladding?dealloying process effectively limit light escape and exhibit a wide light-absorption range. Additionally, these electrodes respond rapidly to light, and the metallurgical bonding between the porous layer and substrate reduces the interface resistance. These factors significantly enhance their photocatalytic performance, providing new insights for the preparation of photocatalytic and other types of electrode materials.

Cr12MoV cold work die steel is widely used in the roll industry due to its advantages, such as minimal deformation, high abrasion resistance, and large bearing capacity. Cr12MoV rolls have high hardness and brittleness and are prone to spalling, pitting, cracking, and other micro-area damage during service. The service conditions of rolls are harsh, as they are often subjected to significant impact, extrusion, and external friction. This makes them highly susceptible to localized damage, such as large-area spalling and cracking defects, ultimately leading to Cr12MoV roll failure. Laser repair is a method used to restore damaged rolls, producing a dense structure and metallurgical bonding at the interface. It has become an important technology for repairing Cr12MoV rolls. However, under conventional laser processing, the surface hardness of the repair layer fluctuates significantly, and the hardness variation in the interface transition zone is pronounced. This negatively affects the repair of cold rolling work rolls used in rolling high-precision plates, severely impacting both rolling accuracy and roll service life. This paper introduces the application of laser beam oscillation in welding to promote the transformation of columnar crystals into equiaxial crystals. This technique addresses issues such as poor surface hardness uniformity in repaired Cr12MoV rolls, large hardness fluctuations in the interface transition zone, and metallurgical defects. The findings provide a valuable reference for achieving high-quality surface repair of Cr12MoV rolls.

Laser cladding of Fe90 alloy powder is applied to the surface of Cr12MoV cold rolling work rolls to create additive layers. A subsequent layer is then deposited on both the conventional manufactured layer and the single manufactured layer. Initially, the forming quality of the additive layers under the two different surface conditions of the manufactured layers is compared. Next, the microstructural features of the cross-sections and longitudinal sections of the additive layers are examined using a metallographic microscope, followed by a comparative analysis of the microstructures at the same depths within the additive layers. X-ray diffraction is then used to analyze the phase compositions of the additive layers. Additionally, the energy dispersive spectrometer (EDS) line scanning is conducted to examine the distribution of the primary elements, Fe and Cr, in the repair layers. Finally, the hardness and wear resistance of the additive layers under different manufactured layer surface conditions are tested. By combining microstructural analysis, element distribution, and other evaluations, the impact mechanism of manufactured layer surface grinding on the uniformity of surface hardness in the additive layers is explored.

Before manufactured layer surface grinding, the additive layer surface is uneven, with an average surface roughness of 24.41 μm (Fig. 3). The microstructure of the cross-sectional and longitudinal sections of the additive layer consists of a mixture of columnar and equiaxed crystals. The surface hardness varies widely, reaching as low as 178.1 HV, with a hardness fluctuation coefficient of 45.7. Additionally, extensive delamination pits are observed under friction wear (Figs. 12 and 14). After manufactured layer surface grinding, the additive layer surface becomes flatter, with an average surface roughness reduced to 15.50 μm (Fig. 3). Furthermore, the manufactured layer enhances the consistency of laser energy absorption on its surface, improving the uniformity of the melt pool temperature and flow field. This stabilization of grain growth conditions transforms the microstructure into a uniform dendritic crystal form (Figs. 11 and 12). The range of surface hardness decreases to 97.5 HV, with the hardness fluctuation coefficient reduced to 23.9, significantly improving surface hardness uniformity. The friction coefficient decreases from 0.65 to 0.47, and both the width and depth of the wear tracks are significantly reduced (Figs. 12 and 13).

