Nickel-based superalloys are typically used as turbine disks and turbine blade materials in aeroengines owing to their excellent high-temperature performance. Because of the different service conditions of the turbine disk and turbine blade, the premature failure of joints can be avoided using gradient materials. However, a change in the composition of the gradient transition zone can change the microstructure, which significantly affects the properties of the alloy. Therefore, the evolution of the microstructure and the hardness of laser melted nickel based superalloys must be investigated, providing a basis for the laser additive manufacturing of gradient nickel-based superalloys.
The materials used in this experiment are IC10 directional superalloy and FGH9X powder superalloy with a particle size of 0.15-0.25 mm. IC10 and FGH9X powders are prepared as nickel-based superalloy powders with different compositions, and alloy ingots with 11 typical gradient components are prepared by changing the ratio of the two alloys. (From samples F100 to F0, the mass fraction of FGH9X decreases by 10% intervals, and the mass fraction of IC10 increases by 10% intervals. For example, the FGH9X mass fraction of the F80 sample is 80% and the IC10 mass fraction is 20%). A 50 g nickel-based superalloy mixed powder is prepared and placed in a copper crucible. The laser power and opening time are set to 5 kW and 2 s, respectively. A melting experiment is performed in an argon atmosphere, and the alloy ingot samples are obtained after air cooling. The microstructures of the samples are observed using an optical microscope (Leica DM4M) and scanning electron microscope (Apreo S LoVac). The primary dendrite spacing, γ' phase size, and γ' phase content are measured using the Image Pro Plus software. Thermo-Calc 2020b is used to simulate the nucleation driving force of the γ' phase in the alloys. A hardness test is performed using a micro Vickers hardness tester (MH-6) with a load of 4.9 N and a holding time of 15 s.
The microstructure of the 11 nickel based alloy samples prepared via laser melting is composed of dendrites, and primary and secondary dendrites. When the sample changes from F100 to F0, the dendrite morphology of the alloys remains almost unchanged, and the primary dendrite spacing is in the range of 100-120 μm. Based on literature review, the dendrite morphology and dendrite spacing of the alloys are primarily affected by the cooling conditions, and the effect of alloy composition is insignificant. An analysis of the alloy microstructure shows that the 11 types of nickel-based alloys are composed of the γ and γ′ phase, carbides, and the γ/γ′ eutectic phase (Fig. 3). When the alloy composition changes from F100 to F0, the content and size of the γ′ phase increase continuously (Fig. 5), which is due to the gradual increase in Al and Ta contents in the alloys. The nucleation driving force of the γ' phase increases and more γ' phases precipitate. In addition, owing to element segregation, the content and size of the interdendritic γ' phase differ significantly from those of the dendritic trunk γ' phase. The content of γ' phase in interdendritic zone is more than that in dendritic trunk and the size of γ' phase in interdendritic zone is larger than that in dendritic trunk. In addition, the results show that the change in alloy composition does not significantly affect the microhardness, and that the overall hardness value is approximately 500 HV. This is because as the alloy composition changes from F100 to F0, the content of solid-solution strengthening elements and the content of carbides with high hardness and brittleness in the alloys decrease gradually, whereas the hardness and size of the γ' phase increase gradually; therefore, the hardness of the 11 types of nickel-based alloys is similar.
Primary dendrite and secondary dendrite arms are developed in the microstructures of the 11 types of nickel-based superalloy ingots prepared via laser rapid melting. The change in alloy composition does not significantly affect the dendrite morphology, and the average primary dendrite spacing is 110 μm. The microstructures of the alloys are composed of the γ phase, γ′ phase, carbides, and the γ/γ′ eutectic phase. As the powder nickel-based alloy content in alloy ingots decreases, the content and size of the γ 'phases increase continuously. As a result of element segregation, Al, Ti, Ta and Nb, which are the constituent elements of the γ 'phase, segregate in the interdendritic regions, thus causing the content and size of the γ' phase in the interdendritic regions are greater than those in the dendritic trunk. The microhardness values of the 11 samples are similar, and the overall hardness value is approximately 500 HV.
Aiming at the dense porosity defects detected using X-ray in laser melting deposition connected AlSi10Mg alloy, which is manufactured using selective laser melting, the characteristics of the defect and its influence on the mechanical properties are analyzed, and the elimination method of the defect is also explored. Selective laser melting methods cannot prepare large-scale aerospace structural components, laser melting deposition overcomes the forming size limitations, and provides a feasible solution for the additive manufacturing of large structural components.
In this study, the AlSi10Mg alloy prepared using selective laser melting was used as the base, and a laser melting deposition connection experiment was performed. First, connection samples under different laser powers were prepared. The relationship between the distribution range of dense porosity defects and the laser power was analyzed by an X-ray inspection, and the effect of dense porosity on the microhardness of the connection region was measured. The dense porosity was observed and analyzed using scanning electron microscopy (SEM) to determine its type and formation mechanism. Then, a substrate preheating experiment was performed to explore the best preheating temperature, which is used to solve the dense porosity defects in the connection region. Finally, the microhardness and tensile properties of the connection region before and after preheating were tested, and the fracture morphology was analyzed.
The inspection of the prepared connection samples with different energy densities shows that the dense porosity primarily appeared at the position of the interface fusion line between the connection region and substrate, and the dense porosity decreases the microhardness of the bonding zone (Fig. 7). Using SEM to observe the dense porosity, it is discovered that the porosity is primarily hydrogen porosity (Fig. 9). The solubility of hydrogen in molten pool decreases fastly, causing most bubbles to escape and form bubbles. However, the bubbles rise slowly, giving numerous bubbles no time to escape. So, some dense porosity defects are formed in this position. The preheating experiment shows that the optimum preheating temperature is higher than 100 ℃, which can effectively solve the dense porosity defects. The hardness at the fusion line reaches 90.8 HV after preheating (Fig. 14); the tensile strength is 287 MPa, and the elongation is 5.0% (Fig. 15). The observation of the fracture morphology shows that the fracture types before and after preheating are all brittle fractures and numerous dimples and quasi-cleavage morphologies can be observed on the fracture surface (Fig. 17).
Dense porosity is the primary defect of the AlSi10Mg alloy, which is produced using laser melting deposition. The defect is located at the position of the interface fusion line between the connection region and substrate, which is characterized as a watermark phenomenon on the X-ray inspection film. The accumulation of dense porosity results in a much lower hardness at the fusion line than in the connection zone and substrate. Preheating can effectively change the agglomeration effect of the dense porosity, making them evenly dispersed from the fusion line to the connection region. The mechanical properties of the samples after preheating were significantly improved. The hardness at the fusion line increased by 45% compared with that without preheating. The tensile strength increased by 19%, and the elongation increased to 5.0%.
Understanding the crater morphology on the surface of a paint layer after a single laser pulse can effectively suppress the superposition effects of multiple laser parameters and the photothermal and photomechanical effects of a pulse overlap. This helps reveal the laser-material interaction mechanism and provides a basis for the optimization of laser parameters. In recent years, many scholars have simulated the morphology of craters on the surface of a paint layer with the help of finite element software after nanosecond pulsed laser action based on the ablation mechanism. The laser parameters are then optimized based on the simulation results. For nanosecond pulsed lasers, the main mechanism of the laser-material interaction varies at different energy densities (the main mechanism is the ablation mechanism at low energy density, and the plasma shock and thermal radiation mechanism at high energy density). The ablation mechanism, plasma shock, and thermal radiation mechanism have different effects on the morphology of the crater. This study aims to establish a model of the damage form and the removal process of the paint layer during a single pulse of a nanosecond laser under different energy densities, to reveal the differences in the influence of the laser-material mechanism on the morphology of craters under different energy densities, and to provide a reference for the precise control and parameter optimization of the paint removal effect at high and low energy densities.
A nanosecond pulsed laser with a wavelength of 1064 nm and beam energy following a Gaussian distribution was applied to the epoxy primer surface. The diameter, depth, and three-dimensional morphology data of the craters on the surface of the paint layer were measured using a 3D optical surface profiler after the laser pulse. A simulation model of crater morphology was established based on the ablation mechanism and the fitting relationship between the depth (d) of the craters and energy densities (F). MATLAB was used to simulate the morphology of the craters in the energy density range of 13.58-27.16 J/cm2, and an experimental verification was carried out. Error analysis of the experimental and simulation results under a high density revealed the influence of the plasma shock and thermal radiation mechanism on the morphology of the crater. The model correction and experimental verification were carried out based on the plasma shock and thermal radiation mechanisms.
The simulation model of crater morphology based on the ablation mechanism has an error of less than 5% for crater depth and diameter at a low energy density (13.58-16.98 J/cm2), less than 5% for crater depth error at a high energy density (20.37-27.16 J/cm2), and up to 40% for diameter error (Fig.5). The error analysis shows that at a high energy density, the plasma shock and thermal radiation mechanisms are the main reason for the diameter error (Fig.7). After the model was corrected based on the above analysis, the diameter and depth errors of the craters under a high energy density were controlled within 5%, which significantly improved the accuracy of the model (Fig.9). The model shows that at a low energy density, the surface of the crater is approximately rotated paraboloid, and the profile of the crater is similar to a parabola; at a high energy density, the surface of the crater can be regarded as a combination of multiple normally-distributed surfaces, and the crater profile as a combination of multiple normal distribution curves.
At different energy densities, differences in the laser-material mechanism are noted; the ablation mechanism at a low energy density and the laser plasma shock and thermal radiation mechanisms at a high energy density are the main interaction mechanisms. Differences in the laser-material interaction mechanisms cause damage to the paint layer. Compared to the ablation mechanism, the plasma shock and thermal radiation mechanisms lead to an increase in the amount of paint removed near the surface of the crater and a wider profile near the crater surface. A simulation model of crater morphology is established for different laser-material mechanisms, thereby effectively improving the model accuracy. The study results provide a reference for the accurate control of the laser paint removal process and the optimization of paint removal parameters under high and low energy densities.