In this study, an Fe90 alloy additive layer is prepared on the surface of Cr12MoV cold rolling work rolls using laser cladding technology, and the effect of manufactured layer surface grinding on the additive layer is studied. Surface grinding of the manufacturing layer improves the forming quality of the additive layer, reduces surface roughness, decreases the temperature gradient between the center and edges of the melt pool, and increases the solidification rate. It also results in a more uniform distribution of Fe and Cr elements within the additive layer, enhances the consistency of the crystal growth environment, and promotes a transition to a uniform microstructure, thereby suppressing microstructural variations in the additive layer. Overall, this study demonstrates that incorporating a surface grinding process into the laser cladding procedure enhances the forming quality, hardness uniformity, and wear resistance of the additive layer.

316L stainless steel is widely used in the aerospace, biomedical, and nuclear power industries owing to its high corrosion resistance, ease of processing, adequate strength, and reasonable cost. Laser directed energy deposition (LDED) is an additive manufacturing technology that has the advantages of high flexibility, high processing efficiency, and high degrees-of-freedom, making it suitable for manufacturing fine and complex components. However, fields such as aerospace, biomedical, and nuclear power require components that achieve low surface roughness. The main objective of this study is to utilize the jet electrochemical polishing method to post-polish an LDEDed 316L stainless steel surface. Before polishing, this study elucidates the anisotropy of the microstructure of the LDEDed front, top, and side faces, and analyzes the differences in their anodic dissolution behavior in a sodium chloride-glycol electrolyte, which provides a sufficient theoretical basis for the subsequent jet electrochemical polishing. The surface morphologies and microstructures of the different faces were analyzed in detail after polishing. The aim of this study is to provide an empirical and technical support for the polishing of LDEDed components to improve their surface quality.

The microstructural anisotropy of three faces—front, top, and side—of LDEDed 316L stainless steel was analyzed using optical microscopy (OM), X-ray diffraction (XRD), and electron backscatter diffraction (EBSD) techniques. Then, the open-circuit potentials, polarization curves, and alternating current (AC) impedance curves of the three faces were measured using an electrochemical workstation to assess the anodic dissolution behavior of the individual surfaces in a sodium chloride-glycol electrolyte. A jet electrochemical polishing device was then utilized to polish the three faces with the face-scan mode to reduce their roughness. The three-dimensional morphology and surface roughness of the three faces were measured using a laser confocal microscope to evaluate the effect of polishing. Finally, scanning electron microscopy (SEM) was used to analyze the micromorphology and chemical composition.

The LDEDed 316L stainless steel shows significant anisotropy in the microstructure (Fig. 4), physical phase (Fig. 5), and texture (Fig. 6). The results of the open-circuit potential (Fig. 7), potentiodynamic polarization curves (Fig. 8), and electrochemical impedance spectra (Fig. 9) indicate that the corrosion-dissolution behavior of the three different faces of the LDEDed 316L stainless steel exhibits anisotropy. Specifically, the corrosion resistance of the three faces decreases in the following order: top face < front face < side face. Scratches on the front, top, and side faces disappear after the jet electrochemical polishing surface sweep, all of which show good mirror effect (Fig. 10). The roughness decreases from the original 1.057 μm to 0.177 μm, 0.200 μm, and 0.171 μm, respectively (Fig. 11). Bright and dark zones exist on all three faces after polishing (Fig. 10), which are caused by the different microstructures of the bright and dark zones (Fig. 12). SEM and energy-dispersive X-ray spectroscopy (EDS) results (Fig. 14) show that after jet electrochemical polishing, a large number of dendrites exist in the dark zone; dendrites in the bright zone are basically dissolved.

In this study, LDEDed 316L stainless steel was subjected to jet electrochemical polishing to reduce its roughness. First, the anisotropy of the microstructures of the front, top, and side faces was analyzed. Subsequently, the electrochemical anodic dissolution behaviors of the three faces in a sodium chloride-glycol electrolyte were analyzed. Finally, the three faces were subjected to jet electrochemical polishing with surface sweep mode to reduce their roughness to less than 0.2 μm. The reasons for the appearance of bright and dark areas on the polished surfaces were elucidated by analyzing the microstructure and chemical composition of the bright and dark zones.