The stern shaft is an important device for the power transmission of ships. However, corrosion is the main failure mode for stern shafts, which are subjected to the attack of Cl- and microorganisms in seawater for a long time in a marine environment with high salt and humidity. The vibration and shock of the ship stern shaft aggravate the wear of the stern shaft. The traditional anticorrosion strengthening of the ship stern shaft surface involves coating its surface with anticorrosion coatings, such as ethylene resin, epoxy resin, and chlorinated rubber. Although these anticorrosion coatings protect ship stern shafts to a certain extent, most are toxic and harm the natural environment, which seriously violates the current trend of green development. Therefore, developing a green, clean, and pollution-free surface modification method for ship stern shafts has not only economic value, but also broad environmental value. An attempt is being made to develop a eutectic high-entropy alloy based on laser cladding technology to provide an effective method for green anticorrosion and wear-resistant modification of ship stern shaft surfaces.
The base material was 42CrMo steel. In an argon atmosphere, FeCoNiCrNb0.5Mo0.25 high-entropy alloy cladding layers were prepared using a laser with a spot diameter of 4 mm and a scanning speed of 3 mm/s at five different laser powers (1200, 1300, 1400, 1500, and 1600 W). The phase compositions of the cladding layers were analyzed using X-ray diffraction. The microstructures of the cladding layers were observed by scanning electron microscopy. The hardness values of the cladding layers were measured using a Vickers hardness tester. Friction and wear experiments were conducted using a multifunctional friction and wear tester. The corrosion resistances of the cladding layers were tested using an electrochemical workstation.
With the increase in laser power, the molten pool depth of the high-entropy alloy cladding layers increases (Fig. 3). FeCoNiCrNb0.5Mo0.25 entropy alloy cladding layers prepared using different laser powers are composed of an incomplete eutectic structure of FCC and Laves phases. With an increase in laser power, the content of lamellar nano-eutectic structure first increases and then decreases. The eutectic microstructure can be promoted by increasing the laser power appropriately, but too high laser power results in a stronger dilution effect of Fe in the substrate on the high-entropy alloy cladding layer, which weakens the promoting effect of Mo and Nb on the eutectic microstructure. The microstructure of the high-entropy alloy cladding layer prepared using a laser power of 1400 W is better than that of high-entropy alloy cladding layers prepared using other laser powers, and the microstructure is nano-eutectic with a lamellar spacing of approximately 86 nm. The increase in the laser power reduces the average hardness of the cladding layer, and the high-entropy alloy cladding layer sample prepared at laser power of 1200 W has the highest microhardness of 665.8 HV1.0, which is approximately 2.34 times that of the substrate (Fig. 5). With an increase in the laser power, the wear resistance of the FeCoNiCrNb0.5Mo0.25 high-entropy alloy cladding layer first increases and then decreases (Fig. 11). The high-entropy alloy cladding layer (1400 W cladding layer) sample owns excellent eutectic structure and the best wear resistance, with the lowest wear rate of 0.079 mm3·N-1·m-1. Compared with the substrate, the FeCoNiCrNb0.5Mo0.25 high-entropy alloy cladding layers have better corrosion resistance. However, there is no obvious linear relationship between the laser power and corrosion resistance of the high-entropy alloy cladding layer (Table 7). The lowest self-corrosion current density of the FeCoNiCrNb0.5Mo0.25 high-entropy alloy cladding layer is 1.716×10-6 A·cm-2. The existence of a eutectic structure reduces the corrosion resistance of the cladding layer to some extent. The corrosion resistance of the 1400 W cladding layer with a better eutectic structure is poor, and the self-corrosion current density is 4.332×10-6 A·cm-2.
Laser power affects the microstructure by changing the content of the cladding layer elements and solidification conditions. Properly increasing the laser power can promote the eutectic microstructure, but too high laser power strengthens the dilution effect of Fe in the matrix on the high-entropy alloy cladding layer and weakens the promotion effect of Mo and Nb on the eutectic microstructure. With an increase in the laser power, the microhardness of the cladding layers decreases owing to the increase in substrate dilution. The wear mechanisms of high-entropy-alloy cladding layers include oxidation wear, adhesion wear, and abrasive wear. The oxide film on the worn surface plays a significant role in protecting the lower metal. The eutectic structure with alternating soft and hard distributions reduces the material loss, and the cladding layer prepared at 1400 W has the lowest wear rate. In 3.5% NaCl solution, corrosion occurs around the oxide on the cladding layer surface, and the existence of a eutectic structure intensifies galvanic corrosion and reduces corrosion resistance.
Thermal barrier coatings can protect hot end components from high temperature, high pressure, and high stresses, which are extensively applied in gas turbine engine blades, combustion chamber, and ducting and nozzle guide vanes. However, many impurities (sodium, sulfur) exist in the operating gases and fuels, condensation of which leads to the formation of molten corrosive salts during the long-time service process, which may lead to the serious hot corrosion failure and reduce the lifetime of thermal barrier coatings. Plasma-sprayed thermal barrier coatings possess the typical characteristic of pores and laminar structure, which may provide penetration paths for molten corrosive salts into the coating. To improve the hot corrosion resistance, it is necessary to produce a denser layer to prevent molten salts from penetrating into porous coatings. Laser alloying technology has the advantages of high energy density, short action time, and reliable processing quality, and thus can change the loose porous structure of the plasma-sprayed thermal barrier coatings. Meanwhile, the oxidation reaction of self-healing materials at high temperature can produce some oxidation products, which can further fill the pores and cracks in the porous coatings. Therefore, the combination of laser alloying technology and self-healing materials is an alternative method to improve the hot corrosion resistance of thermal barrier coatings, which is few reported. The purpose of this paper is to study the effect of laser alloying on the microstructure, phase composition, and hot corrosion properties of thermal barrier coatings.
In this study, the double-layer thermal barrier coatings of NiCrAlY/8YSZ were deposited onto superalloy substrate via air plasma spraying. The mixture powders of TiAl3 particles with 10% mass fraction and Ceria and Yttria-stabilized Zirconia (CYSZ) ceramic were pre-placed on the plasma-sprayed thermal barrier coatings, and then processed using a fiber-coupled semiconductor laser. The mixture of 25% NaCl and 75% Na2SO4 as the corrosive salts was spread on the surface of the plasma-sprayed and laser-alloyed thermal barrier coatings with deposition content of 10 mg/cm2. The hot corrosion test was conducted in a furnace at 900 ℃ for 4 h. Finally, the microstructure, phase composition and hot corrosion behaviors of the plasma-sprayed and laser-alloyed thermal barrier coatings were systematically investigated.
After the laser alloying treatment, the porous and laminar microstructures in the plasma-sprayed thermal barrier coatings were eliminated. As a result, dense columnar crystal structure and some segmented microcracks were formed in the laser-alloyed thermal barrier coatings (Fig. 2). The detrimental monoclinic zirconia (m-ZrO2) disappears after the air plasma spraying and the laser alloying treatment, and all phases in the 8YSZ powder turn to non-equilibrium tetragonal zirconia (t'-ZrO2) and cubic zirconia (c-ZrO2) (Fig. 3). Because of the rapid cooling and solidification rate of air plasma spraying and laser alloying treatment, the phase transformation of t'-ZrO2 to m-ZrO2 is restrained. After the hot corrosion at 900 ℃ for 4 h in molten salts (25% NaCl+75% Na2SO4), the corrosion products of Y2(SO4)3 and m-ZrO2 were found in the plasma-sprayed thermal barrier coatings, and Y2(SO4)3 and Al2O3 were detected in the laser-alloyed thermal barrier coatings [Figs. 5(a) and (b)]. Compared with the plasma-sprayed thermal barrier coatings, there are less corrosion products on the surface of the laser-alloyed thermal barrier coatings [Fig. 4(b) and Fig. 6(b)]. Therefore, the hot corrosion resistance of the laser-alloyed thermal barrier coatings is superior to that of the plasma-sprayed thermal barrier coatings. On the one hand, the dense columnar structure in the laser-alloyed thermal barrier coatings can inhibit the penetration of molten salts; on the other hand, the self-healing agent TiAl3 undergoes oxidation reaction at high temperature, and the formed Al2O3 and a small amount of TiO2 can fill the cracks, which can further reduce the hot corrosion reaction between molten salts and yttria stabilizer [Figs. 9(c), (d), and (e)].
In this study, a typical double-layer 8YSZ/NiCrAlY thermal barrier coating was prepared via air plasma spraying technology, and then the self-healing agent TiAl3 was introduced into the thermal barrier coatings through laser alloying technology. Microstructure, phase composition, and hot corrosion properties of the plasma-sprayed and the laser-alloyed thermal barrier coatings were investigated. The surface of the plasma-sprayed thermal barrier coatings is relatively rough, and there are many microcracks and pores within it. While the surface of the laser-alloyed thermal barrier coatings is smooth, and some fine segmented microcracks and dense columnar crystals are formed. After the hot corrosion in the 25% NaCl+75% Na2SO4 molten salt at 900 ℃ for 4 h, it was found that there were corrosion products Y2(SO4)3 and harmful m-ZrO2 in the plasma-sprayed and the laser-alloyed thermal barrier coatings; however, less corrosion products existed in the latter. The hot corrosion resistance of the laser-alloyed thermal barrier coatings is much better than that of the plasma-sprayed thermal barrier coatings. On the one hand, the oxidation reaction of self-healing agent TiAl3 during the process of high-temperature hot corrosion produces Al2O3 and less TiO2, and they can fill some cracks in the coating and reduce the penetration paths of the corrosion salts. On the other hand, the laser-alloyed layer with dense columnar structure can inhibit the penetration of molten salt and reduce the occurrence of thermal corrosion reaction. Finally, the hot corrosion resistance of laser-alloyed thermal barrier coatings is greatly improved.
Selective laser melting (SLM) is a commonly used technology for the additive manufacturing (AM) of metal material. It uses a high-energy laser beam to melt the metal powder layer-by-layer and finally prints the desired parts. During the SLM process, the printed part on the top plane (the XOY plane in Fig. 2 in the vertical printing direction) and the printed part on the side plane (the YOZ plane in Fig. 2 in the parallel printing direction) have different heating histories and temperature gradients. Therefore, the two planes have significantly different microstructures. This anisotropy in the microstructure is bound to introduce anisotropy to the performance. Recently, several studies have been conducted on the effect of microstructural anisotropy on mechanical properties. The unified conclusion is that printed samples have better mechanical properties in the vertical printing direction than in the parallel printing direction. However, few studies have been conducted on the effect of microstructural anisotropy on the corrosion behavior of printed parts, and their conclusions are different. Therefore, it is necessary to further investigate this issue. The aim of this study is to investigate the corrosion behaviors in different directions (the XOY and YOZ planes) in 316L stainless steel (SS) prepared using SLM through electrochemical measurements and propose internal causes of these corrosion behaviors, which have not yet been described.