Yttria-stabilized zirconia (YSZ), as a material for thermal barrier coatings, is widely used for the protection of high-temperature components in gas turbines and aero-engines. YSZ coatings prepared by atmospheric plasma spraying have high surface roughness, loose structures, and many pores and cracks. These factors can lead to the frictional wear of the coating, the susceptibility to molten salt corrosion causing the coating to fail and peel off, and the reduction in lifespan of the coating. Laser polishing can reduce the surface roughness, improve the aerodynamic performance of the components, hinder the penetration of molten salts, and reduce the stress concentration. Laser glazing can improve the loose structure, repair the surface defects, enhance the resistance to thermal corrosion, and thereby extend the service life. This paper proposes an integrated laser polishing and glazing processing technology. The coating is treated using the femtosecond laser pulse train mode for integrated polishing and glazing. Preliminary conclusions on the mechanism of femtosecond laser polishing and glazing are drawn, providing valuable reference for femtosecond laser pulse train mode processing.

This paper uses atmospheric plasma spraying technology to prepare YSZ thermal barrier coatings on the substrate. Femtosecond laser technology is utilized for the polishing and glazing integration treatment of the coating. A laser confocal microscopic system is employed to observe and measure the surface roughness and the three-dimensional morphology before and after polishing. Scanning electron microscope and metallographic microscope are used to characterize the cross-sectional morphology of the coating. An X-ray diffractometer is used for phase analysis of the original coating and the coating after experimentation. A mixture of NaCl and Na?SO? molten salt powder is first uniformly spread on the coating surface before and after polishing at a deposition amount of 20 mg/cm2, and then heated to 900 ℃ at a rate of 10 ℃/min and held at temperature for 4 h to test the resistance of the coating to molten salt corrosion.

The surface of the coating forms a flat polished area after undergoing the polishing and glazing integration treatment with a femtosecond laser. In the pulse train mode, the coating surface exhibits vaporization and a “melting peak filling valley” effect (Fig. 3). The change in the number of sub-pulses has little impact on the surface roughness of the coating.The coating roughness can be reduced by more than 74% after femtosecond laser scanning in pulse train mode (Fig. 4). As the number of sub-pulses increases, the width of cracks on the coating surface gradually increases. After the polishing and glazing integration treatment, a dense glazed layer forms on the surface of the coating. In the pulse train mode, as the number of sub-pulses increases, the thickness of the glazed layer gradually decreases, but this leads to an increase in crack width. When the number of sub-pulses is 10, the porosity of the glazed layer is reduced to the minimum of 6.8% (Fig. 6). The main phase component of the coating before and after the polishing and glazing integration treatment is the non-equilibrium tetragonal phase of zirconia (t′-ZrO₂), and the polishing and glazing integration treatment does not lead to the formation of harmful monoclinic phase zirconia (m-ZrO₂) (Fig. 8). After thermal corrosion, the original coating develops distinct horizontal cracks between the thermal barrier coating and the bond coat. In contrast, the coating that has undergone the integrated polishing and glazing treatment maintains its surface integrity after thermal corrosion, the internal structure of the coating is largely preserved, and no horizontal cracks are observed (Fig. 9).

The femtosecond laser polishing and glazing integration process can significantly reduce the surface roughness of YSZ coatings, reducing the original roughness S_a of 7.5 μm to below 2.0 μm. Using the pulse train mode can effectively improve the surface morphology and the quality of the glaze layer, with the porosity of the glaze layer being as low as 6.8%. Compared to the original coating, the coating after the polishing and glazing integration treatment exhibits superior resistance to thermal corrosion. Under the process with 10 sub-pulses, the coating adheres well, and no significant transverse cracks are observed. The femtosecond laser pulse train mode optimizes the proportion of material vaporization and remelting under a laser action through a more gentle heating method and a longer duration of action. By adjusting the number of sub-pulses, precise control over the morphology and thickness of the glaze layer can be achieved.