The 316L SS parts are first prepared using SLM. To obtain samples in different directions, including the XOY and YOZ planes, samples are cut according to the diagram shown in Fig. 3. In this study, the forged 316L SS is used as the counterpart after solution treatment. The body and surface density of 316L SS are measured using the Archimedes drainage and metallographic methods, respectively. The microstructures of the SLMed sample on XOY and YOZ planes are characterized by electron backscattered diffraction (EBSD) and a scanning electron microscope (SEM). The phase structures of all samples are measured by X-ray diffractometry (XRD). The corrosion behaviors are explored by measuring the open-circuit potential (OCP), potentiodynamic polarization, and electrochemical impedance spectroscopy (EIS). In addition, the internal causes of this effect can be explained by the potentiostatic polarization and characterization of the surface topographies of all parts after corrosion.
The results show that the body density of 316L SS prepared using SLM is 99.38%, which is close to that of its forged counterpart (99.7%). The surface densities of the SLMed sample on XOY and YOZ planes are 99.7% and 99.87%, respectively, indicating that the surface densities in the different directions are almost similar. The XRD results confirm that the additive manufacturing technology does not change the phase structure of the 316L SS (Fig. 5). However, a clear discrepancy is evident in the grain orientation for both planes from the EBSD tests (Fig. 7). On the XOY plane, more (101)-oriented grains are observed, whereas on the YOZ plane, more (111)-oriented grains are observed. According to the literature , (111)-oriented grains are more resistant to corrosion. The grain sizes in both planes differ slightly according to the EBSD test results (Fig. 8). The average grain size of the SLMed sample on the YOZ plane (9.51 μm) is slightly larger than that of the SLMed sample on the XOY plane (7.35 μm). However, the grain sizes of SLMed sample on XOY and YOZ planes are significantly smaller than that of the forged counterpart (50-100 μm). The results from the electrochemical tests show that the corrosion resistance of the SLMed sample on the XOY plane is better than that of the SLMed sample on the YOZ plane, and the SLMed sample on both planes are superior to the forged counterpart, as confirmed by the OCP measurements (Fig. 9), potentiodynamic polarization curves (Fig. 10), and EIS measurements (Fig. 11). The improved corrosion resistance of the SLMed sample on the XOY plane is attributed to the fewer (111)-oriented grains on the XOY plane, and consequently, to the more compact passive film formed on the XOY plane based on the results of potentiostatic polarization measurements (Fig. 12). These conclusions are further confirmed by observing the SEM morphologies of the three corroded samples. The sizes of the inclusions on the XOY and YOZ planes of the printed samples are much smaller than those of the forged part (Fig. 13). In addition, the inclusion on the XOY plane remains closely combined with the matrix after corrosion, demonstrating outstanding corrosion resistance. However, for both the SLMed sample on the YOZ plane and its forged counterpart, the case worsens. A clear corrosion gap is present around the inclusions after corrosion, particularly for the forged counterpart, indicating poorer corrosion resistance.
First, compact 316L SS samples are produced using SLM. Their densities are 99.38%, which are considerably close to that of the forged parts (99.7%). There is a notable difference in the microstructure between the XOY and YOZ planes in the printed part. On the XOY plane, more (101)-oriented grains are observed. However, on the YOZ plane, more (111)-oriented grains are observed. This microstructural anisotropy has a significant effect on the corrosion behavior of 316L SS printed using SLM. The corrosion resistance of the SLMed sample on the XOY plane is better than that on the YOZ plane, and the SLMed samples on both planes are superior to their forged counterparts. The better corrosion resistance of the SLMed sample on the XOY plane results from fewer (111)-oriented grains, leading to more compact passive films formed on the surface. The SEM morphologies of inclusions in the three corroded samples show that the size of the inclusions on the XOY plane is smaller than that on the YOZ plane and that of forged counterpart. In addition, the corrosion gap between the inclusion and matrix on the XOY plane is far less than that on the YOZ plane and that of forged counterpart, indicating better corrosion resistance.
Owing to the excellent strength, plasticity, and corrosion resistance, 316L stainless steel is widely used in nuclear and chemical industries. The efficient cutting of thick plates is realized using lasers, which are high-energy-density heat sources. During the laser cutting process, the plate material melts and is blown off vertically under the action of a coaxial compressed gas. Therefore, a kerf is formed. During a rapid thermal cycle, an extremely thin recast layer (the order of microns) is formed on the surface of the kerf. During the solidification of the recast layer, a particular temperature gradient and fluid motion significantly influence the morphology and the texture of the structure at room temperature. In previous studies, researchers have mainly focused on the influence of parameters, such as laser power, cutting speed, and pressure of compressed gas, on the cut formation and its quality. Few studies have focused on the microstructural morphology and formation mechanism of the recast layer. The differences between the as-solidified microstructure and the substrate may lead to non-negligible changes in the properties of the edge, which in turn affects the overall characteristics. To study the morphology and microstructural growth of the recast layer, an 18 mm thick 316L austenitic stainless steel plate is taken as the object of laser cutting for this study. The solidification mechanism of the recast layer at different kerf sites during the laser cutting process is revealed.
An 18 mm thick 316L austenitic stainless steel plate was employed as the base metal for this study. A pulsed laser was used to cut the base metal to form a kerf. N2 was chosen as the compressed gas, and its flow direction was coaxial with the laser. Representative specimens were then sampled to analyze their surfaces. Transverse and surficial microstructural morphologies of the recast layer, under the laser action, were analyzed using scanning electron microscopy and electron back scattering diffraction (EBSD). In addition, the recast surfaces were cleaned using anhydrous ethanol. The transverse surfaces were treated using coarse grinding, fine grinding, and polishing techniques. The polished surface was then etched with diluted aqua regia (volume ratio of HCl, HNO3 and H2O is 3∶1∶4).
The distribution of the main elements on the surface of the recast layer is analyzed using energy dispersive spectroscopy. The results indicate that no significant element change occurred along the thickness, except for a slight loss of Fe (Table 5). The grain growth mode of the recast layer is further analyzed using EBSD at the 1/3 site from kerf top and the kerf bottom site . The results indicate that epitaxial growth is the primary growth mode. However, the proportion of non-epitaxial growth at the 1/3 site from kerf top (Fig. 5) is observed to be higher than that at the bottom site (Fig. 7). A comparison between the IPF orientation distribution and pole figures in Figs. 9 and 10 also shows that the grain growth at the 1/3 site from kerf top exhibits some fluctuations with unmixed and unperturbed features.
The results show that a small amount of Fe evaporates from the recast layer surface. A variation in flow state from turbulent at the top to laminar at the bottom surface is observed, with an increase in thickness and needle-like grains. For crystal orientation, the ratio of the epitaxial growth at the top surface of the recast layer is lower than that at the bottom surface. Such a random distribution of epitaxial growth is caused by the turbulent flow at the former, whereas the dominant epitaxial growth is induced by the laminar flow at the latter. Considering the grain profiles, the γ phase in the base metal is equiaxed, whereas the δ phase is arranged in a banded form. The morphology of the γ phase grains in the recast layer is irregular and coarsen by approximately 2 times compared to those of the base metal. However, the δ phase is dispersed and refined from 1/6 to 1/2 of the base metal. Under the conditions of an extremely high-temperature gradient and a disordered disturbance owing to melting, a substantially reduced duration of δ phase formation with considerable dispersion is produced.
Results and Discussions The grains of the deposited layer obtained via the conventional electrodeposition are relatively coarse. Large gaps appear in the intergrain bonding sites. The surface of the deposited layer is loose and porous, and sediment agglomeration is severe. After laser remelting and electrochemical deposition are performed, the pores on the surface of the deposited layer are reduced significantly, the grain-to-grain binding is firm, and the compactness is improved. Although the number of laser remelting/electrochemical deposition interaction processes is increased, the electrode response rate and polarization do not decrease (Figs. 3, 4, 5, and 6). Laser remelting causes mutual diffusion between titanium and copper, thus resulting in a composite remelting layer containing titanium-copper intermetallic compounds such as CuTi, CuTi2, and Cu4Ti. The generation of these intermetallic compounds increases the deposition surface active sites, enhances the polarization of electrodeposition, and accelerates the electrodeposition reaction rate. Under the same electrodeposition time (30 min), the thickness of the deposited layer obtained via the conventional electrodeposition is 79.67 μm, and the thickness of the composite coating obtained via laser remelting and electrochemical deposition is 145.36 μm (Figs. 7, 8, and 9). The copper deposited layer achieved via laser remelting and electrochemical deposition shows a higher binding force with the matrix than that achieved by conventional electrodeposition. Under a 50 N load force, the scratch morphology of the deposited layer obtained via laser remelting and electrochemical deposition remains relatively complete. The composite coating indicates good adhesion to the matrix (Figs. 10, 11, and 12). The composite remelted layer achieved via laser remelting exhibits better resistance to high-temperature oxidation than the copper layer. In addition, the copper layer achieved via conventional electrodeposition is more oxidized than the copper layer achieved via laser remelting and electrochemical deposition (Fig. 13). Owing to the effect of laser remelting, interdiffusion occurs between titanium and copper, thus resulting in the formation of titanium-copper intermetallic compounds and a slight decrease in the conductivity of the composite coating. However, as the deposition time increases, the titanium content in the composite coating achieved via the subsequent laser remelting and electrochemical deposition decreases gradually, thus resulting in an increase in the electrical conductivity of the composite coating (Fig. 14).
To improve the poor deposition quality and binding force of copper electrodeposited directly on titanium, a new laser remelting/electrochemical deposition interaction process is proposed to prepare a titanium-copper alloy layer and thicken the copper layer. The micromorphology, cross-sectional elements, coating thickness, and phase of the composite remelting layer obtained from laser remelting/electrochemical deposition interaction process are investigated. The effect mechanism of the laser electrochemical interaction on the electrical conductivity, high-temperature oxidation resistance, and adhesion to the substrate of the composite coating is discussed.
First, laser melting pretreatment is performed to replace the conventional chemical pretreatment, which simplifies the process, reduces environmental pollution, and reduces hazard to the human body. Thickening of the copper layer and its metallurgical binding to the matrix are achieved via laser remelting and electrochemical deposition. The first laser remelting is performed to obtain a titanium-copper seed layer, whereas the second laser remelting is performed to modify the surface of the deposited layer and reduce defects, such as porosity, crevices, and grain agglomeration. Consequently, the copper layer particles become refined and denser, which can facilitate the subsequent electrodeposition as well as thicken and improve the performance of the copper layer. The morphologies of the composite coating and composite remelting layer are characterized via scanning electron microscopy (Sigma HV-01-043, Carl Zeiss), and the elemental distribution on the cross-section of the composite remelting layer is analyzed using an X-ray energy-dispersive spectroscope connected to a scanning electron microscope. The cross-sectional morphology and thickness of the composite coating and composite remelting layer are observed using an optical microscope (Axio Imager A2M, ZEISS). The composite remelting layer obtained via interactive treatment is analyzed using an Xpert Pro X-ray diffractometer (PANAlytical Company, Netherlands). The adhesion between the composite coating and composite remelting layer achieved via the interactive treatment is evaluated using an automatic adhesion scratch tester (WS-2005).
Compared with conventional electrodeposition, the combination of laser remelting and electrochemical deposition can increase the electrode response rate on the surface of the remelted layer by approximately 44%. The pores on the surface of the deposited layer are reduced significantly, and the grains are bonded firmly. The results show that the combination of laser remelting and electrochemical deposition can significantly improve the deposition quality of the copper deposited layer on the titanium alloy surface, the bonding force with the substrate, and the high-temperature oxidation resistance.
Laser powder bed fusion (LPBF) and laser melting deposition (LMD) are the main processes of metal additive manufacturing. The powder materials used for LPBF or LMD are normally spherical pre-alloyed powders. The high difficulty and cost have limited the number of powder brands available commercially. In-situ alloying, which utilizes different pure elements or alloy powders as raw materials to produce synthesis block alloys, is an efficient and low-cost research method that utilizes the interaction between high-energy laser and powder during the additive manufacturing process. Recently, in-situ alloying of titanium alloys, aluminum alloys, nickel-based superalloys, steel, and high entropy alloys has been explored. 300 series stainless steel is one of the earliest commercial LPBF materials with excellent printing properties. Nevertheless, there are few reports on preparing this material system by elemental powder in-situ alloying. This study examined 304L stainless steel and fabricated multiple groups of block samples with different processes by LPBF in-situ alloying using the elemental mixed powder to explore the effects of the LPBF process parameters on the microstructure and properties of the in-situ alloying samples.
Spherical Fe, Cr, Ni elemental blended powder, and 304L pre-alloyed powder prepared by gas atomization were used as the raw materials. The elemental powder was >99.5% purity and was mixed using a rotary mixer. The samples were prepared using a DLM-280 metal 3D printer produced by Hangzhou DediBot. The laser power and scanning speed were used as variables. For the elemental blended powder, laser powers (W) of 80, 110, 140, 170, 200, 230, 260, and 290, and the scanning speeds (mm/s) of 500, 650, 800, 950, 1100, and 1250, and 48 groups of parameters were selected. For the 304L pre-alloyed powder, the laser power and scanning speed were 230 W and 650 mm/s, respectively. The other parameters were constant: spot diameter of 80 μm, layer thickness of 30 μm, hatching space of 80 μm, and rotation angle of 90° between layers. The sample density was calculated using the Archimedes drainage method. An Olympus BX53 optical microscope and FEI Quanta650 field emission scanning electron microscope were used to characterize the microstructure of the sample. Energy dispersive spectrometry and electron backscatter diffraction were conducted to investigate the elemental content and microstructure. An EM500-2A hardness tester was used to test the microhardness of the samples, and each sample was tested at 16 points with a load of 4.9 N and a loading time of 15 s.
The process window was plotted according to the measured density of the in-situ alloyed sample (Fig. 3). The density of the sample showed a converging trend. The maximum density of the sample was 7.855 g/cm3 at an energy density of 147 J/mm3 (P=230 W, v=650 mm/s), and the corresponding relative density was 99.05%, which was comparable to the sample prepared from the 304L pre-alloyed powder. When the energy density was 71 J/mm3, many lack-of-fusion holes and unmelted particles were observed in the in-situ alloyed sample (Fig. 4). With increasing energy density, the lack-of-fusion holes and unmelted particles gradually disappeared. A typical cellular structure and columnar grains could be observed in the sample when the energy density was 242 J/mm3 (Fig. 5). In terms of composition uniformity, with increasing energy density, the macroscopic and microscopic homogenization driving force of the in-situ alloyed samples during the LPBF process was strengthened, so the composition uniformity was improved. The composition was uniform when the energy density was 242 J/mm3 (Fig. 6). Owing to Fe, Cr, and Ni enrichment zones, the in-situ alloying samples consisted of a face-centered cubic (FCC) and body-centered cubic (BCC) duplex when the energy density was 147 J/mm3. By contrast, the sample consisted of an FCC single phase when the energy density was increased to 242 J/mm3, which is the same as the sample prepared from the 304L pre-alloyed powder (Fig. 8). With increasing energy density, the strengthening effect of fine-grains, duplex structure, dislocations, and residual stress caused by composition inhomogeneity was weakened, and the microhardness of the in-situ alloyed samples decreased gradually from 302 HV to 224 HV (Fig. 10).
The in-situ alloyed 304L stainless steel sample with a uniform microstructure was prepared by LPBF technology using Fe, Cr, and Ni mixed powder. The effects of the process parameters on the density, microstructure morphology, composition uniformity, phase structure, and microhardness of the sample were studied and compared with the LPBF sample fabricated using pre-alloyed 304L powder. The results showed that the density of the in-situ alloyed sample converged gradually in the process window and reached the highest value of 7.855 g/cm3 when the energy density was 147 J/mm3 (P=230 W, v=650 mm/s). The corresponding relative density was 99.05%, indicating that the elemental mixed powder has excellent LPBF printing properties. With increasing energy density, the uniformity of composition was improved; the phase structure changed from FCC + BCC to FCC, and the microhardness gradually decreased from 302 to 224 HV. When the energy density was 242 J/mm3, the average value of each element in the sample reached the proportion of powder mixing, indicating a uniform composition of the structure. The process parameters of LPBF significantly affect the microstructure of the sample. The multiphase structure and fine grain structure generated by hyperinflation in the in-situ alloying samples at a low energy density can improve the hardness of the sample by up to 26.4% compared to the pre-alloyed LPBF sample.
Results and Discussions The prepared sapphire through hole diameter is 200 μm, the taper is 2?, the maximum hole edge damage width is 5.74 μm, and there is almost no heavy condensate at the hole edge (Fig. 9). In general, a suitable laser energy density enables the sample to melt and vaporize, which can reduce the splash of the molten material and improve the quality of the through-hole surface (Fig. 4). The higher the repetition frequency of the laser, the longer the corresponding pulse width, and the thermal effect of the laser is enhanced. Owing to the influence of the thermal action, spatter is produced around the processing zone, affecting the roughness of a through-hole surface (Fig. 6). The laser scanning speed directly affects the pulse overlap rate, and an appropriate pulse overlap rate can effectively remove materials and play an important role in reducing the taper of the through hole (Fig. 7).
Sapphire crystals are commonly used in high temperature pressure sensor chips because of their excellent physical and chemical properties. In particular, in the manufacturing process of an all-sapphire fiber Fabry-Perot cavity high-temperature pressure sensor, the sapphire base must be drilled, and the quality of the through-hole is an important factor to ensure that the fiber can be vertically inserted to realize the effective transmission of optical signals. It is difficult to ensure the quality of sapphire through-holes using traditional mechanical and chemical processing methods. Laser machining technology has non-contact, non-mask, and simple process characteristics, which make it prominent in the field of hard brittle material micro-machining. Recently, there have been reports on the use of ultrashort pulse lasers to punch holes in sapphire; however, its high cost and low processing efficiency make this machining method mainly used in high-precision surface microstructure etching. Short pulse lasers, such as 355 nm ultraviolet nanosecond lasers, are widely used in industrial production owing to their low cost and high efficiency; however, there are few reports on their application in sapphire through hole processing. In this study, we analyze the impact mechanisms of different laser energy densities, repetition frequencies, and scanning speeds on sapphire processing using the control variable method, as well as the parameter space of sapphire through hole processing. We aim to fabricate through holes with diameters of approximately 200 μm and tapering angles of less than 5? on 500 μm thick sapphires without notches, heavy coagulation, or other damages. We hope that our method can provide a new, low-cost, and efficient way to process and cut the internal structure of sapphire.
This study uses sapphire wafers with a diameter of 5.08 cm and thickness of 500 μm. Before and after the experiment, the sapphire is ultrasonically cleaned with anhydrous ethanol to remove impurities on the surface and avoid effects on the laser processing results. First, using the control variable method, different laser energy densities, repetition frequencies, and scanning speeds are adjusted to drill holes into the sapphire. Subsequently, scanning electron microscopy and laser scanning confocal microscopy are used to analyze the surface morphology and 3D structure of the through-hole under different conditions. The effects of different laser parameters on the surface morphology and taper of the through-hole are studied. Finally, the optimal parameter space to obtain high-quality sapphire through holes is selected based on the experimental results.
In this study, a 355 nm ultraviolet nanosecond laser is used to successfully prepare a through hole on a 500 μm-thick sapphire wafer with a quality identical to or better than that of the ultra-short pulse laser (Table 4). Through the analysis of the drilling results assisted by laser scanning confocal microscopy and scanning electron microscopy, we determine a combination of laser energy density of 31.12 J/cm2, laser frequency of 30 kHz, and laser scanning speed of 0.5 mm/s. The results show that a micro through hole with good surface morphology and small taper can be prepared on the sapphire wafer by adjusting the appropriate combination of laser parameters. The results indicate a certain reference value for processing and cutting sapphire internal structures with a 355 nm all-solid-state ultraviolet nanosecond laser.